Superior Pinning Properties in
Nano-Engineered YBa2Cu3O7−δ
Giorgio Ercolano
Department of Materials Science & Metallurgy
Corpus Christi College
This dissertation is submitted for the degree of
Doctor of Philosophy
2011
Si qui forte mearum ineptiarum
lectores eritis manusque uestras
non horrebitis admouere nobis...
Declaration
I declare that, except where otherwise stated, this dissertation is the
result of my own work and includes nothing which is the outcome of
work done in collaboration.
No part of this dissertation had been submitted at Cambridge or any
other University for a degree, diploma or other qualification.
The length including tables, footnotes, bibliography and appendices
do not exceedes the limit of 60000 words.
Giorgio Ercolano
Cambridge
April 2011
Acknowledgements
This dissertation describes the research work on controlled nano-
engineering of high temperature superconducting materials carried
out at Cambridge University Materials Science and Metallurgy De-
partment in the Device Materials Group between November 2007 and
February 2011.
I would like to express my gratitude to my supervisor, Prof. Judith
MacManus-Driscoll, for the help and valuable advice provided during
this project.
I would also like to acknowledge...
Dr Stuart Wimbush for the helpfull discussions on nanoinclusions flux
pinning and on the vortex path model
Dr John Durrell and Dr Marcus Weigand for the trainings and help
on microfabrication and critical current measures
Dr Sophie Harrington for the helpfull discussions on pulsed laser de-
position of doped YBa2Cu3O7−δ thin films and target preparation
Dr Marco Bianchetti for the practical help with the photolithography
and critical currents characterisation of the thin films and for the
usefull discussions
Prof. Haiyan Wang, Joon Lee Hwan and Dr Lata Sahonta for the
amazing TEM work
Dr Xavier Moia and Dr Thomas Fix for the valuable discussions on
structural data interpretation
Dr Nadia Stelmashenko for the help with the atomic force microscope
Mary Vickers for the help with the x-ray diffraction analysis
and Dr Lara San-Emeterio-Alvarez, Dr David Muoz-Rojas, Diana Iza,
Oon Jew Lee and Dr Sohini Kar for the support
The work described in this dissertetion was supported by NESPA,
Nano-Engineered Superconductors for Power Application, a frame-
work of the Marie Curie Research Training Network, funded within
the EUs 6◦ framework program.
Superior Pinning Properties in
Nano-Engineered YBa2Cu3O7−δ
Giorgio Ercolano
Large electrical current transport in the absence of energy losses is the
key factor in commercial applications of high temperature supercon-
ductors. This thesis demonstrates an easy and inexpensive bottom-up
technique to produce self assembled nanorods, segmented nanorods
as well as nanoparticles in YBa2Cu3O7−δ thin films grown by pulsed
laser deposition. The structural and morphological characteristic of
the pinning landscapes produced are investigated and correlated to
their effects on the superconducting properties of the thin films.
In particular two pinning landscapes are investigated: Ba2YNbO6
nanorods are grown in YBa2Cu3O7−δ thin films using a Ba2YNbO6
doped YBa2Cu3O7−δ pulsed laser deposition targets and
Ba2(Y/Gd)(Nb/Ta)O6 segmented nanorods together with (Y/Gd)2O3
nanoparticles are grown in (Y/Gd)Ba2Cu3O7−δ thin films using a
Ba2YNbO6 + Gd3TaO7 doped YBa2Cu3O7−δ pulsed laser deposition
targets.
The Ba2YNbO6 + YBa2Cu3O7−δ is deeply characterised and the ef-
fects of the deposition parameters are analysed. Ba2YNbO6 is demon-
strated to be an interesting novel pinning addition capable to increase
the critical current and to reduce the YBa2Cu3O7−δ critical currents
angular dependencies anisotropy.
The Ba2YNbO6 + Gd3TaO7 + YBa2Cu3O7−δ is found to produce a
new complex pinning landscape extremely effective. At high fields
the synergetic combination of the different defects typology is shown
to generate an interesting new feature in the critical current angular
dependencies.
Chapter 1 is an introduction to superconductivity, the fundamentals of
the field are briefly presented. In chapter 2 the discussion in focused on
pinning in high temperature superconductors. Cuprates and in par-
ticular YBa2Cu3O7−δ are presented. The pinning phenomenon and
the practical pinning engineering in thin films is also discussed in this
chapter. Chapter 3 describes the thin films preparation methods and
the characterisation techniques used in the research work. Chapter 4
and 5 are focused on the Ba2YNbO6 doped YBa2Cu3O7−δ thin films.
Chapter 4 is an introduction to Ba2YNbO6 doped YBa2Cu3O7−δ, the
preliminary results obtained on Ba2YNbO6 doped YBa2Cu3O7−δ thin
films are shown in this chapter. The crystalline structure, the mor-
phology and the superconducting properties of thin films deposited
adopting different deposition parameters are analysed and discussed
in chapter 5. In chapter 6 the new complex pinning landscape of
Ba2(Y/Gd)(Nb/Ta)O6 and (Y/Gd)2O3 in (Y/Gd)Ba2Cu3O7−δ is pre-
sented. Concluding remarks on the research described in the work
ends the thesis in a brief final chapter 7.
Contents
Contents vii
List of Figures xi
1 Introduction 1
1.1 The discovery of superconductivity . . . . . . . . . . . . . . . . . 1
1.2 The Meissner effect and the London penetration depth . . . . . . 2
1.3 Critical magnetic field, critical current density. Two types of su-
perconductors . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4
1.4 Finally a microscopic theory of superconductivity . . . . . . . . . 7
1.5 The high temperature superconductors revolution . . . . . . . . . 7
2 Pinning in YBa2Cu3O7−δ thin films 9
2.1 Cuprates . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9
2.1.1 Weaklinks and the texturing of cuprates . . . . . . . . . . 10
2.2 YBa2Cu3O7−δ . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 11
2.2.1 Crystal structure . . . . . . . . . . . . . . . . . . . . . . . 11
2.2.2 Oxygen deficiency . . . . . . . . . . . . . . . . . . . . . . . 12
2.3 Pinning . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 14
2.3.1 The dissipative phenomenon . . . . . . . . . . . . . . . . . 14
2.3.2 Dissipation-free current flow, the pinning force . . . . . . . 15
2.3.3 Defects as pinning centers . . . . . . . . . . . . . . . . . . 16
2.4 Practical Pinning Engineering . . . . . . . . . . . . . . . . . . . . 17
2.4.1 Non superconducting secondary phase addition . . . . . . 17
2.4.2 Alternative routes . . . . . . . . . . . . . . . . . . . . . . . 19
vii
CONTENTS
3 Experimental Techniques 21
3.1 Sample Preparation . . . . . . . . . . . . . . . . . . . . . . . . . . 21
3.1.1 Pulsed Laser Deposition . . . . . . . . . . . . . . . . . . . 22
3.1.2 Target Preparation . . . . . . . . . . . . . . . . . . . . . . 25
3.1.3 Sample Patterning . . . . . . . . . . . . . . . . . . . . . . 27
3.2 Structural Morphological characterisation . . . . . . . . . . . . . . 29
3.2.1 Structural Analysis . . . . . . . . . . . . . . . . . . . . . . 29
3.2.2 Microscopy . . . . . . . . . . . . . . . . . . . . . . . . . . 35
3.3 Superconductivity Properties characterisation . . . . . . . . . . . 36
4 Ba2YNbO6 doped YBa2Cu3O7−δ: preliminary results 39
4.1 The Nb introduction in the YBa2Cu3O7−δ . . . . . . . . . . . . . 39
4.1.1 Nb doped ceramic samples . . . . . . . . . . . . . . . . . . 40
4.1.2 Ba2YNbO6 use in YBa2Cu3O7−δ thin film . . . . . . . . . 41
4.2 Ba2YNbO6 perovskite additions to YBa2Cu3O7−δ: the preliminary
results . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 43
4.2.1 The Ba2YNbO6 synthesis, the pulsed laser deposition tar-
get preparation and thin films deposition . . . . . . . . . . 43
4.2.2 Crystalline structure analysis: x-ray diffraction data . . . . 45
4.2.2.1 Crystalline phases identification . . . . . . . . . . 46
4.2.2.2 Crystalline phases orientation . . . . . . . . . . . 47
4.2.2.3 Crystallographic matching and strain . . . . . . . 48
4.2.3 Nanostructure analysis: Atomic Force Microscopy and Trans-
mission Electron Microscopy . . . . . . . . . . . . . . . . . 51
4.2.3.1 Surface topography (Atomic Force Microscopy) . 51
4.2.3.2 Cross-section transmission electron microscopy . 53
4.2.4 The superconducting properties: Tc, Jc(B) and Jc(B, θ) . . 58
4.2.4.1 Transition temperature, Tc . . . . . . . . . . . . . 59
4.2.4.2 The critical current density . . . . . . . . . . . . 60
4.2.5 Concluding remarks to the preliminary results . . . . . . . 63
5 Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters 65
viii
CONTENTS
5.1 Ba2YNbO6 perovskite additions to YBa2Cu3O7−δ: the effects of
the deposition rate on the nanostructure and the superconducting
properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 65
5.1.1 Ba2YNbO6 doped YBa2Cu3O7−δ target preparation and
thin film deposition . . . . . . . . . . . . . . . . . . . . . . 66
5.1.2 Crystalline structure analysis: x-ray diffraction data . . . . 67
5.1.2.1 Crystalline phases identification, phases orientation 67
5.1.2.2 Epitaxy quality: rocking curve . . . . . . . . . . 68
5.1.2.3 Epitaxy quality: reciprocal space maps . . . . . . 69
5.1.3 Nanostructure analysis: Atomic Force Microscopy and Trans-
mission Electron Microscopy . . . . . . . . . . . . . . . . . 76
5.1.3.1 Surface topography (Atomic Force Microscopy) . 76
5.1.3.2 Cross-section transmission electron microscopy . 79
5.1.4 The superconducting properties: Tc, Jc(B) and Jc(B, θ) . . 82
5.1.4.1 Transition temperature, Tc . . . . . . . . . . . . . 82
5.1.4.2 The critical current density . . . . . . . . . . . . 84
5.1.5 Concluding remarks to the analysis of the effects of the de-
position rate on the nanostructure and the superconducting
properties . . . . . . . . . . . . . . . . . . . . . . . . . . . 90
5.2 Ba2YNbO6 perovskite additions to YBa2Cu3O7−δ: deposition pa-
rameter optimization . . . . . . . . . . . . . . . . . . . . . . . . . 91
5.2.1 Ba2YNbO6 doped YBa2Cu3O7−δ pulsed laser deposition
target preparation and thin films deposition . . . . . . . . 91
5.2.2 Surface Topology: Grain Connectivity . . . . . . . . . . . 92
5.2.3 The superconducting properties: Tc, Jc(B) and Jc(B, θ) . . 98
5.2.3.1 Transition temperature, Tc . . . . . . . . . . . . . 98
5.2.3.2 The critical current density . . . . . . . . . . . . 100
5.2.4 Concluding remarks to the deposition parameters optimiza-
tion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 103
6 Ba2YNbO6 and Gd3TaO7 simultaneous doping of YBa2Cu3O7−δ 104
6.1 Ba2YNbO6 and Gd3TaO7 doped YBa2Cu3O7−δ target preparation
and thin films deposition . . . . . . . . . . . . . . . . . . . . . . . 105
ix
CONTENTS
6.2 Crystalline structure analysis: x-ray diffraction data . . . . . . . . 106
6.2.0.1 Crystalline phases identification . . . . . . . . . . 106
6.2.1 Crystalline phases orientation . . . . . . . . . . . . . . . . 108
6.3 Cross-section transmission electron microscopy . . . . . . . . . . . 109
6.4 The superconducting properties: Tc, Jc(B) and Jc(B, θ) . . . . . . 112
6.4.1 The critical current density . . . . . . . . . . . . . . . . . 112
6.4.1.1 The critical current density: Jc(B) . . . . . . . . 113
6.4.1.2 The critical current density angular dependence:
Jc(B, θ) . . . . . . . . . . . . . . . . . . . . . . . 115
6.4.1.3 A new pinning feature . . . . . . . . . . . . . . . 117
6.4.2 Concluding remarks to Ba2YNbO6 and Gd3TaO7 simulta-
neous doping of YBa2Cu3O7−δ . . . . . . . . . . . . . . . . 117
7 Conclusions and further work 120
Appendix 123
References 146
x
List of Figures
1.1 Differences in magnetisation behaviours of a perfect conductor and
a superconductor, the Meissner effect. a) Perfect conductor cooled
in zero field; b) Perfect conductor cooled in applied magnetic field;
c) Superconductor cooled in zero field; d) Superconductor cooled
in applied magnetic field. . . . . . . . . . . . . . . . . . . . . . . . 3
1.2 Structure of an isolated Abrikosov vortex [1] . . . . . . . . . . . . 6
1.3 Magnetisation curves for a type I and a type II superconductor [1] 6
2.1 a) Crystal Structure of YBa2Cu3O7−δ; b) CuO2 planes; c) CuO
chains. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 11
2.2 a) Superconductive transition temperature Tc for ten samples of
Ba2YCuOx with varying x values; b) Refined crystallographic cell
parameters for Ba2YCuOx with varying x values; from [2]. . . . . 13
3.1 Schematic of a typical pulsed laser deposition system [3] . . . . . 23
3.2 The target sintering thermal process . . . . . . . . . . . . . . . . 26
3.3 Sketches of a sample at the different stages of a photolithographic
process. a) As deposited sample; b) Sample after the first spin
coating of the photoresist; c) Sample after the developing of the
electrode mask; d) Sample after the silver layer and gold layer
deposition. e) Sample after the lift-off process; f) Sample after
the second spin caoting of the photoresist; g) Sample after the
developing of the trak mask; h) Sample after the milling/cleaning
process. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 30
3.4 Micrography of the surface of a patterned sample . . . . . . . . . 31
xi
LIST OF FIGURES
3.5 Schematic of a diffractometer operated in the Bragg-Brentano Ge-
ometry . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 31
3.6 Schematic of a diffractometer geometry operating a φ scan . . . . 32
3.7 Schematic of a diffractometer geometry operating a rocking curve
scan . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 33
3.8 Schematic of a diffractometer geometry scanning a reciprocal space
map . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 34
3.9 Schematic of the geometry for the determination of Jc(B, θ) . . . 38
4.1 Powder diffraction patter for a) pure Ba2YNbO6, b) pure YBa2Cu3O7−δ,
c) 1:1 mole mixture of YBa2Cu3O7−δ:Ba2YNbO6 heated at 950 ◦C
for 15 hr. [4] . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 42
4.2 Powder diffraction patter for Ba2YNbO6powder produced. Adapted
from [5]. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 44
4.3 X-ray diffraction data for undoped YBa2Cu3O7−δ and 5%mol Ba2YNbO6
doped film deposited on SrTiO3. Adapted from [5]. . . . . . . . . 46
4.4 X-ray diffraction data from φ scans of (101) SrTiO3, (102) YBa2Cu3O7−δ,
(202) Ba2YNbO6from a 5%mol Ba2YNbO6 doped film. Adapted
from [5]. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 48
4.5 Crystallographic matching of YBa2Cu3O7−δ with Ba2YNbO6. a)
b-axis view and b) c-axis view [5]. . . . . . . . . . . . . . . . . . . 49
4.6 Atomic Force Microscopy surface image of a pure YBa2Cu3O7−δ
and a 5%mol Ba2YNbO6 doped YBa2Cu3O7−δ thin films. a) 5
µm × 5 µm surface area of pure YBa2Cu3O7−δ; b) 1 µm × 1 µm
surface area of pure YBa2Cu3O7−δ; c) 5 µm × 5 µm surface area of
5%mol Ba2YNbO6 doped YBa2Cu3O7−δ; d) 1 µm × 1 µm surface
area of 5%mol Ba2YNbO6 doped YBa2Cu3O7−δ . . . . . . . . . . 52
xii
LIST OF FIGURES
4.7 Transmission electron microscope cross-sectional image of a 5%mol
Ba2YNbO6 doped YBa2Cu3O7−δ thin film. White arrows in the di-
rection of the c-axis mark the position of self assembled Ba2YNbO6
nanorods parallel to the c-axis. Larger white arrows parallel to the
YBa2Cu3O7−δ ab-planes mark the position of nanoparticles and of
the interface between the substrate and the thin film. Inset shows
selected area electron diffraction pattern taken from a region in
proximity of a Ba2YNbO6 inclusion showing diffraction spot of
Ba2YNbO6 YBa2Cu3O7−δ and SrTiO3. (TEM images from Prof.
H Wang research group at Texas A&M University). . . . . . . . . 54
4.8 Transmission electron microscope cross-sectional image of a 5%mol
Ba2YNbO6 doped YBa2Cu3O7−δ thin film. White arrows in the di-
rection of the YBa2Cu3O7−δ c-axis mark the position of self assem-
bled Ba2YNbO6 nanorods parallel to the c-axis. Arrows parallel
to the YBa2Cu3O7−δ ab-planes mark the position of nanoparticles.
The Ba2YNbO6 nanorods presence is indicated by the Moire fringes
arising from lattice mismatch between Ba2YNbO6 and YBa2Cu3O7−δ.
(TEM images from Prof. H Wang research group at Texas A&M
University). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 57
4.9 Resistance variation with the temperature measured on a 5%mol
Ba2YNbO6 doped YBa2Cu3O7−δ thin film. The inset shows the
region of the superconducting transition magnified. . . . . . . . . 59
4.10 Critical current density variation with the applied magnetic field
value measured on a pure YBa2Cu3O7−δ and a 5%mol Ba2YNbO6
doped YBa2Cu3O7−δ thin films at 77 K and field applied parallel
to the c-axis of the YBa2Cu3O7−δ. The inset shows data on a
logaritmic axis and the calculated α values. . . . . . . . . . . . . . 61
4.11 Critical current density variation with the direction of the applied
magnetic field measured on a pure YBa2Cu3O7−δ and a 5%mol
Ba2YNbO6 doped YBa2Cu3O7−δ thin films at 77 K and magnetic
flux density = 0.5 T . . . . . . . . . . . . . . . . . . . . . . . . . 62
xiii
LIST OF FIGURES
5.1 X-ray diffraction reciprocal space map measured on a pure YBa2Cu3O7−δ
thin film. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 70
5.2 X-ray diffraction reciprocal space map measured on 5%mol Ba2YNbO6
doped YBa2Cu3O7−δ thin films. a) Film deposition rate of 1 Hz;
b) Film deposition rate of 10 Hz. . . . . . . . . . . . . . . . . . . 72
5.3 Half maximum intesity insoline and peak position comparison for
the (108)(018) YBa2Cu3O7−δ peaks from figure 5.1 and 5.2. . . . 75
5.4 Atomic Force Microscopy surface image (5 µm × 5 µm) of 5%mol
Ba2YNbO6 doped YBa2Cu3O7−δ thin films deposited at a sub-
strate temperature of 780 ◦C, repetition rate of 1 Hz and 10 Hz
with a thickness of 700 nm and 400 nm. a) 1 Hz and 700 nm; b)
10 Hz and 700 nm; c) 1 Hz and 400 nm; d) 10 Hz and 400 nm. . . 77
5.5 Transmission electron microscope cross-sectional image of a 5%mol
Ba2YNbO6 doped YBa2Cu3O7−δ thin film deposited with a laser
pulse rates of 1 Hz (a, b) and 10 Hz (c, d). White arrows in the di-
rection of the YBa2Cu3O7−δ c-axis mark the position of self assem-
bled Ba2YNbO6 nanorods parallel to the c-axis. Arrows parallel
to the YBa2Cu3O7−δ ab-planes mark the position of nanoparticles.
(TEM images from Prof. H Wang research group at Texas A&M
University). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 80
5.6 Resistance variation with the temperature measured on a pure
YBa2Cu3O7−δ thin film deposited at a substrate temperature of
780 ◦C and laser pulses repetition rates of 1 Hz and on 5%mol
Ba2YNbO6 doped YBa2Cu3O7−δ thin films deposited at a sub-
strate temperature of 780 ◦C and laser pulses repetition rates of 1,
5 and 10 Hz. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 83
xiv
LIST OF FIGURES
5.7 Critical current density variation with the applied magnetic field
value measured on a pure YBa2Cu3O7−δ thin film deposited at a
substrate temperature of 780 ◦C and laser pulses repetition rates
of 1 Hz and on 5%mol Ba2YNbO6 doped YBa2Cu3O7−δ thin films
deposited at a substrate temperature of 780 ◦C and laser pulses
repetition rates of 1, 5 and 10 Hz. The field is applied parallel
to the c-axis of the YBa2Cu3O7−δ, T = 77 K. The data are on a
log-linear plot (a) and a log-log plot (b). . . . . . . . . . . . . . . 85
5.8 Pinning force calculated from data reported in figure 5.7 . . . . . 86
5.9 Critical current density variation with the direction of the applied
magnetic field measured on a pure YBa2Cu3O7−δ thin film de-
posited at a substrate temperature of 780 ◦C and laser pulses repe-
tition rates of 1 Hz and on 5%mol Ba2YNbO6 doped YBa2Cu3O7−δ
thin films deposited at a substrate temperature of 780 ◦C and laser
pulses repetition rates of 1, 5 and 10 Hz. a) Applied magnetic flux
density = 0.5 T, T = 77 K; b) Applied magnetic flux density = 1
T, T = 77 K. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 88
5.10 Atomic Force Microscopy surface image (5 µm × 5 µm) of 5%mol
Ba2YNbO6 doped YBa2Cu3O7−δ thin films deposited at a sub-
strate temperature of 740 ◦C; a) repetition rate of 1 Hz; b) repeti-
tion rate of 5 Hz; c) repetition rate of 10 Hz; . . . . . . . . . . . . 93
5.11 Atomic Force Microscopy surface image (5 µm × 5 µm) of 5%mol
Ba2YNbO6 doped YBa2Cu3O7−δ thin films deposited at a sub-
strate temperature of 760 ◦C; a) repetition rate of 1 Hz; b) repeti-
tion rate of 5 Hz; c) repetition rate of 10 Hz; . . . . . . . . . . . . 94
5.12 Atomic Force Microscopy surface image (5 µm × 5 µm) of 5%mol
Ba2YNbO6 doped YBa2Cu3O7−δ thin films deposited at a sub-
strate temperature of 780 ◦C; a) repetition rate of 1 Hz; b) repeti-
tion rate of 5 Hz; c) repetition rate of 10 Hz; . . . . . . . . . . . . 95
5.13 Average grain size measured from AFM analysis of 5%mol Ba2YNbO6
doped YBa2Cu3O7−δ thin film deposited at a substrate tempera-
ture of 740, 760 and 780 ◦C as a function of laser pulse rate . . . 97
xv
LIST OF FIGURES
5.14 Visualisation of the transition temperature values measured for
5%mol Ba2YNbO6 doped YBa2Cu3O7−δ thin films deposited at
a substrate temperature of 740, 760 and 780 ◦C and laser pulses
repetition rates of 1, 5 and 10 Hz . . . . . . . . . . . . . . . . . . 98
5.15 Critical current density variation with the applied magnetic field
value measured on 5%mol Ba2YNbO6 doped YBa2Cu3O7−δ thin
films deposited at a substrate temperature of 760 and 780 ◦C and
laser pulses repetition rates of 1, 5 and 10 Hz. The field is applied
parallel to the c-axis of the YBa2Cu3O7−δ, T = 77 K. . . . . . . . 100
5.16 Critical current density measured at 0.5 T and 1 T on 5%mol
Ba2YNbO6 doped YBa2Cu3O7−δ thin films deposited at a sub-
strate temperature of 760 and 780 ◦C as a function of laser pulse
rate. The field is applied parallel to the c-axis of the YBa2Cu3O7−δ,
T = 77 K. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 101
5.17 Critical current density variation with the direction of the applied
magnetic field measured on 5%mol Ba2YNbO6 doped YBa2Cu3O7−δ
thin films deposited at a substrate temperature of 760 and 780 ◦C
and laser pulses repetition rates of 1, 5 and 10 Hz. a) Applied
magnetic flux density = 0.5 T, T = 77 K; b) Applied magnetic
flux density = 1 T, T = 77 K. . . . . . . . . . . . . . . . . . . . . 102
6.1 Bragg-Brentano scan of a YBa2Cu3O7−δ + 2.5%mol Ba2YNbO6 +
2.5%mol Gd3TaO7 thin film deposited on SrTiO3. . . . . . . . . . 107
6.2 X-ray diffraction data from φ scans of (102) YBa2Cu3O7−δ, (202)
Ba2YNbO6 and (404) (Y/Gd)2O3 from a 2.5%mol Ba2YNbO6 +
2.5%mol Gd3TaO7 doped YBa2Cu3O7−δ thin film. . . . . . . . . . 109
xvi
LIST OF FIGURES
6.3 a) TEM image of a YBa2Cu3O7−δ + 2.5%mol Ba2YNbO6 + 2.5%mol
Gd3TaO7 thin film cross section; b) TEM image of plate-like nanopar-
ticleof (Y/Gd)2O3 nucleated between two Ba2(Y/Gd)(Nb/Ta)O6
nanorods segments; c) Selected area electron diffraction pattern
of the image in (a); d) TEM imafe of plate-like nanoparticle of
(Y/Gd)2O3, inset (d) Fourier transform of the particle image. (TEM
images from Prof. H Wang research group at Texas A&M Univer-
sity). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 110
6.4 Critical current density variation with the applied magnetic field
value measured on a pure YBa2Cu3O7−δ, a 5%mol Ba2YNbO6
doped YBa2Cu3O7−δ, a 2.5%mol Ba2YNbO6 - 2.5%mol Gd3TaO7
doped YBa2Cu3O7−δ and a BaZrO3 doped YBa2Cu3O7−δ thin films
[6]. The field is applied parallel to the c-axis of the YBa2Cu3O7−δ,
T = 77 K. the data are on a log-linear plot (a) and a log-log plot
(b). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 113
6.5 Critical current density variation with the direction of the ap-
plied magnetic field measured on a pure YBa2Cu3O7−δ, a 5%mol
Ba2YNbO6 doped YBa2Cu3O7−δ, a 2.5%mol Ba2YNbO6 - 2.5%mol
Gd3TaO7 doped YBa2Cu3O7−δ and a BaZrO3 doped YBa2Cu3O7−δ
thin films [6]. The field is applied parallel to the c-axis of the
YBa2Cu3O7−δ. a) Applied magnetic flux density = 1 T, T = 77
K; b) Applied magnetic flux density = 3 T, T = 77 K. . . . . . . 116
6.6 a) High field critical current density variation with the direction
of the applied magnetic field measured on a 2.5%mol Ba2YNbO6
- 2.5%mol Gd3TaO7 doped YBa2Cu3O7−δ. Applied magnetic flux
density = 3 T to 6 T, T = 77 K; b) sketch of vortices interacting si-
multaneously with nanorod segments along the c-axis and intrinsic
and extrinsic defects along ab-planes. . . . . . . . . . . . . . . . . 118
xvii
Chapter 1
Introduction
Superconductivity, a state of matter with a century of history. In the introduction
of this dissertation is presented a road map of the main scientific milestones
in superconductivity, from its discovery in 1911 to the latest high temperature
superconductor discovered.
1.1 The discovery of superconductivity
First discovered by the Nobel prize winner H. Kamerling Onnes [7], the supercon-
ductive state is characterized by one of the most intriguing properties of matter,
the capacity to transport an electric current without the appearance of any re-
sistive process. In other words the possibility to transfer and transform energy
with reduced losses.
H. K. Onnes was awarded the Nobel Prize in Physics in 1913 “for his inves-
tigations on the properties of matter at low temperatures which led, inter alia, to
the production of liquid helium”. In 1908 he was the first scientist able to liquefy
helium [8]; in 1911, during his following investigation on the properties of mat-
ter at low temperature, he wrote about The Disappearance of the resistance of
mercury. He discovered that the resistance of mercury becomes practically zero
when this is cooled below 4.2 K [9]. Onnes realized that zero resistance was a new
property that would have characterized a different state of matter, as a matter
of the fact in his Nobel lecture he said “the mercury at 4.2 K has entered a new
1
1. Introduction
state, which owing to its particular electrical properties, can be called the state of
superconductivity.”
1.2 The Meissner effect and the London pene-
tration depth
Another milestone in the superconductivity history was the discovery of the
Meissner effect, in 1933 Walther Meissner and Robert Ochsenfeld, during their
researches on the magnetoelectric properties of matter in very low temperatures,
observed a sudden change in current distribution and in magnetic induction at
the beginning of superconductance. They discovered that tin and lead samples
expel the magnetic field when cooled below the transition temperature, in other
words when the transition from normal to superconductive state occurs [10].
One of the main implications of the Meissner effect is that superconductivity
is something different from perfect conductivity. If two different samples, a su-
perconductor and an ideal perfect conductor, are cooled in their zero resistance
state and after that an external magnetic field is applied, both the sample would
react to the external magnetic field in the same way. The rise of Faraday shielding
superficial currents would prevent the penetration of the external magnetic field
in the samples. On the other hand if the same two samples are cooled in their
zero resistance state while the field is applied, the magnetic field distribution
would not change in the perfect conductor, because there would be no induced
electromotive force, instead the superconductor will expel the flux also in this
second scenario (figure 1.1).
A superconducting phase does not allow magnetic flux density to exist whether
it is field cooled or not, even a magnetic flux preexisting the transition will be
expelled.
A few years later, in 1935, a theoretical description of the Meissner effect
was found by Fritz and Heinz London [11]. They were able to relate the current
distribution to the electromagnetic field. A key factor in the London brothers
research was the definition of the magnetic flux penetration depth. Flux expulsion
is related to the presence of screening currents. Since only a finite current density
2
1. Introduction
Figure 1.1: Differences in magnetisation behaviours of a perfect conductor and a
superconductor, the Meissner effect. a) Perfect conductor cooled in zero field; b)
Perfect conductor cooled in applied magnetic field; c) Superconductor cooled in
zero field; d) Superconductor cooled in applied magnetic field.
3
1. Introduction
can be carried in a superconductor, a surface layer in which the screening currents
flow has to have a finite thickness. For this reason the magnetic flux density
does not decay abruptly at the superconductor surface, but decays gradually
penetrating the material. The flux penetration depth, λ, is commonly called the
London penetration depth. In simple geometries λ is the characteristic length
over which the field density decays exponentially.
1.3 Critical magnetic field, critical current den-
sity. Two types of superconductors
From an experimental point of view it was clear that the superconductive state
disappears when a sufficiently large magnetic field is applied. The field at which
the superconductivity is destroyed is called the critical magnetic field, and it is
usually denoted as Hc. This field value is strictly related to another fundamental
value, the critical current density. If an increasing magnetic field is applied to a
superconductor the shielding currents will also increase to produce a zero mag-
netic flux density, Hc is the value at which the shielding current density reaches
the critical value (Jc). When the applied external field strength is equal to Hc the
superconductor is incapable of carrying any transport current, this is due to the
fact that the total current is the sum of both shielding and transport currents thus
if the shielding current has reached the Jc no others current can be added without
exceed the critical current density and therefore destroy the superconductivity.
In 1950 Vitaly Lazarevich Ginzburg and Lev Davidovicˇ Landau developed a
mathematical phenomenological theory to model superconductivity [12]. The the-
ory, describes the macroscopic phenomena in superconductivity relying on general
thermodynamics arguments. Despite the lack of explanation of the microscopic
mechanisms of superconductivity this theory was fundamental to understand the
differences between two main classes of superconductors. Based on the general
theory of the second order phase transitions proposed by Landau in 1937, the
main variable of the model is the “order parameter” ψ which has a finite value
below the transition and zero above it. One of the key results of the model is
the definition of two characteristic length: the coherence length and penetration
4
1. Introduction
depth. The coherence length ξ, expresses the distance over which the order pa-
rameters can differ by a significant amount, in other words represents the distance
over which a superconductive phase can become normal. The penetration depth
is the characteristic length over which the field density decays, the same char-
acteristic length derived in the London theory. The Ginzburg-Landau equations
allow the calculation of the energy associated to the formation of a boundary be-
tween a normal region and a superconducting region, thus it is possible to predict
at which ξ/λ ratio such a boundary is thermodynamically favored. The equations
of the theory depend only on the dimensionless material constant, κ, defined as
the ratio between λ and ξ.
The energy associated to the formation of a boundary between a normal region
and a superconductive region is at the base of Alexei A. Abrikosov prediction of a
mixed state in which superconductivity and magnetic flux coexist. It was known
that for κ values above 1/
√
2 the surface energy between the superconducting
and normal layers would have been negative. Abrikosov decided to solve the
Ginzburg-Landau equations for κ values above 1/
√
2. He found that the magnetic
flux and superconductivity could indeed coexist as an array of cylinders of normal
conducting material parallel to the field direction surrounded by shielding current
vortices in a matrix of superconducting material, the geometric features of this
array strongly depend on temperature and field strength [13]. He stated that “in
the case κ >> 1 (λ >> ξ) every vortex has a “core” of size ξ, where the order
parameter varies rapidly, and the outer region of the size λ where the magnetic
field decays to zero” [14] (figure 1.2). Between current vortices a repulsive force
arises and the normal core are usually ordered in triangular or square lattice
[15,16].
The last outstanding result of Abrikosov’s calculations was that the magnetic
flux passing through an area bounded by a superconducting electrical current is
quantized and that the quantum of magnetic flux is a constant φ0, it is indepen-
dent from the material as long as it is superconductive. This quantized flux is
normally referred as a fluxon (Φ0 ≈ 2.067× 10−15 Wb).
The presence of the mixed state is confirmed by the magnetization behaviour
of type II superconductors that is different from that of type I. A type II super-
conductor has two different critical field strength values: the lower critical field,
5
1. Introduction
Figure 1.2: Structure of an isolated Abrikosov vortex [1]
Hc1, that is the value below which the Meissner effect is present and the entire
magnetic flux is expelled, and the upper critical field, Hc2, that is the value above
which superconductivity is destroyed. For fields value between Hc1 and Hc2 a type
II superconductor is in the mixed state predicted by Abrikosov (figure 1.3).
Figure 1.3: Magnetisation curves for a type I and a type II superconductor [1]
6
1. Introduction
1.4 Finally a microscopic theory of supercon-
ductivity
Since the discovery of superconductivity most of the brilliant minds of the twen-
tieth century theoretical physics tried to find a microscopic theory to explain
the phenomenon and failed. Almost fifty years passed from the the discovery of
the phenomenon to the publication of a microscopic theory of superconductivity.
In 1957 Bardeen, Cooper and Schrieffer published their theory on a paper en-
titled “Theory of Superconductivity” [17], the theory, known as the BCS theory
after the scientist initials, explain “a second-order phase transition at the critical
temperature, Tc an electronic specific heat varying as exp(−T0/T ) near T = 0K
and other evidence for an energy gap for individual particle-like excitations, the
Meissner-Ochsenfeld effect (B = 0), effects associated with infinite conductivity
(E = 0), and the dependence of Tc on isotopic mass, Tc
√
M = const.”
The theory is based on the coupling of electrons in Cooper pairs [18] due
to phonon-electron interactions. The interactions between the vibrations of the
lattice and the electrons generate an attractive force, when this interaction over-
come the Coulomb repulsive interaction the coupling is probable. In the BCS this
coupling of the electron leads to the formation of a gap in the electrons energy
band, thus any further interactions with an associated energy lower than the gap
can not occurs. In a superconductor below Tc the scattering of coupled electron
by the lattice (dissipative process) has an associated energy that is lower than
the gap, thus it does not happen.
The BCS theory is a rigorous quantum mechanical description of the super-
conductive phenomenon, a more detailed description of the theory would require
a discussion which goes beyond the intent of this chapter.
1.5 The high temperature superconductors rev-
olution
Since Onnes discovery only metallic materials (mainly type I superconductors)
and metal alloys (mainly type II superconductors) were thought to be supercon-
7
1. Introduction
ductive. From 1986 onward an acceleration in the scientific achievements leaded
to discovery of different families of superconductors.
In 1986 Johannes Georg Bednorz and Karl Alexander Mu¨ller discovered high
Tc superconductivity in the Ba-La-Cu-O system [19], a year later they were
awarded with the Noble prize in Physics “for their important breakthrough in
the discovery of superconductivity in ceramic materials”. The discovery of the
first cuprate based ceramic superconductor with a Tc of 30 K obtained a con-
siderable interest and started an incredible and productive research race. The
result of this intellectual spring was the achievement of Tc above 100 K that were
considered impossible only a few years before. In 1987 the Y-Ba-Cu-O system
was discovered, Tc = 93 K [20], was the first material to be in a superconduc-
tive state above the liquid nitrogen boiling temperature. A year later two other
cuprates were discovered to be superconductive at even higher temperatures, the
Bi-Sr-Ca-Cu-O with a Tc of 108 K [21] and the Tl-Ca/Ba-Cu-O with a Tc of 120
K [22]. These materials are hard to process, but the high-Tc together with the
high Hc and Jc made of these materials the ideal candidates for most practical
applications. However, despite the large amount of research and published work,
a complete theoretical knowledge of this class of superconductors is still missing,
in particular the superelectron condensation mechanism is still unknown.
In the last few years other two non-cuprate materials were discovered to be
superconductive and deserve a mention for the interest that the scientific commu-
nity has shown to these newcomers. Magnesium Diboride, Tc of 39 K, discovered
in 2001 [23] is interesting for the low-cost of the raw material and for the cable
manufacturing process easy and inexpensive when compared to that of cuprates
based cables [24–27]. And last, a new family of iron-based superconductors was
discovered in 2006 [28], as with the Bednorz-Mu¨ller discovery of cuprates this
discovery started a productive race in the search of similar compounds with bet-
ter properties [29–32]. However in this case after four years the compound with
the highest Tc of the family is still below the liquid nitrogen boiling point: the
samarium-doped Sr-Fe-As-F has a Tc of 56 K [33].
8
Chapter 2
Pinning in YBa2Cu3O7−δ thin
films
YBa2Cu3O7−δ is undoubtedly one the most promising materials to use in high
temperature superconducting cables. Companies and research centers are cur-
rently working and studying to achieve better performances, lower production
costs and a wider understanding of the physical processes that limit loss less
current transport capability.
2.1 Cuprates
The two most studied superconductors based on cuprate perovskites are the
YBCO (YBa2Cu3O7−δ) and the BSCCO system, the most common supercon-
ductors used from the BSCCO system are the Bi-2212 (Bi2Sr2CaCu2O8) and the
Bi-2223 (Bi2Sr2Ca2Cu3O10). Cuprates are considered as extreme type II super-
conductors and show an intrinsic anisotropy, a layered structure and an extremely
small coherence length ξ of the order of a few nm. The small ξ implies a large
upper critical field, considering that at Hc2 the normal cores are as densely packed
as the coherence length limit allows and that each core carries a quantized flux
φ it is possible to estimate the upper critical field value as:
Hc2 =
φ
piξ2
(2.1)
9
2. Pinning in YBa2Cu3O7−δ thin films
On the other hand disorder on the atomic scale can influence the supercon-
ductive properties.
2.1.1 Weaklinks and the texturing of cuprates
Despite the large critical current densities measured for YBa2Cu3O7−δ single crys-
tals, or for epitaxial YBa2Cu3O7−δ films on single crystals (Jc > 106MAcm−2 at
4.2 K) [34, 35] one of the main problems that kept YBa2Cu3O7−δ, away from
the production of superconducting cables were the low Jc values measured and
in general the poor connectivity in polycrystalline samples [36]. Soon it became
clear that the critical current densities across grain boundaries are a function of
the misorientation angle of the grains’ crystalline lattices. The ratio of the grain
boundary critical current density to the critical current density in the grains
varies by two orders of magnitude, changing from almost 1 when the grain are
well aligned to a drastically reduced value of about 1
50
when the grains tilt angle
is above 20 [37].
The BSCCO cuprates show similar behaviours [38], however a self texturing
process, that aligns the grains when the BSCCO wires are produced by a simple
“powder in tube” method, minimize the problematic related to the weaklinks of
the cuprates. The self grain alignment is the reason why large critical current
densities were “easily” obtained in the BSCCO conductors [39] while, when sim-
ilar method were applied to the YBa2Cu3O7−δ, only rather low critical current
densities were achieved [40,41]. Due to this fact, despite the higher Jc, the lower
anisotropy and the better in-field performance of the epitaxial YBa2Cu3O7−δ
films [35], the first cuprate based superconductors to enter the market were the
BSCCO tapes.
Today a second generation of coated conductors is entering the scene. These
new high temperatures superconducting cables are based on YBa2Cu3O7−δ. The
texturing issues were solved adopting two new production techniques: epitax-
ial YBa2Cu3O7−δ on highly textured buffer layers deposited by using ion beam
assisted deposition (IBAD) [42] and epitaxial YBa2Cu3O7−δ on rolling-assisted
biaxially-textured substrates (RABiTS) [43]. In both the process the
YBa2Cu3O7−δ is deposited epitaxially on a textured buffer layer. In the IBAD
10
2. Pinning in YBa2Cu3O7−δ thin films
technique the texturing is obtained during the buffer layer deposition while in the
RABiTS the substrate (usually a nickel alloy) is textured by a thermo-mechanical
treatment before the buffer layer deposition [44].
Once the production of long length textured YBa2Cu3O7−δ conductors was
achieved the YBa2Cu3O7−δ has become the cuprate in use in new coated super-
conductors, thus an optimisation of the material properties has become one of
the most important topics.
2.2 YBa2Cu3O7−δ
An interesting and complex crystal structure underlie the excellent YBa2Cu3O7−δ
properties. For this reason any optimal description of the material has to start
with a description of the crystalline structure.
2.2.1 Crystal structure
Figure 2.1: a) Crystal Structure of YBa2Cu3O7−δ; b) CuO2 planes; c) CuO chains.
11
2. Pinning in YBa2Cu3O7−δ thin films
The YBa2Cu3O7−δ crystal structure is described in figure 2.1a, it is an or-
thorhombic, distorted, oxygen deficient perovskite [45].
The lattice parameters of this orthorhombic structure are a = 0.3822 nm, b =
0.3891 nm and c = 1.1677 nm [46,47]. The lattice difference between a and b is
small enough to allow epitaxial growth of YBa2Cu3O7−δ on cubic substrates. As a
result a large number of twin boundaries are usually formed [48,49], therefore the
results of most of the macroscopic measurements are an average of the properties
measured over the a and b directions. However since the anisotropy within the
ab-plane is small usually an uniaxial anisotropy ab-plane c-axis is assumed.
Another usual description of the crystal structure is focused on the Cu-O
system. In this description YBa2Cu3O7−δ is pictured as a layered structure made
of two distorted planes of Cu-O2 separated by Y
+3 ions, and Cu-O copper chains
coordinated with Ba+2 ions at the edge of the unit cell. The Cu-O2 copper planes
are evidenced in figure 2.1b and the Cu-O copper chains in figure 2.1c.
2.2.2 Oxygen deficiency
Oxygen content is a key factor in YBa2Cu3O7−δ as it determines to the structure
and properties of the materials [2, 50, 51]. A detailed study of the effects of the
oxygen content on the superconductive properties as well as the crystal structure
was published in 1990 by Cava R J et al. [2].
The key results of their work was the discovery that on changing the oxygen
content form 7 to 6 (δ changes from 0 to 1) the YBa2Cu3O7−δ shows changes in
Tc from 92 K to 60 K followed by the disappearance of superconductivity (figure
2.2a) as well as a structural transition from orthorhombic to tetragonal (figure
2.2b). Furthermore Cava R J et al. found that the oxygen content variation
involves only the oxygen sites along the Cu-O chains.
A direct consequence of this knowledge is that good oxygenation is crucial to
achieve high Tc and that any effort to improve the YBa2Cu3O7−δ performance
has to take into account the importance of oxygen content. As a matter of fact
most of the YBa2Cu3O7−δ production processes include an annealing step in a
concentrated oxygen environment.
12
2. Pinning in YBa2Cu3O7−δ thin films
Figure 2.2: a) Superconductive transition temperature Tc for ten samples of
Ba2YCuOx with varying x values; b) Refined crystallographic cell parameters
for Ba2YCuOx with varying x values; from [2].
13
2. Pinning in YBa2Cu3O7−δ thin films
2.3 Pinning
In YBa2Cu3O7−δ the loss-less current transport capabilities are usually also lim-
ited by the mobility of the Abrikosov vortices. More deeply the critical current
density is determined by complex interactions between the transport currents and
the induced magnetic flux lines and the interactions between the flux lines and
the nanostructure of the materials.
Vortex phenomena is a complex field of solid state matter which is still evolv-
ing. The aim of this section is to give a brief introduction to the phenomenon
in order to understand the possible interactions, and the improvements that an
engineered nanostructure can induce on the critical currents. The following brief
introduction is based on an exhaustive description of the phenomenon that was
published in 1994 by Blatter G. et al. [52].
2.3.1 The dissipative phenomenon
When a transport current is applied, the flux lines start to move. The movement
is due to the action of the Lorentz force (note that c is a constant and η is the
friction coefficient):
FL = j ∧B/c (2.2)
fL = (Φ0/c)j ∧ n (2.3)
In the absence of external flux pinning mechanisms the only counter force is
the friction force:
Fη = −ηv (2.4)
At the equilibrium v is the steady-state velocity and FL = Fη thus v can be
derived as:
v = j ∧B/cη (2.5)
The flux motion generate a finite electric field E:
14
2. Pinning in YBa2Cu3O7−δ thin films
E = B ∧ v/c (2.6)
A finite electric field coexisting with a current density generate a dissipation,
E and j are parallel thus the power dissipated is:
P = (j ∧B)2/c2η (2.7)
In a scenario where the friction force is the only force opposing the Lorentz
force it would be impossible to realize a dissipation-free current flow.
2.3.2 Dissipation-free current flow, the pinning force
To achieve a dissipation-free current flow, the flux lines have to be pinned in
place. In other words v has to be equal 0 also when FL 6= 0. An additional static
force, a pinning force Fpin, has to counter the Lorentz force and be active when
the flux lines are not moving (v = 0).
The flux lines interact with any defects in the lattice. Since the formation of
a normal region leads to the loss of the condensate state increasing the energy
associated to the system, any region of the material in which the superconducting
order parameter is already depressed constitute a region in which the presence of
a flux line is energetically favourable. Therefore any region in which the super-
conducting order parameter is depressed will contribute to a finite pinning force
density Fpin.
The critical current density is the current at which the Lorentz force is equal
to the pinning force since any further increment of the current flow would lead
to an imbalanced increment of the Lorentz force (FL > Fpin) and subsequently a
flux motion (v 6= 0) and power dissipation. Therefore assuming FL = Fpin and
j⊥B, Jc can be derived as:
jc = cFpin/B (2.8)
In order to increase Jc it is necessary to increase the pinning force Fpin.
15
2. Pinning in YBa2Cu3O7−δ thin films
2.3.3 Defects as pinning centers
A single flux line can interact with different class of defects [53, 54]. Each class
is characterised by a characteristic dimension. Defects smaller than ξ in each
dimension are called point defects and the saved condensation energy is propor-
tional to the defect volume. In anisotropic materials, since ξ is a function of
direction, the pinning force provided by this class of defects is also a function of
the force direction. Defects with one dimension larger than the others are named
linear defects while defects with two dimensions larger than the other are planar
defects. In high temperature superconductors naturally occurring defects like dis-
location are linear defects and act like pinning lines while twin planes, stacking
faults are planar defects and act as pinning planes. In YBa2Cu3O7−δ and cuprates
in general the layered crystal structure and the consequent ordered stacking of
layers with different superconducting order parameter are an additional array of
pinning planes parallels to the ab-planes of the crystal figure 2.1.
Linear and planar defects, as well as defects with specific spatial arrangements
can lead to directional pinning, increasing the pinning potential with respect to a
specific direction of the applied magnetic field. As an example the pinning force
provided by a dislocation to a flux line will be higher if the flux line and the
dislocation are parallel. In the same way, an ordered array of nanoparticles can
act as a linear defect providing the higher pinning potential to flux lines that are
able to penetrate the sample in the same direction of the array.
In conventional YBa2Cu3O7−δ thin films the Jc is the highest when measured
with the applied magnetic field parallel to the ab-planes of the crystal. This is
a direct consequence of the directional planar pinning provided by the specific
cuprates layers.
A last thing to note is that, even if the single interactions between flux lines
and defects are at the base of the material response, in most field values and
temperature conditions (especially high temperatures) the materials properties
are given by the simultaneous interactions of many flux lines with many different
defects [55,56]. For this reason a theoretical approach to describe the overall ma-
terial behaviour is usually too complex, thus an experimental approach is usually
more productive when dealing with pinning engineering. Nevertheless the theo-
16
2. Pinning in YBa2Cu3O7−δ thin films
retical achievements greatly enhance the understanding of the physics processes
underling the complex behaviours of pinning engineered enhanced superconduc-
tors.
2.4 Practical Pinning Engineering
Defects engineering of YBa2Cu3O7−δ to increase the Jc values has been and is
today one of the hot topics of superconductive materials science. Large electri-
cal current transport in absence of energy losses is the key factor in commer-
cial application of high temperature superconductors. Since energy losses arise
from vortices movements a solution has been researched in nanostructured defects
landscape capable to pin the vortices in place. The higher is the efficiency of the
pinning landscape the higher is the current value threshold at which the losses
appears.
Different pinning landscapes and different techniques have been adopted dur-
ing the years by researchers, in the following sections an overview of the most
interesting works is given.
2.4.1 Non superconducting secondary phase addition
The secondary phase addition is an attractive way to introduce defects. It can be
adopted both in films grown by physical processes and in films grown by chemical
processes.
The first example of pinning engineering in YBa2Cu3O7−δ thin films by the
introduction of a non superconducting epitaxial second phase is the addition of
BaZrO3 [57]. MacManus-Driscoll et al. in 2004 were able to produce high quality
YBa2Cu3O7−δ thin films with improved Jc (up to a factor of 5) in the magnetic
field by introducing BaZrO3 powder in a Pulsed Laser Deposition (PLD) target of
YBa2Cu3O7−δ. They achieved an epitaxial growth of a cubic non-superconducting
phase in an epitaxial grown high quality YBa2Cu3O7−δ thin film. An improvement
of both the random and directional pinning was obtained. In particular a large
increment of Jc was measured with the magnetic field applied parallel to the c-
axis of the film. Furthermore, by using TEM (Transmission Electron Microscopy)
17
2. Pinning in YBa2Cu3O7−δ thin films
analysis, c-axis oriented columnar defects were found evidencing the correlation
between the defects landscape and the pinning potential. This pioneering work
is the first of a long series of research articles aiming to produce the ideal pinning
landscape adopting a standard industrial ready deposition technique and avoiding
technological complications [58–63].
BaZrO3 is not the only secondary phase to produce columnar defects, during
the years also BaSnO3 [64–66], RE3TaO7 [67], Ba2YTaO7 [68] where found to pro-
duce similar columnar defects. All these phases produce c-axis oriented columns,
but the mean width, the spacing and the linearity of the columnar defects pro-
duced varies. The variation is dependent on both the phase and the process
parameters adopted. A common feature of all this heteroepitaxial columnar de-
fects is the perovskite crystal structure with the only exception of a pyrochlore
phase (RE3TaO7 [67]).
In this work the same technique is used to produce nanostructured
YBa2Cu3O7−δ thin film with superior pinning properties adopting a Nb based
perovskite as secondary phase as well as in a Nb and Ta simultaneous addition
that will result in the complex ideal pinning landscape that will be described later
in the thesis.
When giving also a small overview of the milestones in the YBa2Cu3O7−δ
pinning engineering fields the achievements obtained in chemical solution nanos-
tructured YBa2Cu3O7−δ films have to be reported. Gutierrez et al., in 2007, were
able to obtain a dense nanodispersion of defects in YBa2Cu3O7−δ thin films by
introducing BaZrO3 [69]. In their work an entirely chemical process was adopted
and randomly oriented BaZrO3 nanoparticles were reported as the basis of a
strong isotropic pinning.
Another widely used technique used to produce pinning enhanced
YBa2Cu3O7−δ is the multiple target ablation in pulsed laser deposition [70–73].
A secondary phase can be introduced by ablating multiple targets in a so called
pseudo-multilayer deposition [70,71]. Haugan et al. were able to improve the pin-
ning properties of YBa2Cu3O7−δ by introducing particles of nanometric size of
YBa2CuO5 (referred as 211) by growth of alternating layers of ultra thin 221 and
YBa2Cu3O7−δ [70]. This multiple bilayer structures were obtained by ablating in
sequence a target made of pure YBa2Cu3O7−δ and a target of pure YBa2CuO5 us-
18
2. Pinning in YBa2Cu3O7−δ thin films
ing a computer driven target’s carousel. Adopting the same technique Campbell
et al. deposited Y2O3 - YBa2Cu3O7−δ pseudo-multilayer films achieving similar
results to Haugan [71].
Recently a Nb doped YBa2Cu3O7−δ was also produced adopting the multiple
target deposition [74].
Multiple target ablation, when compared to single composite target ablation,
has the advantage of having a certain degree of control over some dispersion
parameters allowing a change in the number of laser pulses for each layer (the
amount of material deposited from each target), and, in theory some control can
also be achieved on the deposition parameters adopted for each target. In prac-
tice the change of certain deposition parameters like oxygen pressure or substrate
temperature during the film deposition would over complicate the deposition pro-
cess, therefore in almost all the works adopting this technique the only parameters
that are tuned during the secondary phase target ablation are pulse frequency
and laser energy.
2.4.2 Alternative routes
Other possible routes to increase defects concentration in YBa2Cu3O7−δ thin films
have been experimented in recent years and deserve citation. The first is substrate
decoration which can be achieved adopting different techniques. The idea is to
deposit nanoparticles or nanoislands on the substrate before the YBa2Cu3O7−δ
film []. The nanoparticles can be deposited by pulsed laser deposition [75–80],
sputtering [81–84] or even with gas-phase prepared nanoparticles [85] and TFA-
MOD [86]. These particles will then induce stresses and therefore defects in the
film by introducing distortions at the substrate-film interface.
The last two techniques differ from the others as they are not based on the
addition of a secondary phase. Improved pinning can be achieved by exchanging
a certain amount of Y with different rare-earth elements [87]. Enhanced low-
field pinning is shown by mixed rare-earth barium cuprate films, while Y based
film with large radius rare-earth elements substitution have improved pinning up
to at least 7 T. Irradiation with high energy heavy ions has also been adopted
in YBa2Cu3O7−δ single crystal pinning research [88–93] and in thin film studies
19
2. Pinning in YBa2Cu3O7−δ thin films
[94–99] for its capability to introduce linear defects along the ballistic directions.
Despite the interesting and unique features of this technique its high complexity
make the irradiation process unpractical.
20
Chapter 3
Experimental Techniques
This chapter describes the experimental techniques adopted in the sample prepa-
ration and characterisation. The first section will focus on the sample deposition
followed by a description of the micro bridge patterning together with the elec-
trode deposition. The structural and morphological characterisation techniques
are described in the second section. These measurements are mainly performed on
unpatterned samples (as deposited samples), and with the only exception of the
TEM are non destructive measures. The last section of the chapter describes the
transport measures, these are the main characterisation of the superconducting
properties.
3.1 Sample Preparation
As anticipated in the previous chapter, the nanostructuring technique adopted in
this thesis is the introduction of one or more secondary phases in the YBa2Cu3O7−δ.
These additional phases are included in the films by introducing the materials
directly to the targets used in the pulsed laser deposition.
The collection of experimental techniques used to prepare the samples anal-
ysed in this work are presented in this section. Below is given a description of the
pulsed laser deposition system utilized and a discussion of the technical solutions
adopted to obtain the best possible control of the process parameters. A tight
control of process parameters is crucial in obtaining a good repeatability. The
21
3. Experimental Techniques
sintering procedure applied to the pulsed laser deposition target preparation is
described after the deposition technique while the photolithographic process per-
formed to obtain the micro bridge patterning and the electrodes deposition are
explained at the end of the section.
3.1.1 Pulsed Laser Deposition
Pulsed laser deposition (PLD) is a valuable choice for depositing extremely pure
films because it reproduces exactly the target composition, multilayer materials
can be done rather easily, it is the fastest route to prototyping any thin film
coating and, most importantly, very high quality YBa2Cu3O7−δ thin films can be
easily produced. As a matter of the fact of its extreme versatility, a considerable
amount of research work in YBa2Cu3O7−δ pinning engineering has been carried
out adopting pulsed laser deposition as deposition technique [100–104].
Pulsed laser deposition is one of the physical vapor deposition (PVD) adopted
to grow high quality thin films. A target of the desired composition is ablated by
a high power pulsed laser, the material ablated from the target forms a plasma
plume of energized ions that after traveling though a controlled atmosphere are
deposited on an heated substrate forming a thin coating film. Once these ions
reach the film surface thay are usually referred to as adatoms. In particular
adatom is the term used to describe a single atom adsorbed on a surface. To
obtain high quality thin films it is necessary to control several parameters during
the deposition. These parameters are the substrate temperature, the oxygen
pressure in the chamber, the laser energy and the frequency of the laser pulses
[105,106].
All the samples in this work were deposited in a YBa2Cu3O7−δ dedicated de-
position system to minimize contamination issues. This system is equipped with
a Lambda Physik KrF excimer laser (λ = 248 nm, fluence = 2 mJcm−2) and the
laser beam is admitted to a ultra high vacuum (UHV) chamber through a quartz
window. Periodic cleaning of the window as well as a constant monitoring of the
energy admitted in chamber ensure homogeneity of the ablating energies over the
time. Others quartz windows are mounted on the sides of the UHV chamber and
are generally used for growth and plume monitoring, secondary substrate surface
22
3. Experimental Techniques
temperature monitoring and for laser-target-substrate alignment.
A focal lens is positioned outside the chamber and it is used to focus the laser
on the targer in order to form a 2 mm x 3 mm rectangular spot on the target
surface. The laser spot geometry influence the plume geometry: a large spot
generates a small plume thus a small region in which the deposition is uniform;
a small spot generates a large plume and a large region in which the deposition
is uniform but as a consequence the deposition rate is reduced. The laser spot
geometry adopted in this work allows a uniform deposition in a region ≈ 1.4 cm
x 1.4 cm at ≈ 5 cm away from the target surface and uniform films are easily
obtained on 1 cm x 0.5 cm substrates.
The substrate surface temperature is monitored using an external pyrometer
while a secondary visible light laser, aligned to the primary excimer laser, is used
during the aligning task. Figure 3.1 shows a schematic of a typical pulsed laser
deposition system.
Figure 3.1: Schematic of a typical pulsed laser deposition system [3]
The substrates used are (001) aligned single crystals of SrTiO3 (STO), these
23
3. Experimental Techniques
are glued to an heater/holder with an high thermal conducting silver paste. The
temperature of the heater is controlled using a PID controller that controls the
power supply of the heating elements measuring the internal temperature of the
heater block with a thermocouple. Since the thermocouple is inserted in the
heater the use of the external pyrometer to monitor the substrate surface tem-
perature is adopted to avoid substrate temperature variations between different
depositions. This variations could rise from unavoidable differences in the ther-
mal conductivity between the heater and the substrate. A variable ∆T between
the thermocouple the pyrometer measure is usually present. This ∆T is mainly
due to the position of the thermocouple. The thermocouple is placed inside the
heater and at high temperatures a large energy dissipation due to radiating pro-
cesses generates large thermal gradients between the center of the heater, where
the thermocouple is positioned, and its surface, where the substrate is glued.
Furthermore the impossibility to reproduce the same thermal contact in every
deposition adds a variability to the discussed ∆T . This variable ∆T and the fact
that the substrate surface temperature is the one that determines the atoms mo-
bility, make the substrate surface temperature directly measured by the external
pyrometer the only possible choice as reference temperature.
The target is mounted on a rotating carousel equipped with a plume shutter.
The target is rotated to provide a uniform ablation. while the shutter is used
to protect the substrate from the plume during a pre-deposition ablation pro-
cess. This pre-ablation process is adopted to remove the target surface layer and
to ensure that the entire deposition is realized under similar target conditions.
The laser ablation induces modification to the surface morphology of the target
inducing a continuous fusion-recrystallization cycle. It is important to induce
these morphologic modifications before starting the thin film deposition without
allowing the plume to reach the substrate.
The vacuum system is a common two stage system equipped with a rotary
pump and turbomolecular pump, an additional multi channel mass flow controller
is also present. The rotary pump is used to pump down the chamber pressure
from ambient pressure to mid vacuum (∼ 103 mbar to ∼ 10−3 mbar) and to
provide a backing pressure ≤ 10−3 mbar to the turbomolecular pump. The tur-
bomolecular pump is used from mid vacuum to high vacuum (∼ 10−3 mbar to
24
3. Experimental Techniques
≤ 10−6 mbar). Depositions are performed in a pure oxygen atmosphere with
an oxygen pressure of 0.3 mbar. These pure oxygen deposition atmospheres are
realized by first reaching high vacuum (≤ 10−6 mbar) using the turbomolecular
pump and then venting oxygen while controlling the chamber pressure by simul-
taneously controlling the oxygen incoming flow with the mass controller and the
extraction pump power.
In a typical deposition the laser is operated in constant energy mode and the
fluence is monitored and kept at ≈ 2 mJcm−2, the pulse frequencies adopted are
between 1 and 10 Hz. Usually each deposition is performed using ∼ 4500 pulses
to obtain a film thickness of ∼ 500 nm. The distance between the target and the
substrate is ∼ 5 cm. An optimal oxygenation of the samples is obtained with
an in situ low temperature annealing of 1 hr. This step is realized cooling the
sample to ≈ 520 ◦C and raising the oxygen pressure to 500 mbar.
All thin films produced in this research work were deposited on SrTiO3 (001)
single crystal substrates. This substrate is commonly used in YBa2Cu3O7−δ epi-
taxial growth by pulsed laser deposition. It allows high quality films to be de-
posited and it was chosen for ease of comparison with previous work. In particular
the use of SrTiO3 (001) allows the comparison of the properties of the sample
produced with a large amount of existing data without introducing unnecessary
variables.
Oxygen content, grain orientation and film morphology are determinant pa-
rameters of the films superconducting properties. All these features are strongly
influenced by the deposition parameters adopted. Almost twenty years of research
in the field has provided a good knowledge of the processability windows of pure
YBa2Cu3O7−δ. High quality YBa2Cu3O7−δ thin films deposition is technological
knowledge already acquired and thus will not be discussed. An exhaustive re-
view on pulsed laser deposition of pure YBa2Cu3O7−δ thin films was published
by Singh et al. [3].
3.1.2 Target Preparation
All the pulsed laser deposition targets used to produce the samples used in
this work were prepared from sintered powders. Although commercial pure
25
3. Experimental Techniques
YBa2Cu3O7−δ targets are available it was chosen to produce pure YBa2Cu3O7−δ
targets in laboratory in order to avoid differences in the superconducting prop-
erties which may arise from different target quality rather than different targets
composition.
The process adopted to produce the targets is the same for both the pure and
the composite ones. Powder are pressed in the form of a cylindrical target and
the sintered at 950 ◦C in oxygen flow for 12hr in a dedicated tubular furnace. A
detailed scheme of the thermal process is reported in figure 3.2.
Figure 3.2: The target sintering thermal process
Pure YBa2Cu3O7−δ powder (SCI Engineered Materials 99.999%) is used in
both pure and composite targets. The secondary phases used in this work are
Ba2YNbO6 and Gd3TaO7. Targets produced by mixing pure YBa2Cu3O7−δ with
5 mol% of Ba2YNbO6 powder were adopted in the deposition of the samples
discussed in the next chapter. Targets produced by mixing pure YBa2Cu3O7−δ
with 2.5 mol% of Ba2YNbO6 powder and 2.5 mol% of Gd3TaO7 precursor were
used in the deposition of the samples discussed in the last chapter.
The Gd3TaO7 precursor mixed with the pure YBa2Cu3O7−δ and the Ba2YNbO6
in the target were 99.99% Gd2O3 and Ta2O5.
26
3. Experimental Techniques
3.1.3 Sample Patterning
While measuring the critical current density Jc a limited critical current Ic is
desired. A low Ic reduce a series of problematic side effects and technological
complication during the measure. A large Ic would require large current transfer,
this implies powerful currents sources, large conducting cables connecting the
sample and large electrodes trough which transfer the currents to the sample.
Further more since it is impossible to eliminate the contact resistances a large
transfer current passing trough any residual contact resistance will cause the
dissipation of large energies. To avoid an increment of the sample temperature,
caused by the heat generated in the dissipations at the contact, a powerful cooling
system should also be provided.
Considering equation 3.1 the easiest way to reduce Ic is the reduction of the
superconductor’s cross section.
Ic = jcSSupercond. (3.1)
A reduction of the cross section Ssupercond. simplifies the measure of the large Jc
typical of the YBa2Cu3O7−δ thin films (usually Jc > 106 MAcm−2) and is obtained
by introducing current tracks of micrometric width with a lithographic process.
A reduction of the contact resistances can be obtained with the deposition of an
electrode array through which the currents are transferred to the sample. Also
the electrodes are deposited with a lithographic process.
The electrode deposition is realized with the lift-off process described below.
First a layer of a commercial positive photoresist (AZ 4533) is spin coated on
the sample (spin at 6000 rpm for 45 s), then the resist layer is baked at 110 ◦C
for 1 minute to allow the resist network reticulation (figure 3.3b).
An intense blue light is projected on the coated sample surface using a pro-
jection mask aligner for 24 s, the mask projected at this stage is composed of two
rows of six square holes, the intense blue light is projected on the sample only
through these twelve squares.
The exposed regions of the positive photoresist become soluble to a developing
solution while the unexposed regions remain insoluble, thus after a developing
stage the two rows of six square holes printed on the mask are transferred to
27
3. Experimental Techniques
the photoresist coating layer. The developing stage is realized by the submersion
of the samples in a 3vol : 1vol H2O:AZ351B developing solution (AZ351B is a
commercial developer). The duration of the developing stage is a key factor not
always predictable: it has to be long enough to ensure that the resist layer is
completely removed in the areas that were exposed to the light but at the same
time an over lasting developing stage is to be avoided since it can cause resist
removal from the unexposed zones. In order to optimally time the duration, the
developing step is divided in several steps of reduced duration followed by the
observation of the sample surface with an optical microscope until the resist layer
is completely removed from the exposed areas (figure 3.3c).
Once the developing stage is successfully completed a silver layer and a gold
layer are deposited on the sample surface in an ion milling magnetron sputtering
hybrid system. The Ag / Au bilayer deposition is preceded by an ion milling step
to remove the passivated layer on the superconducting film ensuring the optimal
connectivity between the Ag layer and the YBa2Cu3O7−δ surface (figure 3.3d).
The last step of the lift-off process, which gives the name to the technique,
is realized by submerging the samples in acetone in an ultrasonic bath. The
acetone dissolves the photoresist layer and the resist removal causes the lift-off
(delamination) of the Ag / Au bilayer deposited on it while the metals deposited
directly on the YBa2Cu3O7−δ surface in the areas that were not covered by the
resist forms the electrodes. In this way two rows of six Ag / Au electrodes are
deposited on the films (figure 3.3e).
The track pattern transfer (the micro bridge formation) is obtained with a
process similar to the electrode deposition until the developing stage. In this
process the mask used project the light in the areas of the samples from which
is necessary to remove the superconducting material. Thus after the developing
stage, the resist layer is the exact copy of the desired tracks pattern (figure 3.3g),
the effective film etching can then by realized either by a physical or chemical
process (dry etching or wet etching). Very small features ( ≤ 10 µm) are usually
realized with a dry etching because the chemical etching could blur the features.
Since the smallest feature in this work, the tracks width, is ≈ 50 µm both
processes are equally effective. The dry etching is performed in an argon ion
milling system where energized ions removes material from the sample surfaces.
28
3. Experimental Techniques
The wet etching is realized by submerging the sample in a HCl diluted solution
(0.01M) with the acid dissolving the unprotected YBa2Cu3O7−δ from the sample
(figure 3.3h).
A schematic of the sample at the different stages of the process is reported
in figure 3.3 while a micrography showing a single track and fractions of two
electrodes in shown in figure 3.4
Five tracks (≈ 50 µm in width), each one connected to four Ag / Au elec-
trodes, are patterned on the samples using the described techniques.
3.2 Structural Morphological characterisation
In this section are listed the scientific investigation techniques adopted to charac-
terise the crystal structures and the morphology at the nanoscale of the samples.
Since these techniques are commonly adopted by materials scientists working in
thin films, as well as different fields, they will not be discussed in detail. In this
section only a presentation of the answers that can be achieved with the analysis
performed is given.
3.2.1 Structural Analysis
Phase and orientation analysis are performed using x-ray diffraction. Performing
this non destructive technique it is possible to gather information on the crystal
phases present in the thin films and their orientation. The distribution of grain
orientations is described by the crystalline texture. Random texture describes
films with randomly oriented grains, fibre texture is a term used to describe films
in which grains have one crystallographic axis parallel to the substrate normal
but are randomly oriented in plane, and epitaxial growth (or epitaxial alignment)
describes films made of grains that have a fixed orientation in plane. A scan in
Bragg-Brentano geometry will give information on the spacing of the crystalline
planes that are oriented perpendicularly to the substrate surface. The analysis
are usually performed in the geometrical configuration reported in figure 3.5, and
in this configuration only the diffraction peaks associated to the planes with the
normal oriented perpendicular to the substrate surface will be registered.
29
3. Experimental Techniques
Figure 3.3: Sketches of a sample at the different stages of a photolithographic
process. a) As deposited sample; b) Sample after the first spin coating of the
photoresist; c) Sample after the developing of the electrode mask; d) Sample
after the silver layer and gold layer deposition. e) Sample after the lift-off process;
f) Sample after the second spin caoting of the photoresist; g) Sample after the
developing of the trak mask; h) Sample after the milling/cleaning process.
30
3. Experimental Techniques
Figure 3.4: Micrography of the surface of a patterned sample
Figure 3.5: Schematic of a diffractometer operated in the Bragg-Brentano Geom-
etry
31
3. Experimental Techniques
The possible presence of an epitaxial growth is easily evidenced by the presence
of only the peaks indexed (00l) since this will be the experimental proof that the
crystalline c-axis is the only crystal axis oriented perpendicular to the substrate
surface. The analysis of the diffraction peaks gathered in this geometry allows the
determination of the crystalline composition of the samples and a first information
on the orientation of the phases. A limitation of this typology of x-ray diffraction
analysis is to only gather information of the so called out of planes diffraction.
The determination of the out of plane orientations allows to distinguish a random
texture from a fibre texture or an epitaxial growth.
In order to determine whether a film can be described by a fiber texture
or an epitaxial growth it is necessary to analyse the in plane orientation of the
crystalline phases. The in plane orientation analysis is performed with the so
called φ scan. In figure 3.6 a sketch of the geometrical configuration in which the
φ scans are done is reported.
Figure 3.6: Schematic of a diffractometer geometry operating a φ scan
In a φ scan the sample orientation χ, the tube and detectors position θ are
predetermined to focus the machinery on a specific in plane diffraction peak. The
scan is then performed varying the sample φ angle and measuring the intensity
32
3. Experimental Techniques
of the diffracted x-rays at the different φ values. In this way it is possible to asses
the in plane orientation of a given crystalline phase and determine whether that
specific crystalline phase is growing following the orientation of the substrate
crystalline axis or is growing with a determined offset or is not following any
specific orientation. This technique is valid if the phases were rightly determined
and if the peaks on which the analysis system is focused are unique.
These first two x-ray diffraction techniques described are an optimal solution
in the determination of the crystalline phases and their orientation but are limited
when determining the epitaxy quality. They give little to no information on the on
the spread of the orientation angles that is usually below the resolution of systems
adopted when operated in the geometrical configuration described. Further more
the presence of small fraction of randomly oriented crystalline phases could easily
remain undetermined by these analysis. A randomly oriented fraction of a given
oriented phase could be interpreted as noise when performing a φ scan.
Figure 3.7: Schematic of a diffractometer geometry operating a rocking curve
scan
In order to address the epitaxy quality and the strain levels rocking curves
and reciprocal space maps are gathered.
A rocking curve (ω scan) is an analysis of the angular spread of a determined
33
3. Experimental Techniques
planar direction. It is realised by scanning with a fixed 2θ (a fixed pi − 2θ) a
small windows of positive and negative offsets ω (figure 5.2). The full width half
maximum of the peaks measured gives an indication of the angular spread of the
planar direction analysed.
Figure 3.8: Schematic of a diffractometer geometry scanning a reciprocal space
map
A reciprocal space map is the detailed study of the x-ray diffraction in a
small window of 2θ and ω angles. A reciprocal space map can be obtained by
measuring a set of ω/2θ scans with different ω-offset angles around a specific
reciprocal space reflection, determined by ω and 2θ (figure 3.8). The intensity of
the diffracted signal is then plotted on a map with coordinates qz and qx that are
the reciprocal of the spacing dz and dx thus it is possible to directly calculate the
lattice parameters. The position of the diffraction peaks and their shape gives
information on the mean lattice spacing and on the strain of the phase analysed.
It is also possible to establish if a phase is relaxed or strained with respect to
the substrate. As a downside the reciprocal space maps are time consuming mea-
sures and therefore the windows around a specific diffraction angles are usually
the smallest possible. Even if in theory it would be possible to gather a complete
maps this are always avoided for the impracticality of the time required.
34
3. Experimental Techniques
3.2.2 Microscopy
The knowledge of the crystalline phases, their orientation and strains level rep-
resents only a partial knowledge of a film quality. Furthermore since the pinning
potential is related to the nanostructuration of the phases, a study of the mor-
phology on the nanometric scale is fundamental. Only by knowing the grain mor-
phology and the shape, spacing and orientation of a secondary phase it is then
possible to relate the defect landscape to the pinning properties. Furthermore it
is possible to evaluate the strength and weaknesses of a defects distribution in
order to improve the properties increasing the points of strength and decreasing
the weaknesses of a pinning landscape generating ideals defects distribution.
There are two main scientific investigation tools capable to fulfill the task:
the atomic force microscopy (AFM) and the transmission electron microscopy
(TEM).
AFM [107] is used to investigate the surface topography of the samples, from
the topography is possible to evaluate the homogeneity of a film and the roughness
of the surface. In addition to shapes and dimensions of the grain in samples
characterised by a smooth surface is also possible to investigate the presence
and distribution of secondary phases. The technique is non destructive and it is
also fast and reliable, unfortunately the only information gathered concern the
samples surface, thus only a partial information of the defects distribution can be
obtained. As an example it would be impossible from an AFM only investigation
to establish whether a secondary phase particle is a c-axis oriented defects or
a plate like ab-planes oriented particles. For these reasons it is necessary to
complete the information with TEM.
TEM gives information on the distribution and orientation of the phases in the
sample. It is a powerful tool but is a time consuming and destructive technique.
Nevertheless it is fundamental when picturing the pinning landscapes produced
in the films. Essentially with this technique are taken images of the samples
cross-section at the nanometric scale. A direct observation of the nanoparticles
introduced in the samples is obtained. Furthermore analysing the electron diffrac-
tion pattern taken with the same technique it is possible to confirm and complete
the phase and orientation analysis produced the the x-ray diffraction. A TEM
35
3. Experimental Techniques
allows to take different electron diffraction patterns from different locations of
the sample thus allows to gather separated electron diffraction pattern from the
different phases present in the sample. TEM is also adopted to perform local
strain analysis to complete the information on the average strain obtained with
the reciprocal space maps.
Cross-section TEM images were taken by JEOL 2010 analytical microscope
with a point-to-point resolution of 0.20 nm. High resolution TEM and scanning
transmission electron microscopy analysis was conducted using FEI Tecnai F20
with a point-to-point resolution of 0.18 nm. Cross-sectional samples for TEM
analysis were prepared by a standard manual grinding and thinning procedure
followed by a final ion polishing step (Gatan PIPS 691 precision ion polishing
system).
3.3 Superconductivity Properties characterisa-
tion
The final part of this chapter is dedicated to on overview of the techniques adopted
to characterize the superconducting transport properties of the sample produced.
The three main properties analysed are the transition temperature Tc, the critical
current density as a function of the applied magnetic field value Jc(B) and the
critical current density as a function of the direction of the applied magnetic field
Jc(B, θ). It is evident that the target is to achieve the highest possible Jc and
Tc and to minimize the detrimental effect of an increasing value of the applied
magnetic field on Jc. This set of measurements gives complete information on
the quality and effectiveness of the defects landscape in the creation of an ideal
pinning enhanced thin film YBa2Cu3O7−δ based superconductors.
The first measure always performed is the determination of the transition
temperature Tc. This measure is performed by a straight forward technique. A
4 point resistance measure is realized applying the smallest transport current
possible that minimize the noise while not affecting the transition temperature
(usually I ≤ 10−6 A). The resistance is measured while reducing the tempera-
ture from TAmb to 77 K. The transition temperature is then associated with the
36
3. Experimental Techniques
temperature at which the resistance disappears. The attribution of Tc is not
unique, in some work it is attributed to the temperature with the maximum of
the ∂ρ
∂T
or with the mean point of the transition, in this work the most restrictive
condition is chosen. The reason to chose the most restrictive condition for the
determination of Tc is that this condition takes into account the homogeneity of a
sample and that a sample will not be considered superconductive until a complete
superconductive link is generated between the probing electrodes.
The critical current density as a function of the applied magnetic field value
Jc(B) and the critical current density as a function of the direction of the applied
magnetic field Jc(B, θ) are measured with the same technique. The only difference
is the parameter that is changing during the measurement. In the Jc(B) measure
the direction of the applied magnetic field is constant (usually B‖ c-axis and B
‖ ab-planes) and the applied magnetic field value is increased. In the Jc(B, θ)
measurement the applied magnetic field value is constant and the direction is
changed.
The core of the measurement is the individuation of the critical transport
current Ic. To fulfill this task the current is increased in small step and a voltage
value is measured for each current value until the voltage reaches a maximum
value (V = 10−5V ), from the I−V curve obtained the current value at which the
voltage is equal to a defined criterion (Vc = 10
−6V cm−1) is derived and that value
is defined as the Ic value. Knowing the Ic and the geometry of the micro bridge
the Jc is easily calculated using equation 3.1. The derivation of the Ic value can
be achieved by adopting different technique: a linear interpolation around the
criterion voltage, a polynomial interpolation over a n-point windows around the
criterion or fitting the I − V data with the following equation:
I = Ic(
V
Vc
)n (3.2)
The different Ic derivation techniques lead to little or no variation of the value
derived. Never the less the technique adopted in the work is the fitting with the
equation 3.2 since it is more reliable in the case of noisy measures.
The Jc(B) and the Jc(B, θ) measurements are basically realised with a set of
I − V measures in different magnetic fields intensity and geometry. Both Jc(B)
37
3. Experimental Techniques
Figure 3.9: Schematic of the geometry for the determination of Jc(B, θ)
and Jc(B, θ) measurements are performed in maximum Lorentz force configura-
tion (I ⊥ B) figure 3.9.
38
Chapter 4
Ba2YNbO6 doped YBa2Cu3O7−δ:
preliminary results
In this chapter are reported the results of the study of the Nb doping of
YBa2Cu3O7−δ thin films. The effects on the nanostructure of the films as well as
the superconducting properties are discussed. The study in this chapter includes
the preliminary results obtained on Ba2YNbO6 doped YBa2Cu3O7−δ thin films
deposited by pulsed laser deposition.
4.1 The Nb introduction in the YBa2Cu3O7−δ
The introduction of a secondary phase in the YBa2Cu3O7−δ that has been de-
scribed in the previous chapter is now a widely used tool to increase the pinning
properties of commercial conductor as well as research labs advanced materials.
When introducing new pinning additions to YBa2Cu3O7−δ adopting the pulsed
laser deposition technique, it is important to consider that the phase which forms
may not be the one which is added to the YBa2Cu3O7−δ in the target. The
phase that forms is the most thermodynamically and epitaxially stable. As a
matter of the fact if Zr4+ is added to the YBa2Cu3O7−δ the phase that forms is
Ba(Zr,Y)O3 [57] while when Sn
4+ is added BaSnO3 forms [64–66]. The addition
of Ta5+ was reported to form both RE3TaO7 [67] and Ba2YTaO7 [68]. Consid-
ering the optimal results obtained with the Ta5+ in this chapter the effects of
39
4. Ba2YNbO6 doped YBa2Cu3O7−δ: preliminary results
another large, highly charged ion, Nb5+, are presented.
4.1.1 Nb doped ceramic samples
The first introduction of Nb in YBa2Cu3O7−δ was reported soon after the dis-
covery of the YBa2Cu3O7−δ. In 1988 Kuwabara M et al. reported an increased
Jc value in Nb2O5 doped YBa2Cu3O7−δ ceramic samples [108]. The effects of Nb
was unclear but the authors reported that almost all the Nb added was not in-
corporated in the YBa2Cu3O7−δ and was instead segregated to form a secondary
phase. An elemental distribution map showing that the Cu element did not co-
exist with the Nb while both Y and Ba did and the discovery that Nb was no
longer in the from of oxide made the authors claim the formation of a secondary
compound with Y and Ba.
In 1989 Greaves C. and Slater P.R. published a study on the Nb and Ta
substitution in the REBa2Cu3O7 (RE = Y, Eu, La) [109]. They suggest the
presence of two possible products, the REBa2Cu2MO8 and the Ba2REMO6 +
CuO (M = Nb, Ta). They observed that the energy difference between the two
possible products is small and that, since in the double perovskite structure the
rare earth cation is in the small octahedral site, the stability of the perovskite
structure is reduced when large RE3+ ions are used. The authors observed that
for RE = Y, Eu the phases formed where REBa2Cu3O7, Ba2REMO6 and CuO
while a substitution and the consequent formation of the REBa2Cu2MO8 phase
was only observed for RE = La. From their studies it is evident that doping
YBa2Cu3O7−δ (RE = Y) with Nb or Ta leads to the formation of the Ba2YNbO6
and Ba2YTaO6.
A few years later, in 1993, the diffraction pattern of a sample with composition
YBa2Cu2.95Nb0.005O7−δ was reported in a paper from Strukova G. K. et al.
[110,111]. They also found the presence of an YBa2Cu3O7−δ orthorhombic phase
(a = 0.3827 nm, b = 0.3894 nm, c = 1.1718 nm), a Ba2YNbO6 cubic perovskite
phase (a = 0.4218 nm) and a small amount of CuO. They also reported that the
intensity of the lines corresponding to the Ba2YNbO6 increased with increasing
the Nb content. They attributed a greater stability of the Ba2YNbO6 secondary
phase when compared to YBa2Cu3−xNbxO7−δ to the low solubility of the Nb in
40
4. Ba2YNbO6 doped YBa2Cu3O7−δ: preliminary results
the YBa2Cu3O7−δ.
Recently, in 2003, Babu N. H. et al. reported the presence of
Y2Ba4CuxNb1−xOy, a phase similar to Ba2YNbO6 with Cu partially substituting
Nb, in YBa2Cu3O7−δ bulk doped with Nb [112]. Improved Jc was observed in
the doped samples [113]. In 2009 Yeoh W. K. et al. reported enhanced flux pin-
ning at high field in YBa2Cu3O7−δ single grain bulk samples doped with NbO2
and related the improved pinning to the formation of a Nb rich phase within the
YBa2Cu3O7−δ [114].
4.1.2 Ba2YNbO6 use in YBa2Cu3O7−δ thin film
The stability and compatibility of the Nb-based double perovskite Ba2YNbO6
with the YBa2Cu3O7−δ was studied by Paulose K V et al. in 1992 [4]. They
found that annealing a 1:1 mole mixture of YBa2Cu3O7−δ and Ba2YNbO6 for 16
hr at 950 ◦C there was absolutely no reaction (figure 4.1).
The stability and compatibility of the Ba2YNbO6 with the YBa2Cu3O7−δ
made it a candidate for a novel substrate in the YBa2Cu3O7−δ thin film depo-
sition. The difficulties encountered in the production of large Ba2YNbO6 single
crystal have prevented the use this material as a substrate, nevertheless high
quality high temperature superconductor thin films and electronic devices were
fabricated adopting Ba2YNbO6 as a new buffer layer [115,116].
The only early report on the use of Ba2YNbO6 to enhance flux pinning in
YBa2Cu3O7−δ thin film is a work published in 1995 from Jia J. H. et al. [117].
The authors report on thin Nb added YBa2Cu3O7−δ deposited on ZrO2 substrates
by a dc magnetron sputtering method. The secondary phase formed in the film
is reported to be the Ba2YNbO6 perovskite but the effects on the critical current
registered were negative. The conclusions of this early research were that the
Ba2YNbO6 secondary phase does not play a role in the flux pinning process.
More then 10 years later, in 2007, Nb doped ErBa2Cu3O7−δ was successfully
produced by pulsed laser deposition [118–122], in this work the authors report
that c-axis oriented nanorods of Ba2ErNbO6 were formed in the ErBa2Cu3O7−δ.
However the effects of the nanorods on the critical current values were not clearly
explained. The presence on these c-axis oriented nanorods seemed to have effects
41
4. Ba2YNbO6 doped YBa2Cu3O7−δ: preliminary results
Figure 4.1: Powder diffraction patter for a) pure Ba2YNbO6, b) pure
YBa2Cu3O7−δ, c) 1:1 mole mixture of YBa2Cu3O7−δ:Ba2YNbO6 heated at 950
◦C for 15 hr. [4]
42
4. Ba2YNbO6 doped YBa2Cu3O7−δ: preliminary results
on the pinning performances of the films only when the latter were grown under
certain condition and a window of processability was undefined.
Recently, successful formation of Ba2YNbO6 nanorods in YBa2Cu3O7−δ thin
films grown by pulsed laser deposition has been reported by several authors [5,
74,123,124]. The article [5] is part of this thesis, it is the first attempt to produce
Ba2YNbO6 - YBa2Cu3O7−δ thin films by pulsed laser ablation of a mixed target,
and contains the first results obtained. These results will be described in the
following section of this chapter. In conclusion Ba2YNbO6 is a good pinning
phase candidate for the following reasons:
• The large Nb+5 does not substitute the Cu in YBa2Cu3O7−δ
• Ba2YNbO6 is the most stable Nb compound in a YBa2Cu3O7−δ matrix
• Ba2YNbO6 forms in a wide processing window
• Ba2YNbO6 nanoparticles forms self-assemble in nanorods in a similar man-
ner to BaZrO3
• The large lattice mismatch between Ba2YNbO6 and YBa2Cu3O7−δ should
provide additional pinning inducing lattice distortion in YBa2Cu3O7−δ
4.2 Ba2YNbO6 perovskite additions to YBa2Cu3O7−δ:
the preliminary results
In this section is reported the first results obtained by the addition of the Ba2YNbO6
perovskite pinning phase to the YBa2Cu3O7−δ. The results were discussed in a
feasibility study published in early 2010 in Superconductor Science and Technol-
ogy [5].
4.2.1 The Ba2YNbO6 synthesis, the pulsed laser deposi-
tion target preparation and thin films deposition
The technique adopted to produce the single phase Ba2YNbO6 powder is the
same described and successfully used by Paulose K V et al. [4]. Ba2YNbO6
43
4. Ba2YNbO6 doped YBa2Cu3O7−δ: preliminary results
powder were produced mixing and grinding stoichiometric quantity of 99.99%
Y2O3, Ba(NO3)2, Nb2O5 and a minimal amount of CuO2 followed by a solid
state reaction at 1450 ◦C for 24 h in flowing O2. In order to achieve a complete
reaction and thus produce a single phase Ba2YNbO6 two successive grinding and
reaction were needed. The minimal amount of CuO2 (≈ 0.5%wt.) was added to the
oxides mixture because it is described to greatly enhance the powder reactivity.
The reaction enhancing effect of the CuO2 is explained by the fact that CuO2
presence introduce in the system a low energy reaction intermediate lowering the
energy barrier of the overall reaction [4].
Figure 4.2: Powder diffraction patter for Ba2YNbO6powder produced. Adapted
from [5].
In the figure 4.2 are shown the results of an x-ray analysis in the Bragg-
Brentano geometry (θ − 2θ scan) of the Ba2YNbO6 powder produced by the
described solid state reaction. No impurities or secondary phase peaks are found,
given the minimal amount of CuO2 that was added is not expected to be easily
evidenced. The quantity involved is too small to be separated from the noise levels
that are usually present in this typologies of measures. The lattice parameters
calculated form the x-ray data is a = 0.844 nm ± 0.001 nm. The lattice parameter
derived is in agreement with the data recorded in the Joint Committee on Powder
44
4. Ba2YNbO6 doped YBa2Cu3O7−δ: preliminary results
Diffraction Standards database (JCPDS) as well as the data presented in more
recent studies [125].
The Ba2YNbO6 production process described constitutes a reliable source
of pure single phase Ba2YNbO6 to use, together the YBa2Cu3O7−δ commercial
powder, in the pulsed laser deposition target preparation.
The target sintering procedure adopted for the samples discussed in this
section is exactly the general procedure described in the previous chapter, the
Ba2YNbO6 doping amount in the target adopted in this study is 5%mol. The
control pure YBa2Cu3O7−δ target is sintered adopting the same technique. More
sophisticated procedures were also adopted to produce advanced targets, the
modification to the processes will be described in the appropriate spaces.
The pulsed laser deposition parameters adopted for both pure YBa2Cu3O7−δ
control thin films and the Ba2YNbO6 added samples are summarized in table 4.1.
Parameter Value
Substrate Temperature 770 ◦C
Chamber Pressure 0.3 mbar flowing O2
Laser Fluence 2 Jcm−2
Repetition Rate 5 Hz
Number of Pulses 4500
Annealing Time 1 hr
Annealing Temperature 520 ◦C
Annealing Pressure 500 mbar O2
Table 4.1: Pulsed laser deposition parameters
All the thin films deposited were measured to be of ≈ 0.5 µm thickness.
4.2.2 Crystalline structure analysis: x-ray diffraction data
The first analysis realized on deposited films when using a new target composi-
tion are usually x-ray diffraction structural analysis to investigate the crystalline
phases produced in the samples, their orientation and possibly the strains and
lattice distortions arising form the large interfaces of nanostructured composite
materials.
45
4. Ba2YNbO6 doped YBa2Cu3O7−δ: preliminary results
4.2.2.1 Crystalline phases identification
Figure 4.3: X-ray diffraction data for undoped YBa2Cu3O7−δ and 5%mol
Ba2YNbO6 doped film deposited on SrTiO3. Adapted from [5].
The first relevant results reported in figure 4.3 is that the x-ray diffraction data
collected in the Bragg-Brentano geometry from the deposited thin films (both
doped and pure) shows only the diffraction peaks related to the (00l) planes from
the YBa2Cu3O7−δ, the Ba2YNbO6 and the SrTiO3. This is evidence that also
the doped thin films have the usual c-axis orientation of pure YBa2Cu3O7−δ thin
films deposited on (001) SrTiO3 single crystal.
The only difference evidenced by the comparison of the diffraction data gath-
ered from the pure YBa2Cu3O7−δ control thin films and the Ba2YNbO6 doped
ones is the presence of the peaks related to the (00l) planes of the Ba2YNbO6. It
is possible to conclude that the stability of the Ba2YNbO6 - YBa2Cu3O7−δ sys-
tem during the pulsed laser deposition does not differ from the proved stability at
high temperatures [4], thus the only secondary crystalline phase introduced in the
thin films deposited is the Ba2YNbO6. Furthermore a first hint of the Ba2YNbO6
orientation is revealed by the presence of only the (00l) planes diffraction peaks,
the non superconducting phase is aligned out-of-plane with the YBa2Cu3O7−δ
46
4. Ba2YNbO6 doped YBa2Cu3O7−δ: preliminary results
matrix.
To conclude the crystalline phases identification a small discussion has to be
given about a small peak at 2θ ≈ 34◦. This peak is present in both the pure
YBa2Cu3O7−δ and the Ba2YNbO6 doped thin films thus its presence can not be
associated with the Ba2YNbO6 doping. The intensity is lower than that from
the (004) and (008) Ba2YNbO6 peaks, and the Ba2YNbO6 doping level is only
5%mol. A peak at 2θ ≈ 34◦ in the YBa2Cu3O7−δ system could be associated
with the (111) YBa2Cu3O7−δ peaks or with the (004) Y2O3. The first could be
due to a minimal fraction of misaligned grains, a small fraction a grown grains
can explain the peaks presence and it is commonly found at lower deposition
temperature that the one adopted in these films. The second could be formed
starting from a minimal Y imbalance in the target and considering that the peaks
is also present in the pure YBa2Cu3O7−δ control sample, this minimal imbalance
should be in the commercial pure YBa2Cu3O7−δ powder adopted in the target
sintering process. Nevertheless, since the intensity is very small and the peak is
not related to the Ba2YNbO6 doping, its presence does not influence the results.
4.2.2.2 Crystalline phases orientation
The assessment of the in-plane texture is realized by φ scans. The technique is
described in the previous chapter and the results obtained analysing a 5%mol
Ba2YNbO6 doped film are summarized in figure 4.4.
The (202) Ba2YNbO6 peak (χ = 45
◦, 2θ = 30.03◦) reveal an in-plane align-
ment since it matches both the (101) SrTiO3 (χ = 45
◦, 2θ = 32.46◦) and the
(202) YBa2Cu3O7−δ (χ = 57.06◦, 2θ = 27.62◦) peaks. This direct observation
of a cube on cube growth of the Ba2YNbO6 with the YBa2Cu3O7−δ combined
with the out-of-plane c-axis homogeneous orientation demonstrated in figure 4.3
suggest full heteroepitaxy between the YBa2Cu3O7−δ, the Ba2YNbO6 and the
SrTiO3.
A last point should be made on the diffraction peaks choice for the φ scans.
Some authors perform the analysis on the (103) YBa2Cu3O7−δ (χ = 45◦, 2θ =
32.8◦) peak and not the (102) YBa2Cu3O7−δ used in this work. Unfortunately
the first is extremely close to (101) SrTiO3 and in most analysers the two peaks
47
4. Ba2YNbO6 doped YBa2Cu3O7−δ: preliminary results
Figure 4.4: X-ray diffraction data from φ scans of (101) SrTiO3, (102)
YBa2Cu3O7−δ, (202) Ba2YNbO6from a 5%mol Ba2YNbO6 doped film. Adapted
from [5].
overlap. Since the intensity of the (101) SrTiO3 peak in much higher that the
intensity of the (103) YBa2Cu3O7−δ peak choosing the latter as reference peak
for the YBa2Cu3O7−δ in-plane orientation assessment could lead to misinterpre-
tation.
4.2.2.3 Crystallographic matching and strain
Once the phase orientations are known it is possible to propose a crystallo-
graphic matching model and to evaluate the lattice mismatch of Ba2YNbO6 with
YBa2Cu3O7−δ.
In figure 4.5a is presented the matching of YBa2Cu3O7−δ with Ba2YNbO6
viewed along the b-axis ([010]Y BCO direction). In this view 3 unit cells of the cubic
Ba2YNbO6 match 2 unit cells of YBa2Cu3O7−δ. The c-axis mismatch (related to
YBa2Cu3O7−δ) can be calculated with the following equation:
3aBY NO − 2cY BCO
2cY BCO
= +8.34% (4.1)
48
4. Ba2YNbO6 doped YBa2Cu3O7−δ: preliminary results
Figure 4.5: Crystallographic matching of YBa2Cu3O7−δ with Ba2YNbO6. a) b-
axis view and b) c-axis view [5].
49
4. Ba2YNbO6 doped YBa2Cu3O7−δ: preliminary results
In order to have a perfect lattice match along the c-axis the Ba2YNbO6 lattice
should be compressed along the c-axis. In the doped film the (00l) peaks from the
Ba2YNbO6 are indeed shifted to higher angles compared to the bulk values (figure
4.3). The resulting lattice parameter is of 0.830 nm ± 0.001 nm instead of the
0.844 nm ± 0.001 nm resulting from the powder scan (figure 4.2). This confirms
the presence of a compressive strain of ≈ 1.7%, this strain is much lower than the
theoretically derived mismatch. In addition the (00l) YBa2Cu3O7−δ diffraction
peaks are slightly shifted towards lower 2θ values (in the opposite direction of the
Ba2YNbO6 diffraction peaks) thus the YBa2Cu3O7−δ c-axis lattice parameter
seems to be extended to higher values. A tensile strain of ≈ 0.4% can be derived
from the lattice value of 1.173 nm calculated from the diffraction data presented in
figure 4.3. Nevertheless the overall strain given by the Ba2YNbO6 contraction and
the YBa2Cu3O7−δ extension is lower than the theoretical value thus an additional
strain relief mechanism is most likely to be playing a matching role. The formation
of misfit dislocations is one the possible strain relief mechanism and such defects
formation has been previously reported in BaZrO3 doped YBa2Cu3O7−δ [126]. A
semicoherent interface could also be responsible of the reduced strain levels in
the crystalline structures.
In figure 4.5b the matching of YBa2Cu3O7−δ with Ba2YNbO6 looking along
the c-axis ([001]Y BCO direction) is presented. The matching between the
YBa2Cu3O7−δ and the Ba2YNbO6 crystalline lattices is realised with 1 unit cell
of Ba2YNbO6 and 2 unit cell of YBa2Cu3O7−δ. The averaged in-plane lattice
mismatch (averaged along the a-axis and b-axis) is given by the following equa-
tion:
aBY NO − 2aY BCO
4aY BCO
+
aBY NO − 2bY BCO
4bY BCO
= +9.42% (4.2)
This large lattice mismatch, larger than the others studied non superconduct-
ing phase additions, should be beneficial for low field pinning thanks to possible
localized strain fields at the Ba2YNbO6-YBa2Cu3O7−δ interface [127], but it likely
to produce self-assembled columnar arrays of defects that are shorter and wider
than those produced with pinning addition with lower lattice mismatch [64], and
similar to those produced with BaZrO3 addition [57].
50
4. Ba2YNbO6 doped YBa2Cu3O7−δ: preliminary results
4.2.3 Nanostructure analysis: Atomic Force Microscopy
and Transmission Electron Microscopy
X-ray diffraction analysis is a powerful tool to investigate the crystalline struc-
tures and their orientation but they reveals little information on the nanostruc-
turing of these phases. To investigate the nanostructure nanoscale microscopy has
to be performed. In this section the surface topography is discussed by analysing
atomic force microscopy images while the morphological distribution of the sec-
ondary pinning phase and the distortion in the YBa2Cu3O7−δ lattice are studied
with cross-section transmission electron microscopy.
4.2.3.1 Surface topography (Atomic Force Microscopy)
In figure 4.6 are reported atomic force micrographs of a pure YBa2Cu3O7−δ and
a 5%mol Ba2YNbO6 doped film surface.
An analysis of the root mean square roughness performed on the surface to-
pography of the 5 µm × 5 µm scans gives information on the quality of the two
films in examination. The root main square roughness of the pure YBa2Cu3O7−δ
thin film is ≈ 9.1 nm (figure 4.6a), a slightly lower value of ≈ 7.8 nm was calcu-
lated for the 5%mol Ba2YNbO6 doped thin film (figure 4.6c).
Carefully analysing both the 5 µm × 5 µm images it is evident that a small
amount of particulate is present on the surfaces of both thin films, these superficial
particles of diameter ≥ 250 nm were deposited on the growing film during the
deposition and are a well known and studied in consequence of a pulsed laser
deposition process in which the substrate is perpendicular to the plume and the
target density is not optimal [105, 128]. Improved techniques could be adopted
to reduce and eliminate this phenomenon [128–132]; in applications that require
a perfectly clean surface adopting these methods is recommended. In this case
the presence of the particulate is not a major issue and as shown later is not
evidently reducing the overall superconducting properties. Furthermore since the
particulate is present on the surfaces of both doped and undoped thin films it
should not be related to the Ba2YNbO6 doping. Nevertheless in the next section
higher quality films obtained from higher density target are discussed.
The grain growth morphology is affected by the Ba2YNbO6 presence in the
51
4. Ba2YNbO6 doped YBa2Cu3O7−δ: preliminary results
Figure 4.6: Atomic Force Microscopy surface image of a pure YBa2Cu3O7−δ and
a 5%mol Ba2YNbO6 doped YBa2Cu3O7−δ thin films. a) 5 µm × 5 µm surface
area of pure YBa2Cu3O7−δ; b) 1 µm × 1 µm surface area of pure YBa2Cu3O7−δ;
c) 5 µm × 5 µm surface area of 5%mol Ba2YNbO6 doped YBa2Cu3O7−δ; d) 1
µm × 1 µm surface area of 5%mol Ba2YNbO6 doped YBa2Cu3O7−δ
52
4. Ba2YNbO6 doped YBa2Cu3O7−δ: preliminary results
doped YBa2Cu3O7−δ thin film. The first difference is the fact that in the Ba2YNbO6
sample the diameter of the growth grains appear to be slightly larger, the mean
diameter measured in the undoped YBa2Cu3O7−δ film in ≈ 500 nm while the
one measured in the Ba2YNbO6 doped YBa2Cu3O7−δ is ≈ 1 µm. This can be
explained by a lower nucleation rate in the Ba2YNbO6 doped film. The second
difference is the grain boundary shape, rough in Ba2YNbO6 doped YBa2Cu3O7−δ
films and smooth in undoped YBa2Cu3O7−δ films.
The main difference between the surface morphology of the doped film and
that of the pure YBa2Cu3O7−δ is evidenced by the comparison of the 1µm ×1µm
scans, figure 4.6b and 4.6d. The latter shows superficial particles of ≈ 15-20 nm
in diameter that are absent in the image of taken from the pure YBa2Cu3O7−δ
thin film. These superficial particles, that are to be attributed to the Ba2YNbO6
presence, are uniformly distributed over the sample surface and are not segre-
gated at the grain boundary. This is a first experimental evidence that a uni-
formly distributed array of defects based on Ba2YNbO6 is possibile given that the
Ba2YNbO6 is able to grow inside a YBa2Cu3O7−δ crystal in the form of nanomet-
ric inclusion and is not expelled from the growing YBa2Cu3O7−δ grains during
the deposition (figure 4.6d).
4.2.3.2 Cross-section transmission electron microscopy
The morphology of the Ba2YNbO6 inclusions was studied with the use of cross
sectional imaging of the sample by transmission electron microscopy. In figure 4.7
is shown cross-section transmission micrography of a 5%mol Ba2YNbO6 doped
thin film. Nanorods aligned with the c-axis of the films are evidently shown.
The nanorods imaged are ≈ 10 nm in diameter and the spacing between two
adjoining rods is ≈ 40 nm. A similar spacing is observed in the superficial par-
ticles shown by the atomic force microscopy scan reported in figure 4.6d. These
similar spacing values is a validation of the assumption that the particles observed
on the surface by using the atomic force microscopy are related to the nanorods
terminating at the film surface. The diameter of the particles observed on the
surface with the atomic force microscopy appears to be larger then the diameter
of the nanorods observed with the transmission electron microscopy. There are
53
4. Ba2YNbO6 doped YBa2Cu3O7−δ: preliminary results
Figure 4.7: Transmission electron microscope cross-sectional image of a 5%mol
Ba2YNbO6 doped YBa2Cu3O7−δ thin film. White arrows in the direction of the
c-axis mark the position of self assembled Ba2YNbO6 nanorods parallel to the
c-axis. Larger white arrows parallel to the YBa2Cu3O7−δ ab-planes mark the
position of nanoparticles and of the interface between the substrate and the thin
film. Inset shows selected area electron diffraction pattern taken from a region
in proximity of a Ba2YNbO6 inclusion showing diffraction spot of Ba2YNbO6
YBa2Cu3O7−δ and SrTiO3. (TEM images from Prof. H Wang research group at
Texas A&M University).
54
4. Ba2YNbO6 doped YBa2Cu3O7−δ: preliminary results
two plausible explanations that are not mutually exclusive but may be jointly
responsible for the observed difference. The first explanation is found in the fact
that the superficial particles might actually have a larger diameter because at the
surface the ions mobility is higher and there may also be further agglomeration
immediately after the deposition is terminated. The second possible explanation
is related to differences in the resolution of the two methods of analysis used. In
particular, the resolution of the atomic force microscopy depends on the diam-
eter of the tip adopted. Tips with a larger diameter can make features appear
larger than they are in reality. This aberration is absent in the transmission elec-
tron microscopy where the captured images are not subjected to the described
deformation.
Additional information that can be derived from the nanorods spacing is the
matching field. The matching field is the field value at which the spatial density
of defects is equal to that of the flux lines passing through the superconductor. At
this field value the pinning potential is the highest since in theory each and every
flux line should be pinned by a single nanorod. Unfortunately, unlike a perfect
Abrikosov lattice, it is not possible to identify a triangular or square arrangement
of flux lines in superconductors with defects. In fact, the flux lines tend to change
their spatial arrangement to overlap the defects distribution. Nevertheless it is
at least possible to have an estimation of the matching field by using the mean
spacing of the nanorods noticing that for a random distribution this tends to
be higher than that derived for a square distribution and lower that that of a
triangular distribution. The equations reported below refer to the matching field
calculation derived for a square and a triangular distribution (note that Φ0 is the
quantized flux ≈ 2.067× 10−15 Wb and d is the defects average spacing)
H∗ =
Φ0
d2
(4.3)
H∗ =
2
√
3Φ0
3d2
(4.4)
Calculating the matching field for a defects average spacing of ≈ 40 nm
applying the formula derived for a square distribution 4.3 gives a matching field
H∗ ≈ 1.29 T. Applying the formula derived for triangular distribution 4.4 the
55
4. Ba2YNbO6 doped YBa2Cu3O7−δ: preliminary results
calculated matching field H∗ is ≈ 1.49 T. The applied magnetic field value in
which the pinning potential of the analysed thin films should be higher is between
1.29 T and 1.49 T.
In order to validate and complete the x-ray diffraction data reported in the
previous section an electron diffraction pattern of the cross section is shown in
the inset to figure 4.7. It is possible to observe a set of streaky diffraction dots for
the (004) Ba2YNbO6. The presence of streaky dots is related to large Ba2YNbO6
c-axis distortions. These distortions were expected due to the fact that the x-
ray analysis reported a compressive strain along the c-axis of the Ba2YNbO6 in
the order of ≈ 1.7%. The lattice spacing d (004) calculated from the electron
diffraction pattern is ≈ 0.214 nm, indicating a lattice parameter a ≈ 0.856
nm. This lattice parameter value, while being broadly consistent with the x-
ray measurements and confirming the secondary phase as the Ba2YNbO6 double
perovskite, does not confirm the compressive nature of the strain.
In the cross sectional images the nanorod inclusions are not the only nanoin-
clusions evidenced. A small fraction of nanoparticles are evidenced by the white
arrows parallel to the ab-planes. The presence of both self assembled nanorods
and nanoparticles is a feature which was already observed in BaZrO3 added
YBa2Cu3O7−δ as well as RETa3O7 added YBa2Cu3O7−δ, the amount of nanopar-
ticles is related to the deposition condition [133,134]. In the work of Boris Maiorov
et al., published in 2009, an interesting discussion of the beneficial effects result-
ing from simultaneous presence of both nanorods and nanoparticles is reported.
In their work it is shows how the synergetic combination of such 1D defects
(nanorods) and point defects (nanoparticles) can strongly enhance the flux pin-
ning by preventing some low energy depinning mechanism to take place. In
particular it is discussed how the flux jump from two adjacent nanorods with a
double kink propagation mechanism is associated with a higher energetic barrier
when a nanoparticle is introduced between the two adjacent nanorods or when
the nanorods are not highly linear, continuous and are not all perfectly parallel
to the c-axis.
Analysing the higher magnification images (figure 4.8a and 4.8b) it is clear how
the nanorods formed by the Ba2YNbO6 are wider (≈ 10 nm) than the RETa3O7
(≈ 5 nm) and BaZrO3 (≈ 7-8 nm), with a larger angle spread around the c-
56
4. Ba2YNbO6 doped YBa2Cu3O7−δ: preliminary results
Figure 4.8: Transmission electron microscope cross-sectional image of a 5%mol
Ba2YNbO6 doped YBa2Cu3O7−δ thin film. White arrows in the direction of the
YBa2Cu3O7−δ c-axis mark the position of self assembled Ba2YNbO6 nanorods
parallel to the c-axis. Arrows parallel to the YBa2Cu3O7−δ ab-planes mark
the position of nanoparticles. The Ba2YNbO6 nanorods presence is indicated
by the Moire fringes arising from lattice mismatch between Ba2YNbO6 and
YBa2Cu3O7−δ. (TEM images from Prof. H Wang research group at Texas A&M
University).
57
4. Ba2YNbO6 doped YBa2Cu3O7−δ: preliminary results
axis and shorter ≈ 100 nm as opposed to whole film thickness. Furthermore
adding the simultaneous presence of nanoparticles and nanorods the resulting
pinning landscape should be optimal in minimizing the low energy depinning
mechanism efficiency. On the downside the fact that the nanorods tend to be
wider indicates a lower rate of nucleation than RETa3O7, BaZrO3 and BaSnO3.
The main consequence is a lower defects density at equal volumetric doping that
should result in a lower pinning efficiency at high field values.
One last thing to notice in this transmission electron microscopy study of the
Ba2YNbO6 doped YBa2Cu3O7−δ sample is the evidence of how the YBa2Cu3O7−δ
lattice is deformed by the nanorod inclusions. As perfectly pictured in figure
4.8b the YBa2Cu3O7−δ ab-planes are curved around the nanorods. The lattice
mismatch between the YBa2Cu3O7−δ and the Ba2YNbO6 induce buckling of the
YBa2Cu3O7−δ crystalline planes around the nanorods. Superconductivity is then
expected to be depressed in the region surrounding the Ba2YNbO6 nanorods.
The high possibility of additional pinning from strain fields in the vicinity of
Ba2YNbO6-YBa2Cu3O7−δ interfacial regions is once more evidenced.
4.2.4 The superconducting properties: Tc, Jc(B) and Jc(B, θ)
A preliminary study of the effects of Ba2YNbO6 inclusions in YBa2Cu3O7−δ can
not be considered complete without the analysis of the main superconducting
properties. Once the crystalline structure and the nanostructure of the sam-
ple have been discussed it is necessary, in order to have an overall picture of
Ba2YNbO6 performance, to evaluate the effective superconducting properties.
Considering a power application such as superconducting cable or a more sophis-
ticated application such as superconducting coils for magnets, the most important
properties to evaluate are the transition temperature Tc, the variation of the crit-
ical current density as a function of the applied magnetic field value Jc(B) and
the variation of the critical current density as a function of the direction at which
the magnetic field is applied Jc(B, θ).
58
4. Ba2YNbO6 doped YBa2Cu3O7−δ: preliminary results
Figure 4.9: Resistance variation with the temperature measured on a 5%mol
Ba2YNbO6 doped YBa2Cu3O7−δ thin film. The inset shows the region of the
superconducting transition magnified.
4.2.4.1 Transition temperature, Tc
In figure 4.9 is reported the resistance measured on a 50 µm width current track
patterned on 5%mol Ba2YNbO6 doped YBa2Cu3O7−δ thin film. The transition
temperature Tc measured as the temperature at which the resistance disappears
is Tc = 89 K. The reduction in the Tc is lower than expected. Reduced Tcs have
been reported for most of the doped YBa2Cu3O7−δ films produced [57,68,118,119].
The Tc reduction is usually a function of the volumetric amount of the doping and
the lattice distortion induced. It is important to notice that similar Tc reductions
were observed in 5%mol BaZrO3 doped YBa2Cu3O7−δ thin films [67], but the
volumetric amount of Ba2YNbO6 in 5%mol Ba2YNbO6 doped YBa2Cu3O7−δ thin
films is higher than that of BaZrO3 in a 5%mol BaZrO3 doped YBa2Cu3O7−δ thin
film. In particular a 5%mol Ba2YNbO6 doped YBa2Cu3O7−δ thin film contains a
4.2%vol of Ba2YNbO6 while a BaZrO3 in a 5%mol BaZrO3 doped YBa2Cu3O7−δ
thin film contains a 1.8%vol of BaZrO3 thus the Tc reduction registered in the
Ba2YNbO6 doped scenario is lower than explected.
59
4. Ba2YNbO6 doped YBa2Cu3O7−δ: preliminary results
Two main informations can be derived from the observation of this small Tc
reduction. First, the YBa2Cu3O7−δ planes buckling around the nanorods (pic-
tured in the transmission electron microscopy figure 4.8b) is an effective strain
sink capable of limiting the strain field propagation in a region close to the
nanorods, for this reason the detrimental effects of the strain on the transition
temperature are minimal. Second, as it was expected from earlier report on Nb
doped YBa2Cu3O7−δ ceramics, the Nb do not substitute in the YBa2Cu3O7−δ
and the Ba2YNbO6 do not react with the superconducting lattice. In other
words the Nb remains enclosed in the stable Ba2YNbO6 and the only effects on
the YBa2Cu3O7−δ crystalline structure are related to the lattice accommodation
of the mismatch.
4.2.4.2 The critical current density
Despite the slightly reduced Tc the critical current density Jc is improved over
the whole field values window analysed.
In figure 4.10 is reported Jc(B) measurements up to 1 T with the field applied
parallel to the c-axis (B‖c) at 77 K for a pure YBa2Cu3O7−δ thin film and a
5%mol Ba2YNbO6 doped YBa2Cu3O7−δ thin film. It is evident that the Jc is
higher for the doped thin film and that at 1 T the Jc is a higher by a factor of
≈ 2.
In the Jc(B) measurements there is a region for low values of the applied
magnetic field in which the critical current Jc can be approximated with the
following equation:
Jc(B) ≈ Jc(0)−α∗B (4.5)
The lower the α value the lower is the Jc reduction with increasing the applied
magnetic field value. In the log-log plot shown in the inset are reported the calcu-
tad α values for both the films. The lower α value calculated for the Ba2YNbO6
doped YBa2Cu3O7−δ thin film is an indication of a higher pinning potential. The
calculated α value for the Ba2YNbO6 doped YBa2Cu3O7−δ sample is α = 0.37
that is lower than the α = 0.51 calculated for YBa2Cu3O7−δ films of this study.
The measurements shown in figure 4.10 are exhaustive in describing the Ba2YNbO6
60
4. Ba2YNbO6 doped YBa2Cu3O7−δ: preliminary results
effects when the field is applied parallel to the c-axis, but in order to have a more
complete picture of the Ba2YNbO6 efficiency as a pinning additive it is necessary
to investigate the angular dependence of Jc. The angular dependence of the criti-
cal current density at 77 K for an applied field value of 0.5 T is reported in figure
4.11.
Figure 4.10: Critical current density variation with the applied magnetic field
value measured on a pure YBa2Cu3O7−δ and a 5%mol Ba2YNbO6 doped
YBa2Cu3O7−δ thin films at 77 K and field applied parallel to the c-axis of the
YBa2Cu3O7−δ. The inset shows data on a logaritmic axis and the calculated α
values.
The main characteristic of the Jc(B, θ) measured for the 5%mol Ba2YNbO6
doped YBa2Cu3O7−δ thin film is a strongly reduced anisotropy compared to
the Jc(B, θ) measured for the undoped YBa2Cu3O7−δ thin film. The reduced
anisotropy is related to an increase of the pinning potential when field is applied
parallel to the Ba2YNbO6 nanorods. Since the nanorods are growing parallel to
the c-axis of the YBa2Cu3O7−δ (figure 4.7) the pinning potential, and thus the
Jc, is increased when the field is applied parallel to the c-axis of the thin film,
61
4. Ba2YNbO6 doped YBa2Cu3O7−δ: preliminary results
Figure 4.11: Critical current density variation with the direction of the applied
magnetic field measured on a pure YBa2Cu3O7−δ and a 5%mol Ba2YNbO6 doped
YBa2Cu3O7−δ thin films at 77 K and magnetic flux density = 0.5 T
which is orthogonal to the sample surface. Furthermore the anisotropy is reduced
also by a reduction of the Jc when the field is applied parallel to the ab-plane of
the YBa2Cu3O7−δ.
The two changes in the pinning potential described can be related to the
presence of Ba2YNbO6. The increment of the critical current density, Jc, when
the field is applied parallel to the Ba2YNbO6 nanorods was expected and has been
described for all the secondary phases forming c-axis aligned nanorods. In general
when the field is applied parallel to the direction of any form of one dimensional
defects, such as Ba2YNbO6 nanorods, the interaction between the flux lines and
the defects is maximized and for this reason the pinning effectiveness of the defects
is optimal. At the same time the interruption of the intrinsic pinning layer of
the YBa2Cu3O7−δ (ab-planes) generated by the introduction of the Ba2YNbO6
nanorods together with the induced buckling of the YBa2Cu3O7−δ crystalline
planes around the nanorods could be responsible for the observed reduction of
the critical current density when the field is applied parallel to the ab-planes of
the YBa2Cu3O7−δ.
62
4. Ba2YNbO6 doped YBa2Cu3O7−δ: preliminary results
An interesting feature that can be observed when comparing the effects of
the Ba2YNbO6 introduction in the YBa2Cu3O7−δ to the angular dependence
of the critical current density with the effects of the other pinning secondary
phases (Ba(Zr,Y)O3, RETa3O7, BaSnO3 and Ba2YTaO7) is that the Ba2YNbO6
nanorods appears to be effective over a larger angular range producing a broader
c-axis peak. The formation of shorted nanorods not perfectly oriented parallel to
the c-axis of the YBa2Cu3O7−δ (discussed in the transmission electron microscopy
analysis) could be the reason behind this increased effectiveness over the angles
the magnetic field application.
To conclude the Jc(B, θ) analysis a small discussion on the effects of the
Ba2YNbO6 nanoparticles observed in transmission electron microscopy of the
thin film cross-section (figure 4.8a) has to be done. Nanoparticles are usually
related to an increase of the isotropic pinning [69]. A strong additional isotropic
pinning increases the critical current density, Jc, independently of the angle at
which the magnetic field is applied. Thus the presence of nanoparticles would
have no effect on the shape of the Jc(B, θ). On the other hand the reduced
pinning potential observed when the magnetic field is applied parallel to the ab-
planes of YBa2Cu3O7−δ would suggest that an addition of isotropic pinning is
absent. In reality the conclusion that can be made is that if nanoparticles are
effectively adding an isotropic pinning to the Ba2YNbO6 doped YBa2Cu3O7−δ
thin film this is lower than the reduction of the ab-pinning potential generated by
the Ba2YNbO6 introduction. Furthermore it is important to remember that the
synergetic combination of Ba2YNbO6 nanorods and nanoparticles can strongly
enhance the pinning provided by the Ba2YNbO6 nanorods to the fields applied
parallel to the c-axis of the YBa2Cu3O7−δ [134].
4.2.5 Concluding remarks to the preliminary results
In this section were reported the first results obtained by the addition of the
Ba2YNbO6 perovskite pinning phase to YBa2Cu3O7−δ. The double perovskite
Ba2YNbO6 results as a stable secondary phase that can be used with positive
results as a pinning additive. Ba2YNbO6 is capable of forming c-axis oriented
nanorods which result in shorter, wider nanorods than those generated by the
63
4. Ba2YNbO6 doped YBa2Cu3O7−δ: preliminary results
other known pinning additives; furthermore the Ba2YNbO6 nanorods are oriented
parallel to the c-axis of the YBa2Cu3O7−δ but with a larger splay angles range.
For these reason the Ba2YNbO6 nanorods appear to be effective over a relatively
large angular range of magnetic field application and not only when the field is
applied parallel to c-axis of the YBa2Cu3O7−δ.
64
Chapter 5
Ba2YNbO6 doped YBa2Cu3O7−δ:
the deposition parameters
The study in this chapter is an evaluation of how the production process param-
eters affect the formation of the secondary phase and the overall epitaxial quality
of the thin films. Furthermore it is discussed how the epitaxial quality of the
film and the morphology of the secondary phase influence the superconductive
properties.
5.1 Ba2YNbO6 perovskite additions to YBa2Cu3O7−δ:
the effects of the deposition rate on the nanos-
tructure and the superconducting proper-
ties
The results obtained from the preliminary study on Ba2YNbO6 doped
YBa2Cu3O7−δ thin films showed that the secondary phase Ba2YNbO6 can be
added into a mixed YBa2Cu3O7−δ-Ba2YNbO6 pulsed laser deposition target to
obtain pinning enhanced superconducting thin films. In order to investigate the
real potential of the double perovskite Ba2YNbO6 a study of the influence of
the two principal deposition parameters on the nanostructure and the supercon-
ducting properties of Ba2YNbO6 doped YBa2Cu3O7−δ thin films was performed.
65
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
A set of 5%mol Ba2YNbO6 doped YBa2Cu3O7−δ thin films was deposited by
varying the deposition rate. The crystalline structure, nanostructure and su-
perconducting properties of the thin films produced were analyzed. In this sec-
tion are reported the analysis results that allows a better understanding of the
Ba2YNbO6-YBa2Cu3O7−δ system overall and they demonstrate the possibility
to tune the nano structure and the pinning properties of the Ba2YNbO6 doped
YBa2Cu3O7−δ thin films by adjusting the deposition parameters.
5.1.1 Ba2YNbO6 doped YBa2Cu3O7−δ target preparation
and thin film deposition
The target used in the pulsed laser deposition of Ba2YNbO6 doped thin films
analysed in this section is similar to the one that was described in the previous
section. Since the preliminary study reported in the previous section evidenced
a small fraction of particulates on the thin film surfaces (see 4.2.3.1), a variation
to the production process of the target adopted for the analysis described in this
section was implemented. Densification of the pulsed laser deposition target is
a first countermeasure to the formation of particulate on the deposited thin film
surface. In order to achieve a more dense target, once sintered a first time it
was subjected to a subsequent regrinding. The powders obtained by the second
grinding were again pressed in the form of a cylindrical target and the sintered
at 950 ◦C in oxygen flow for additional 12hr.
The laser pulses repetition rate was systematically modified to study its ef-
fects on the nanostructure and pinning properties of the deposited thin films. The
laser pulse repetition rate variation changes the maximum time allowed to the
ions migration on the surface of a growing film before new higher energy ions are
deposited from a subsequent laser pulse. A higher rate increases the deposition
rate but lowers the “free migration” time. This parameter has a direct effect on
the migration time and thus a direct effect on the crystalline and nanostructure
of the growing films. In particular repetition rates of 1, and 10 Hz and substrate
temperatures of 780 ◦C were investigated. In the table below 5.1 are summa-
rized the pulsed laser deposition parameters adopted for the deposition of the
Ba2YNbO6 doped YBa2Cu3O7−δ thin films analysed in this section.
66
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
Parameter Value
Substrate Temperature 780 ◦C
Chamber Pressure 0.3 mbar flowing O2
Laser Fluence 2 Jcm−2
Repetition Rate 1, 10 Hz
Number of Pulses 4500
Annealing Time 1 hr
Annealing Temperature 520 ◦C
Annealing Pressure 500 mbar O2
Table 5.1: Pulsed laser deposition parameters
All the Ba2YNbO6 doped YBa2Cu3O7−δ thin films deposited were measured
to be of ≈ 0.4 µm thickness. Some additional thicker samples of ≈ 0.7 µm
thickness where also deposited using 8000 laser pulses.
5.1.2 Crystalline structure analysis: x-ray diffraction data
Similarly to the crystalline phases characterization realised for the first Ba2YNbO6
doped YBa2Cu3O7−δ system, reported in the previous section, the crystalline
structures of all the samples produced for the study of the effects of the depo-
sition parameters were analysed with an extensive use of the x-ray diffraction
technique.
5.1.2.1 Crystalline phases identification, phases orientation
The x-ray diffraction data collected in the Bragg-Brentano geometry to identify
the crystalline phases as well as the in-plane texture analysis by examination of
φ scans don’t add additional information to the ones gathered in the previous
section. The results obtained at the different growth condition closely matches
each other. For this reason it is hard to understand if there are differences in the
crystalline structures of the Ba2YNbO6 doped YBa2Cu3O7−δ thin films that were
deposited in the different conditions described.
At a first x-ray diffraction analysis all the samples appears to have similar
crystalline features. The only diffraction peaks evidenced by x-ray diffraction
analysis in the Bragg-Brentano geometry (θ − 2θ scan) are those related to the
67
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
(00l) planes from the YBa2Cu3O7−δ, the Ba2YNbO6 and the SrTiO3. Further-
more also the φ scans closely matches for the different thin films and confirms
the in-plane orientation discussed in the previous section.
In conclusion the θ− 2θ and the φ scans are not capable at identifying differ-
ences in the crystalline structures and varying laser pulses repetition rate do not
change the fact that the Ba2YNbO6 doped YBa2Cu3O7−δ thin films deposited
con be mainly described as epitaxial YBa2Cu3O7−δ films with the presence of a
cube on cube oriented Ba2YNbO6 perovskite within the YBa2Cu3O7−δ.
5.1.2.2 Epitaxy quality: rocking curve
Even if the standard x-ray diffraction analysis does not detect differences in the
crystalline structure, a first classification can be made on the basis of the epitaxial
quality. The quality of the YBa2Cu3O7−δ epitaxy with the SrTiO3 can be anal-
ysed by the rocking curve detection. The half maximum width of the diffraction
peaks resulting from a rocking curve scan is an indication of the angular spread of
the c-axis orientation. The lower the spread, the higher is the c-axis alignment.
Thin Film SrTiO3 YBa2Cu3O7−δ Ba2YNbO6
(002) (005) (004)
YBa2Cu3O7−δ (1 Hz) 0.10 0.18
YBa2Cu3O7−δ (10 Hz) 0.12 0.18
doped YBa2Cu3O7−δ (1 Hz) 0.11 0.15 1.0
doped YBa2Cu3O7−δ (10 Hz) 0.10 0.20 2.1
Table 5.2: Full Width Half Maximum values of the SrTiO3 (002), YBa2Cu3O7−δ
(005) and Ba2YNbO6 (004) rocking curves measured on pure YBa2Cu3O7−δ and
5%mol Ba2YNbO6 doped YBa2Cu3O7−δ thin films deposited at a substrate tem-
perature of 780 ◦C and repetition rate of 1 and 10 Hz.
In the table 5.2 is reported the full width half maximum (FWHM) val-
ues of the rocking curves around the (002) SrTiO3, (005) YBa2Cu3O7−δ and
(004) Ba2YNbO6 planes of the different samples. The FWHM values of the
YBa2Cu3O7−δ do not reveal an evident pattern and the most relevant informa-
tion is that the YBa2Cu3O7−δ c-axis orientation do not appear to be influenced by
the deposition rate in absence of the Ba2YNbO6 phase. The Ba2YNbO6 FWHM
values instead shows that the sample deposited at higher deposition rate (higher
68
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
laser repetition rate) shows a larger FWHM value than the one obtained for the
sample deposited at lower deposition rate. Larger “free migration” time allow the
formation of a Ba2YNbO6 crystalline lattice with a higher c-axis alignment.
One thing to notice is that since the Ba2YNbO6 is growing inside a
YBa2Cu3O7−δ lattice its orientation quality is strictly related to YBa2Cu3O7−δ.
More precisely the Ba2YNbO6 lattice does not present a better c-axis orientation
alignment than the YBa2Cu3O7−δ lattice in which it is formed, on the other hand
it is also true that a poorly oriented Ba2YNbO6 can be formed in a perfectly
oriented YBa2Cu3O7−δ matrix.
While it is possible to use a technique developed to investigate the c-axis
alignment quality of thin films to also analyse that of the Ba2YNbO6 nanopar-
ticles and nanorods it is not possible to directly compare the FWHM obtained
for the Ba2YNbO6 nanoinclusions with those obtained with the YBa2Cu3O7−δ
thin films. In general, since the Ba2YNbO6 is a dopant and its quantity is
only a fraction of the YBa2Cu3O7−δ, the intensity x-ray diffraction peaks re-
lated to the Ba2YNbO6 are ≈ 2 orders of magnitude lowers than those related
to the YBa2Cu3O7−δ. Furthermore the morphological nature of the Ba2YNbO6,
nanorods and nanoparticles, is responsible for an additional broadening of the
Ba2YNbO6 peaks in the diffraction pattern that is related to the size of the phase
and not the strain or orientation. For these reason the FWHM values obtained
from the Ba2YNbO6 should be taken as a comparison between the different c-axis
degrees of the Ba2YNbO6 phase only.
5.1.2.3 Epitaxy quality: reciprocal space maps
The rocking curve were used to investigate the quality of the c-axis alignment for
both the Ba2YNbO6 and YBa2Cu3O7−δ separately but while a first information
on the Ba2YNbO6 phase was revealed it was impossible to obtain additional
information on the YBa2Cu3O7−δ . In order to further investigate the effects of
the deposition parameters on the overall epitaxial quality of the YBa2Cu3O7−δ
and Ba2YNbO6 nanoparticles reciprocal space maps were gathered for a pure
YBa2Cu3O7−δ thin film, and two Ba2YNbO6 doped YBa2Cu3O7−δ thin films. In
particular the Ba2YNbO6 doped films were deposited adopting different repetition
69
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
rate of the ablating pulsed laser, 1 and 10 Hz, but maintaining the same substrate
temperature of 780 ◦C; the pure YBa2Cu3O7−δ was deposited at 1 Hz and the
same substrate temperature.
Figure 5.1: X-ray diffraction reciprocal space map measured on a pure
YBa2Cu3O7−δ thin film.
In figure 5.1 is reported the reciprocal space map gathered from a pure
YBa2Cu3O7−δ sample. There are three main diffraction peaks, the (103) of the
SrTiO3 and the (109) - (019) of the YBa2Cu3O7−δ are overlapping each other
while is clearly visible the (108) - (018) of the YBa2Cu3O7−δ. The YBa2Cu3O7−δ
peaks are actually formed by two different peaks related to the inevitable twins
formation. It is impossible to distinguish between the (108) and the (018)
YBa2Cu3O7−δ peaks because the two peaks are too close in the reciprocal space
that they appear as a single large peak, the theoretical peaks position indicated in
the reciprocal space maps are well between the peaks boundary and the center of
the peak is between this two theoretical position as expected in twin formations.
It is possible from the reciprocal space maps to calculate the lattice parameters
directly from the coordinate of a peak in the reciprocal space if the (h,k,l) index
is known. Remembering that the reciprocal space coordinates (qx and qz) are the
reciprocal of the lattice spacing (dx and dz) the lattice parameters a,b and c for
70
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
a (h,0,l) and a (0,k,l) can be calculated applying the following formulas:
for a (h,0,l) peak
a =
h
qx
(5.1)
c =
l
qz
(5.2)
for a (0,k,l) peak
b =
k
qx
(5.3)
c =
l
qz
(5.4)
In figure 5.1 it is easier to calculate the YBa2Cu3O7−δ lattice parameters from
the (108)-(018) peaks since these are not overlapping with the any other phases
peak. Its important to notice that similarly to the lattice mismatch calculated in
the last section since it not possible to separate the (108) peak from the (018) the
lattice parameter calculated from the qx coordinate of the reciprocal space map
refers to an average value between the a and b lattice parameters. The average
ab calculated from the a and b bulk values (a = 0.382 nm and b = 0.388 nm), is
abbulk = 0.385 nm and it is perfectly in line with the average value calculated from
the (108)-(018) YBa2Cu3O7−δ taken from the reciprocal space map in figure 5.1
where a qx = 2.60 nm
−1 also gives a ab = 0.385 nm. Similarly the YBa2Cu3O7−δ
c lattice parameter bulk value is c = 1.168 nm and from a qz = 6.85 nm
−1 a c =
1.168 nm is calculated. The YBa2Cu3O7−δ thin film while being epitaxial with
the SrTiO3 substrate it is relaxed and the lattice parameters have the same value
reported in literature for bulk YBa2Cu3O7−δ.
Focusing on the (103) SrTiO3 diffraction peak at qx = 2.56 nm
−1 and qz = 7.68
nm−1, it is evident that additional diffraction peaks of (103) are present. The
presence of these additional peaks is related to the use of a x-ray incident beam
that is not monochrome and, even if the main quantity of x-rays are the Cu Kα
with λ = 0.15405 nm, there are also secondary x-rays generated at different λ that
are the cause of the additional diffraction peaks presence. The intensity of the
secondary peaks is lower and these are only visible for the high intensity SrTiO3
71
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
diffraction peaks. Another feature related to the (103) SrTiO3 is a diagonal line
also only visible for the high intensity diffraction generated by the substrate, also
this line is an artifact generated by the absence of a monochromator in particular
the line is generated by non copper x-rays emissions.
Figure 5.2: X-ray diffraction reciprocal space map measured on 5%mol Ba2YNbO6
doped YBa2Cu3O7−δ thin films. a) Film deposition rate of 1 Hz; b) Film deposi-
tion rate of 10 Hz.
72
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
In figure 5.2 are reported the reciprocal space map gathered from a Ba2YNbO6
doped YBa2Cu3O7−δ sample deposited at 1 Hz repetition rate (figure 5.2a) and
a Ba2YNbO6 doped YBa2Cu3O7−δ thin film deposited at 10 Hz repetition rate
(figure 5.2b). In addition to the three main diffraction peaks also present in figure
5.1, that are related to the YBa2Cu3O7−δ and the substrate, the (103) Ba2YNbO6
peak is also present. Comparing the image in figure 5.1 with the images in figure
5.2 the (103) Ba2YNbO6 peak is not the only evident difference, a broadening of
the YBa2Cu3O7−δ related peaks is a second clear effect of the Ba2YNbO6 doping.
The broadening of the YBa2Cu3O7−δ peaks in the Ba2YNbO6 doped samples
could be related to the buckling of the crystalline YBa2Cu3O7−δ planes around
Ba2YNbO6 nanoinclusion discussed in the last section. The fact that overall the
lattice parameter of the YBa2Cu3O7−δ does not seem to change confirms that the
broadening is related to a local distortion, like planes buckling.
Finally a direct indication of the effects of the deposition parameters on the
crystalline structure can be obtained by the comparison of the reciprocal space
maps of the two different Ba2YNbO6 doped YBa2Cu3O7−δ thin films (figure 5.2a
and 5.2b). First the (103) Ba2YNbO6 peak shape and intensity appear to be
different in the two samples. In particular the Ba2YNbO6 peak of the 10 Hz
sample (figure 5.2b) has a lower intensity and a long tail is present. These two
features could be evidence of the existence of a fraction of randomly oriented
Ba2YNbO6 that is absent in the 1 Hz sample (figure 5.2a). From the (103)
Ba2YNbO6 peak of the 1 Hz sample it is possible to calculate the lattice param-
eters a of the Ba2YNbO6 phase in the planar direction and in the YBa2Cu3O7−δ
c-axis direction from the qx and qz values respectively. Applying the equations
5.1 and 5.2 on the qx = 2.32 nm
−1 and qz = 7.22 nm−1 taken from figure 5.2a the
lattice parameters a = 0.430 nm along the planar direction and c = 0.415 nm
are calculated. The Ba2YNbO6 phase lattice parameter bulk value is a = 0.422
nm, and comparing this value with those calculated from the reciprocal space
map it is clear that the Ba2YNbO6 is subjected to a compressive strain along the
YBa2Cu3O7−δ c-axis direction and an elongation strain along the YBa2Cu3O7−δ
ab-planes direction. The deformation calculated for the two different direction
are + 1.9 % along the ab-planes direction and - 1.66 % along the c-axis direction.
Considering the morphological nature of the Ba2YNbO6 inclusions it is possible
73
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
to explain why the compressive strain is only acting in the c-axis direction even if
the lattice mismatch calculated from the bulk values (in-plane lattice mismatch =
+ 9.42 % and c-axis lattice mismatch = + 8.34 %) indicate a compressive strain
in all directions. The Ba2YNbO6 is mainly forming nanorods aligned parallel to
the c-axis of the YBa2Cu3O7−δ thus the interface between the Ba2YNbO6 and
the YBa2Cu3O7−δ is developing prevalentely along the c-axis direction; for this
reason the minimization of the mismatch along the c-axis direction is energeti-
cally more relevant. The difficulty to identify unique qx and qz values in the low
intensity peak of the Ba2YNbO6 measured from the the 10 Hz deposited thin
film makes not possible the repetition of the calculations performed on the 1 Hz
deposited sample.
Analysing the (108)-(018) YBa2Cu3O7−δ peaks in figure 5.2 additional infor-
mation on the effects of the deposition rate on the YBa2Cu3O7−δ lattice is ob-
tained. The broadening of the peak related to the Ba2YNbO6 doping is slightly
more relevant in the Ba2YNbO6 doped YBa2Cu3O7−δ thin film deposited with
a laser repetition rate of 10 Hz. This indicates that the YBa2Cu3O7−δ region
affected by lattice distortions (buckling of the crystalline planes) is larger when
the deposition rate is higher. There are two possible explanation of this evidence,
in particular either a larger region of YBa2Cu3O7−δ may be affected by lattice
distortions around each nanoinclusion or the Ba2YNbO6 nanoinclusions density
may be higher. The phase that is more directly affected by diffusion processes,
and thus by the deposition rate, is the Ba2YNbO6. The Nb ions must in fact
diffuse on larger region of the surface in order to form the Ba2YNbO6 nanoinclu-
sions. For this reason the higher density of Ba2YNbO6 inclusions is more likely
the cause of the increased lattice distortion.
In figure ?? is shown a direct comparison of the (108)-(018) YBa2Cu3O7−δ
peaks in the different films, the indication of the peaks’ position as well as the
half maximum intensity isolines allow direct visualization of the effects of the
presence of Ba2YNbO6 on the YBa2Cu3O7−δ crystalline structure.
The broadening of the YBa2Cu3O7−δ peaks is not the only difference, the lat-
tice parameters measured from the reciprocal space map of the Ba2YNbO6 doped
YBa2Cu3O7−δ thin film deposited at 10 Hz is different from the one measured from
the film deposited at 1 Hz. In particular, from the (108)-(018) YBa2Cu3O7−δ peak
74
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
Figure 5.3: Half maximum intesity insoline and peak position comparison for the
(108)(018) YBa2Cu3O7−δ peaks from figure 5.1 and 5.2.
of the reciprocal space map measured on the film deposited at 1 Hz (qx = 2.58
nm−1 and qz = 6.85 nm−1) the lattice parameters ab = 0.388 nm and c = 1.168
nm are calculated, while from the (108)-(018) YBa2Cu3O7−δ peak of the recip-
rocal space map measured on the film deposited at 10 Hz (qx = 2.60 nm
−1 and
qz = 6.83 nm
−1) are calculated the lattice parameters ab = 0.385 nm and c =
1.171 nm. In conclusion the YBa2Cu3O7−δ in the film deposited at 1 Hz appears
to better match the SrTiO3 substrate while not being affected on the overall c-axis
spacing by the Ba2YNbO6 presence at all. On the other hand the YBa2Cu3O7−δ
in the film deposited at 10 Hz maintains the ab bulk value (complete relaxation)
but the c lattice parameter is slightly larger. This is an indication that the
lattice distortion generated by the Ba2YNbO6 doping is influencing a region of
the YBa2Cu3O7−δ lattice that is large enough to be noticed when measuring the
overall c lattice parameter value.
The information obtained from the analysis of the (108)-(018) YBa2Cu3O7−δ
peak measured in the reciprocal space is summarised in table 5.3.
75
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
Sample Peak coordinates Lattice parameters
qx / qz ab
2θ / ω c
Undoped YBa2Cu3O7−δ 2.60 nm−1 / 6.85 nm−1 0.385 nm
68.69 ◦ / 13.58 ◦ 1.168 nm
Doped YBa2Cu3O7−δ (1 HZ) 2.58 nm−1 / 6.85 nm−1 0.388 nm
68.60 ◦ / 13.66 ◦ 1.168 nm
Doped YBa2Cu3O7−δ (10 Hz) 2.60 nm−1 / 6.83 nm−1 0.385 nm
68.59 ◦ / 13.42 ◦ 1.171 nm
Table 5.3: (108)-(018) YBa2Cu3O7−δ peak analysis.
5.1.3 Nanostructure analysis: Atomic Force Microscopy
and Transmission Electron Microscopy
The analysis of the nanostructure reveals interesting differences between the
Ba2YNbO6 doped YBa2Cu3O7−δ thin film deposited at 10 Hz and those de-
posited at 1 Hz. Similarly to what was done in the previous section the sur-
face topography is discussed by analysing atomic force microscopy images while
the morphological distribution of the secondary pinning phase and the distortion
in the YBa2Cu3O7−δ lattice are studied with cross-section transmission electron
micrography.
5.1.3.1 Surface topography (Atomic Force Microscopy)
A comparison between the surface topography of samples deposited at different
rates is reported in figure 5.4, in particular the surfaces of samples of different
thickness (≈ 700 nm and 400 nm) deposited at 1 and 10 Hz are investigated.
The first evident effect of increasing the laser repetition rate is a the presence of
smaller growth grains. Comparing the surfaces of sample with similar thickness
(figure 5.4a and 5.4b or 5.4c and 5.4d) it is clearly shown that the crystalline
grains are smaller in the samples grown at 10 Hz repetition rate. As expected,
a reduced “free migration” time can then be associated with a reduced accretion
of the grains. The first result is that the thin film deposited with a repetition
rate of 10 Hz is formed of a higher number of smaller grains when compared to
a thin film deposited with a repetition rate of 1 Hz. The direct consequence of
76
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
Figure 5.4: Atomic Force Microscopy surface image (5 µm × 5 µm) of 5%mol
Ba2YNbO6 doped YBa2Cu3O7−δ thin films deposited at a substrate temperature
of 780 ◦C, repetition rate of 1 Hz and 10 Hz with a thickness of 700 nm and 400
nm. a) 1 Hz and 700 nm; b) 10 Hz and 700 nm; c) 1 Hz and 400 nm; d) 10 Hz
and 400 nm.
77
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
the morphology described is clearly shown by the comparison of the thicker films
surfaces, figure 5.4a and 5.4b. It is in fact clear that the connectivity shown
by the grains in figure 5.4b is reduced. Furthermore by the same comparison it
appears that the surfacial morphology of the thin film deposited at the higher
repetition rate is a scaled down version of the one shown by the film deposited
at 1 Hz. A similar analysis could be made comparing the surface morphologies
of the thinner films (figure 5.4c and 5.4d) but it is difficult to visualise the grains
boundaries in the smooth surface of the thinner film deposited at 1 Hz.
With a more accurate analysis of the images another feature can be deduced.
In particular, as expected, the grain dimension is not affected by the film thick-
ness. Focusing on the thin films deposited with a repetition rate of 10 Hz (figure
5.4b and 5.4d) the grains dimension appears to be similar. However the shape of
the boundaries turn from the well defined straight lines of the thinner film (figure
5.4d) to the more chaotic boundaries of the thicker sample (figure 5.4b). This
is a evidence that the grains connectivity is also reduced when the thickness is
increased.
A fine dispersion of superficial nanoparticles is shown in figure 5.4c. Simi-
lar nanoparticles where also evidenced in the previous section (figure 4.6d) and
where attributed to the Ba2YNbO6 presence. It is impossible to picture the same
small nanoparticles in atomic force micrography taken on samples with a higher
roughness (figure 5.4a and 5.4b) because the increased scale flatten the small
topographic features. It is unclear why the superficial nanoparticles are absent
(or invisible) in the thinner film deposited with a repetition rate of 10 Hz (figure
5.4d). Since its surface is shown in a similar scale to the 1 Hz film the absence of
the nanoparticles is an additional indication of a morphological difference induced
by the different deposition rate.
The root mean square roughness values calculated from the surface topogra-
phies shown in figure 5.4 are reported in table 5.4. The calculated roughness
values are higher in the 700 nm thick films then in the 400 nm thick films. How-
ever, there is no evidence of a direct influence of the repetition rate on these
values.
A last note on this surface topography study has to be made on the target
improved quality. As stated earlier a densified target was prepared to avoid the
78
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
Repetition Rate, Thickness Root mean square roughness
1 Hz, 400 nm 2.12 nm
10 Hz, 400 nm 7.12 nm
1 Hz, 700 nm 34.14 nm
10 Hz, 700 nm 21.41 nm
Table 5.4: Root mean square roughness calculated from the surface topography
of 5%mol Ba2YNbO6 doped YBa2Cu3O7−δ thin films deposited at a substrate
temperature of 780 ◦C, repetition rate of 1 Hz and 10 Hz with a thickness of 700
nm and 400 nm.
presence of particulates on the film surface. Comparing the surface topographies
shows in this section with the earlier study reported (section 4.2.3.1) it is evident
that the changes in the production process of the pulsed laser deposition targets
solved the particulate issue.
5.1.3.2 Cross-section transmission electron microscopy
The cross-section images taken using the transmission electron microscopy clearly
show the strong influence of the laser pulse repetition rate on the nanostructure
of the Ba2YNbO6 doped YBa2Cu3O7−δ thin films.
In figure 5.5 are reported the transmission electron micrographs taken on
the films grown at 1 Hz and 10 Hz at different magnification. Columnar de-
fects are shown with a good contrast only in the film deposited at 1 Hz (fig-
ure 5.5a and 5.5b) and are indicated by arrows parallel to the c-axis of the
YBa2Cu3O7−δ. In the film deposited at 10 Hz some shadows parallel to the c-
axis of the YBa2Cu3O7−δ are present and could be related to the presence of
Ba2YNbO6 nanorods. Nevertheless the contrast in figure 5.5c and 5.5d is lower
and these columnar defects are not clearly identified. The reduced contrast can
be related to a lower orientation degree as well as to larger distortion. Once more
the Ba2YNbO6 inclusions nanostructure is shown to be the phase influenced more
by the deposition rate.
Focusing on the images related to the Ba2YNbO6 doped YBa2Cu3O7−δ thin
film deposited at laser pulse rate of 1 Hz, it is possible to analyse the morphology
of the nanorods parallel to the c-axis of the YBa2Cu3O7−δ. The nanorods diam-
eter is ≈ 10 nm and their spacing is ≈ 40 nm. Similarly to the case reported in
79
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
Figure 5.5: Transmission electron microscope cross-sectional image of a 5%mol
Ba2YNbO6 doped YBa2Cu3O7−δ thin film deposited with a laser pulse rates of 1
Hz (a, b) and 10 Hz (c, d). White arrows in the direction of the YBa2Cu3O7−δ
c-axis mark the position of self assembled Ba2YNbO6 nanorods parallel to the
c-axis. Arrows parallel to the YBa2Cu3O7−δ ab-planes mark the position of
nanoparticles. (TEM images from Prof. H Wang research group at Texas A&M
University).
80
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
the last section where a Ba2YNbO6 doped YBa2Cu3O7−δ deposited with a laser
pulse rate of 5 Hz showed same structures with similar spacing. Repeating once
more the matching field calculation also the films deposited at 1 Hz should have
a matching field value between 1.29 T and 1.49 T. Comparing figure 5.5a with
figure 4.8a two differences are evident: the nanorods in the film deposited at a
laser pulse rate of 1 Hz appears to be continuous and longer than the ≈ 100 nm
length of the nanorods showed in the film deposited at 5 Hz; in the film deposited
at 1 Hz there is no sign of the nanoparticles that where shown in the 5 Hz case.
The film deposited at a laser pulses rate of 10 Hz (figure 5.5c and 5.5d) is the
one that differs the most. The nanorods are not clearly visible but a large number
of nanoparticles is present. In general from the transmission electron microscopy
analysis this film appears to have the less ordered YBa2Cu3O7−δ lattice and an
almost completely disordered Ba2YNbO6 nanostructure.
It is interesting to see that the nanoparticles and nanorods presence is largely
influenced by the laser pulses rate. The nanoparticles are absent in the 1 Hz films,
their presence increase increasing the laser pulse rate up to the 10 Hz film where
they appear to be the dominant nanoinclusion. On the other hand the nanorods
are long, linear and almost perfectly parallel to the YBa2Cu3O7−δ c-axis in the
film deposited at a laser pulse of 1 Hz; are shorter (≈ 100 nm) with a large angle
spread to the YBa2Cu3O7−δ c-axis in the film deposited at a laser pulse of 5 Hz
(see section 4.2.3.2); are distorted and not clearly imaged in the film deposited
at a laser pulse of 10 Hz.
The Ba2YNbO6 inclusions generate distortion in the YBa2Cu3O7−δ lattice.
The influence of the laser pulse rate on the nanostructure of the Ba2YNbO6
inclusions is then transferred to the YBa2Cu3O7−δ lattice distorted by those in-
clusions. In transmission electron microscope images of the Ba2YNbO6 doped
YBa2Cu3O7−δ films deposited at pulse rates of 1 Hz and 5 Hz the YBa2Cu3O7−δ
crystalline planes parallel to the ab direction appear to be distorted only in the re-
gions close the nanorods and the planes buckling in this small region appear as an
effective strain sink. On the other hand a similar buckling of the planes does not
appear to be effective in respect of the high density of the nanoparticles growing
in the Ba2YNbO6 doped YBa2Cu3O7−δ deposited at 10 Hz. The YBa2Cu3O7−δ
lattice appears to be distorted with continuity and the entire lattice seems to be
81
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
affected by the Ba2YNbO6 presence. The high density of nanoparticles appears
to be the most likely cause. In particular, the reduced distance between the
Ba2YNbO6 inclusions do not allow full relaxation of the YBa2Cu3O7−δ lattice.
5.1.4 The superconducting properties: Tc, Jc(B) and Jc(B, θ)
The crystalline structure and the nanostructure of Ba2YNbO6 doped YBa2Cu3O7−δ
thin films deposited at different rates have been analysed. In particular the nanos-
tructure was found to be heavily influenced by the laser pulse rate. In this sec-
tion the superconducting transport properties (Tc, Jc(B) and the Jc(B, θ)) are
analysed and related to the nanostructural feature evidenced. In addition to the
Ba2YNbO6 doped YBa2Cu3O7−δ thin films deposited at 1 Hz and 10 Hz also films
deposited at 5 Hz where characterized. The crystalline structure and nanostruc-
ture morphology of Ba2YNbO6 doped YBa2Cu3O7−δ thin films deposited at 5 Hz
was analysed in the preliminary study. Nevertheless, since the transport proper-
ties characterisation performed in the preliminary study presented was limited to
low field values (Jc(B) up to 1 T and Jc(B, θ) measured at 0.5 T), the transport
properties of new films deposited at laser pulses repetition rate of 5 Hz adopting
the densified pulsed laser deposition target are characterised in order to obtain a
more complete picture of the effects of the deposition rate.
5.1.4.1 Transition temperature, Tc
The transition temperature does not appear to be effected by the laser pulses repe-
tition rate adopted. The Tc measured from all the Ba2YNbO6 doped YBa2Cu3O7−δ
thin films produced is reduced by a small amount (≈ 3 K) when compared to
the Tc measured from a pure YBa2Cu3O7−δ thin film.
In figure 5.6 are reported the resistance measured on a 50 µm width current
track patterned on pure YBa2Cu3O7−δ thin film and 5%mol Ba2YNbO6 doped
YBa2Cu3O7−δ thin films deposited at a laser pulses repetition rate of 1, 5 and 10
Hz. The transition temperature Tc, measured as the temperature at which the
resistance disappears, is reported in table 5.5.
The Tc reduction is in line with the one measured in the preliminary study
(section 4.2.4.1), thus the same considerations can be done and the Ba2YNbO6
82
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
Figure 5.6: Resistance variation with the temperature measured on a pure
YBa2Cu3O7−δ thin film deposited at a substrate temperature of 780 ◦C and laser
pulses repetition rates of 1 Hz and on 5%mol Ba2YNbO6 doped YBa2Cu3O7−δ
thin films deposited at a substrate temperature of 780 ◦C and laser pulses repe-
tition rates of 1, 5 and 10 Hz.
83
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
stability within the YBa2Cu3O7−δ matrix is confirmed. Furthermore it is impor-
tant to notice that despite the higher disorder pictured by the transmission elec-
tron microscopy the transition temperature measured on the Ba2YNbO6 doped
YBa2Cu3O7−δ thin film deposited at 10 Hz is almost equal to the one measured
on the thin film deposited at a laser pulses repetition rate of 1 Hz. In conclusion
the fact that the higher defects density is not related to a further reduction of
the transition temperature could be the evidence that connected regions of the
YBa2Cu3O7−δ lattice only marginally distorted by the Ba2YNbO6 inclusion may
be present also in the thin films deposited at laser pulses repetition rate of 10 Hz.
Thin Film Transition Temperature Tc
pure YBa2Cu3O7−δ (1 Hz) 91.5 K
Ba2YNbO6 doped YBa2Cu3O7−δ (1 Hz) 87.8 K
Ba2YNbO6 doped YBa2Cu3O7−δ (5 Hz) 88.8 K
Ba2YNbO6 doped YBa2Cu3O7−δ (10 Hz) 88.0 K
Table 5.5: Transition Temperature Tc collected from the resistance variation with
temperature measurements shown in figure 5.6.
5.1.4.2 The critical current density
Once again despite the slightly reduced transition temperatures Tc the first ev-
ident effect of the Ba2YNbO6 doping is improved critical current densities Jc
values over the whole field range analysed.
In figure 5.7 are reported Jc(B) measurements up to 6 T with the field applied
parallel to the c-axis (B‖c) at 77 K for a pure YBa2Cu3O7−δ thin film and 5%mol
Ba2YNbO6 doped YBa2Cu3O7−δ thin films deposited at a laser pulses repetition
rate of 1, 5 and 10 Hz. The Jc(B) values measured on the Ba2YNbO6 doped
YBa2Cu3O7−δ thin films deposited at laser pulses repetition rate of 1 and 5 Hz
are similar. This is an evidence that the pinning potential is similar when the
field is applied parallel to the YBa2Cu3O7−δ c-axis.
The Jc values measured on the film deposited at 10 Hz are lower then those
measured on the other Ba2YNbO6 doped YBa2Cu3O7−δ films in the low field
region (B ≤ 2 T) but are the highest when the applied field is above 2 T. A
possible explanation can be found in the higher density of Ba2YNbO6 nanoinclu-
84
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
Figure 5.7: Critical current density variation with the applied magnetic field value
measured on a pure YBa2Cu3O7−δ thin film deposited at a substrate temperature
of 780 ◦C and laser pulses repetition rates of 1 Hz and on 5%mol Ba2YNbO6
doped YBa2Cu3O7−δ thin films deposited at a substrate temperature of 780 ◦C
and laser pulses repetition rates of 1, 5 and 10 Hz. The field is applied parallel
to the c-axis of the YBa2Cu3O7−δ, T = 77 K. The data are on a log-linear plot
(a) and a log-log plot (b).
85
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
sions. The matching field values (1.29 T ≤ H∗ ≤ 1.59 T) calculated from the
Ba2YNbO6 nanorods were similar in the 1 Hz and 5 Hz films. Unfortunately, due
to the high distortions resulting in low transmission electron image contrast it
was not possible to calculate a matching field for the 10 Hz film. However, since
the lower spacing of the inclusion was evidenced, a higher matching field value
was expected.
Figure 5.8: Pinning force calculated from data reported in figure 5.7
To better visualize the effect of different matching field values in figure 5.8
is reported the pinning force calculated as Jc(B) × B from the data reported in
figure 5.7. The pinning force maximum for the Ba2YNbO6 doped YBa2Cu3O7−δ
thin films deposited at 1 and 5 Hz are located at a field valued of ≈ 1.4 T.
It is worth noticing that this value is perfectly in line with the matching field
values predicted from the nanorods spacing measured on the images acquired
with the electron transmission microscopy. The fact that the pinning force max-
imum for the Ba2YNbO6 doped YBa2Cu3O7−δ thin film deposited at 10 Hz is
shifted towards higher field values confirm the higher matching field value that
was expected from the higher Ba2YNbO6 inclusion density.
Table 5.6 reporting the α values calculated from data in figure 5.7 complete
86
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
Thin Film α
pure YBa2Cu3O7−δ (1 Hz) 0.47
Ba2YNbO6 doped YBa2Cu3O7−δ (1 Hz) 0.32
Ba2YNbO6 doped YBa2Cu3O7−δ (5 Hz) 0.37
Ba2YNbO6 doped YBa2Cu3O7−δ (10 Hz) 0.39
Table 5.6: α calculated from the data reported in figure 5.7.
the Jc(B) analysis. The α values are reduced when the Ba2YNbO6 is added
to YBa2Cu3O7−δ and the higher value is the one calculated for the Ba2YNbO6
doped YBa2Cu3O7−δ deposited at a laser pulse repetition rate of 10 Hz is α = 0.39
which is lower than α = 0.47 calculated for the pure YBa2Cu3O7−δ sample. The
reason why the highest α value is the one calculated the 10 Hz Ba2YNbO6 doped
YBa2Cu3O7−δ film is that the pinning landscape generated in this film is effective
at high field more then it is at low field and that α is calculated in the very low field
regime (B ≤ 1 T). At low field values the pinning potential is not fully developed
since the high density of nanoinclusions may reduce the energy associated with
flux jumps between adjoining defects reducing the pinning potential.
To complete the analysis of the effects that different laser pulses repetition
rates have on the nanostructure and properties of Ba2YNbO6 doped YBa2Cu3O7−δ
thin films the critical current density variation as a function of the external mag-
netic field direction is measured. In figure 5.9 are reported Jc(B, θ) measurements
at 77 K and 0.5, 1 T for a pure YBa2Cu3O7−δ thin film and 5%mol Ba2YNbO6
doped YBa2Cu3O7−δ thin films deposited at a laser pulses repetition rate of 1, 5
and 10 Hz.
The first important result is that even if all the Ba2YNbO6 doped YBa2Cu3O7−δ
thin films shows an higher Jc value when the field is applied parallel to the
YBa2Cu3O7−δ c-axis the only thin film that shows higher Jc values than pure
YBa2Cu3O7−δ when the field is applied parallel to the ab-planes is the thin film
deposited at a laser pulses repetition rate of 1 Hz. It is evident that the pinning
potential along the ab-planes of the Ba2YNbO6 doped YBa2Cu3O7−δ thin film de-
posited at 5 and 10 Hz is reduced. The cause of this reduction can not be indicated
in the nanorods inclusion because these are the dominant inclusion type in the
Ba2YNbO6 doped YBa2Cu3O7−δ deposited at 1 Hz and its ab-direction pinning
87
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
Figure 5.9: Critical current density variation with the direction of the applied
magnetic field measured on a pure YBa2Cu3O7−δ thin film deposited at a sub-
strate temperature of 780 ◦C and laser pulses repetition rates of 1 Hz and on
5%mol Ba2YNbO6 doped YBa2Cu3O7−δ thin films deposited at a substrate tem-
perature of 780 ◦C and laser pulses repetition rates of 1, 5 and 10 Hz. a) Applied
magnetic flux density = 0.5 T, T = 77 K; b) Applied magnetic flux density = 1
T, T = 77 K.
88
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
potential is similar if not better then the one shown by the pure YBa2Cu3O7−δ
film. The ab-direction reduction of the Jc could be associated with the lower con-
nectivity evidenced in the Ba2YNbO6 doped YBa2Cu3O7−δ thin films deposited
at 5 Hz and 10 Hz. It is certain that the loss of pinning capability when the
field is applied parallel to the YBa2Cu3O7−δ ab-planar direction is related to the
interruption of the crystalline ab-planes continuity. There are two main sources
of planar continuity interruption: the nanorods and the grain boundaries. The
nanorods generate a local hole through the planes and a possible loss of conti-
nuity in the small region around them that is affected by planar buckling. The
grain boundaries are a two dimensional interruption of the planar continuity that
is extended over lengths that are order of magnitude larger that the nanorods
interruption. This together with the evidence that one of the main effects of
higher repetition rates is the reduction of the growth grains dimension (figure
5.4) suggest that the deficient pinning potential measured when the field is ap-
plied parallel to the ab-planar direction is caused by the interruption of the ab
crystalline planes occurring at the grain boundaries.
It is worth remembering that the scenario presented for the Ba2YNbO6 doped
YBa2Cu3O7−δ deposited at laser pulses repetition rates of 5 and 10 Hz is com-
mon in pinning enhanced YBa2Cu3O7−δ thin films in which the secondary phases
forms heteroepitaxial nanorods [57, 65, 74, 118] while the presence of a strong c-
axis pinning without a reduction of the ab-planes pinning was achieved only in
carefully tuned BaZrO3 doped YBa2Cu3O7−δ [6, 134]. In general while it is com-
mon to increase the isotropic pinning without disrupting the ab-planar continuity,
it is complex to introduce a strong anisotropic pinning along the c-axis direction
without depressing the planar Jc.
The Ba2YNbO6 doped YBa2Cu3O7−δ thin film deposited at a laser pulses
repetition rate of 10 Hz shows a strong c-axis pinning potential revealing the
presence of c-axis directional defects that were not clearly shown by the trans-
mission electron microscopy. Furthermore the isotropic pinning introduced by
the densely packed nanoparticles is not effective at low field (B ≤ 1 T) and
at these field values the connectivity issues and low energies flux jumps nullify
the pinning effects. Similarly the pinning enhancement that should be provided
by the simultaneous presence of nanorods and nanoparticles in the Ba2YNbO6
89
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
doped YBa2Cu3O7−δ deposited at 5 Hz are also probably denied by the reduced
connectivity.
5.1.5 Concluding remarks to the analysis of the effects of
the deposition rate on the nanostructure and the
superconducting properties
The analysis performed in this section demonstrated that the laser pulses repeti-
tion rate heavily influence the YBa2Cu3O7−δ grains morphology and consequently
the overall connectivity and continuity of the ab crystalline planes. Higher repe-
tition rates generates smaller growth grains with reduced connectivity and planar
continuity. Such reduced connectivity and continuity is found to generate a re-
duction of the critical current values Jc when the field is applied parallel to the
YBa2Cu3O7−δ ab-planes.
The laser pulses repetition rate influences also the nanostructure of the
Ba2YNbO6. Long linear nanorods are obtained at low deposition rates (1 Hz), at 5
Hz are obtained shorter nanorods and simultaneously nanoparticles and at 10 Hz
the nanoparticles are the dominant Ba2YNbO6 nanostructure. It would be pos-
sible to tune the Ba2YNbO6 nanorods-nanoparticles ratio to achieve an optimal
pinning landscape by tuning the repetition rate as done with the BaZrO3 [134] but
the connectivity issues evidenced have a larger detrimental effect on the pinning
properties than the enhancing effect expected from the sinergetic combination of
0D nanoparticles and 1D c-axis oriented nanorods.
The best results in mid-low field (0 T to 2 T) values where obtained by
Ba2YNbO6 doped YBa2Cu3O7−δ thin films deposited at 1 Hz. On the other
hand the high density pinning landscape generated in the Ba2YNbO6 doped
YBa2Cu3O7−δ thin film deposited at 10 Hz provides enough pinning force to
overcome the related connectivity issue when the field applied is above 2 T; a
possible application niche for the 10 Hz thin films could then be found in special-
ized equipment for high magnetic field.
90
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
5.2 Ba2YNbO6 perovskite additions to YBa2Cu3O7−δ:
deposition parameter optimization
There are several parameters that define the process in a pulsed laser deposition.
The substrate - target distance, the laser fluence, the oxygen pressure, the sub-
strate temperature and the laser pulse repetition rate all influence the quality of
the deposited thin film. The substrate - target distance, the laser fluence and
the oxygen pressure are related more to the target ablation and plume diffusion
process than to the film growth, furthermore they have been optimized over the
years. The last two are the parameters that directly influence the film growth
process. A complete study on the effects of the laser puses repetition rate is
discussed in the previous section, while in this brief final section of the chapter
the study of the properties of the Ba2YNbO6 doped YBa2Cu3O7−δ thin film is
concluded with the optimisation of the substrate temperature in the pulsed laser
deposition. This final section will not include an in depth study of the nanos-
tructure similar to the one reported in the previous sections of the chapter but
will focus on the grain connectivity and the superconducting properties of a large
number of films.
5.2.1 Ba2YNbO6 doped YBa2Cu3O7−δ pulsed laser depo-
sition target preparation and thin films deposition
The target used in the pulsed laser deposition is the one that was ablated to
deposit the thin films analysed in the previous section.
The substrate temperature was modified to study its effects on the growth and
the superconducting properties of Ba2YNbO6 doped YBa2Cu3O7−δ thin films
deposited at different laser pulses repetition rate. It is difficult to predict the
effects of the substrate temperatures on the film growth as it simultaneously
affects multiple mechanism of the process. As an example an higher substrate
temperature is certainly related to higher mobility of the ions moving on the
growing film surface and at the same time high substrate temperatures in a low
oxygen pressure environment may speed up an oxygen content reduction process.
Thus in order to identify a processability window, if not the optimum, a set
91
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
of Ba2YNbO6 doped YBa2Cu3O7−δ thin films deposited at different substrate
temperatures and different laser pulse repetition rate were grown. In particular
repetition rate of 1, 5 and 10 Hz and substrate temperatures of 740 and 760 ◦C
were investigated and compared to the results obtained on thin films deposited
at a substrate temperature of 780 ◦C. In the table 5.7 below are summarized the
pulsed laser deposition parameters adopted for the deposition of the Ba2YNbO6
doped YBa2Cu3O7−δ thin films analysed in this section.
Parameter Value
Substrate Temperature 740,760,780 ◦C
Chamber Pressure 0.3 mbar flowing O2
Laser Fluence 2 Jcm−2
Repetition Rate 1, 5, 10 Hz
Number of Pulses 4500
Annealing Time 1 hr
Annealing Temperature 520 ◦C
Annealing Pressure 500 mbar O2
Table 5.7: Pulsed laser deposition parameters
All the Ba2YNbO6 doped YBa2Cu3O7−δ thin films deposited were measured
to be of ≈ 0.4 µm thickness.
5.2.2 Surface Topology: Grain Connectivity
The analysis discussed in the previous section showed that the grain connectivity
has a major influence on the superconducting properties, in particular when the
magnetic field is applied parallel to the YBa2Cu3O7−δ ab-planar direction the crit-
ical current values appear to be reduced in the Ba2YNbO6 doped YBa2Cu3O7−δ
thin films featuring small growth grains. For this reason a study of surface to-
pography can not be neglected even in this brief analysis of the superconducting
properties.
In order to provide clear visual impact the images of the surface topography
were separated into three. In figure 5.10 are reported the surface topographies for
all the thin films produced with a substrate temperature of 740 ◦C, in figure 5.11
those deposited at 760 ◦C and in figure 5.12 those at 780 ◦C. In all the figures
from top to down the laser pulses rate varies from 1 Hz to 10 Hz.
92
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
Figure 5.10: Atomic Force Microscopy surface image (5 µm × 5 µm) of 5%mol
Ba2YNbO6 doped YBa2Cu3O7−δ thin films deposited at a substrate temperature
of 740 ◦C; a) repetition rate of 1 Hz; b) repetition rate of 5 Hz; c) repetition rate
of 10 Hz;
93
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
Figure 5.11: Atomic Force Microscopy surface image (5 µm × 5 µm) of 5%mol
Ba2YNbO6 doped YBa2Cu3O7−δ thin films deposited at a substrate temperature
of 760 ◦C; a) repetition rate of 1 Hz; b) repetition rate of 5 Hz; c) repetition rate
of 10 Hz;
94
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
Figure 5.12: Atomic Force Microscopy surface image (5 µm × 5 µm) of 5%mol
Ba2YNbO6 doped YBa2Cu3O7−δ thin films deposited at a substrate temperature
of 780 ◦C; a) repetition rate of 1 Hz; b) repetition rate of 5 Hz; c) repetition rate
of 10 Hz;
95
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
The first clear feature evidenced by these atomic force microscopy study is the
presence of outgrowths on the surfaces of the thin films deposited at a substrate
temperature of 740 ◦C (figure 5.10). These outgrowths where recently reported
in the case of pure YBa2Cu3O7−δ on RABiTS for substrate temperatures from
740 ◦C and below [135]. Nevertheless several authors have investigated and char-
acterized the outgrowths on the surface of pulsed laser ablated YBa2Cu3O7−δ on
different substrates [136,137].
The outgrowths are reported to be a-axis YBa2Cu3O7−δ grains, misaligned c-
axis grains and non superconducting grains made by Y2BaCuO5, CuYO2, CuO,
BaCuO2 and Ba2CuO3. It is possible to distinguish between the outgrowths
formed by superconducting YBa2Cu3O7−δ grains and those formed by non su-
perconducting phases with the atomic force microscopy. The outgrowths formed
by YBa2Cu3O7−δ (a-axis YBa2Cu3O7−δ and misaligned c-axis YBa2Cu3O7−δ) are
reported as block-shaped while those formed by non superconducting phases are
semi-spherical. From the images reported in figure 5.10 it is evident the out-
growths are block-shaped thus are YBa2Cu3O7−δ grains. The superconducting
outgrowths are described as predominantly c-axis YBa2Cu3O7−δ grains paral-
lel to the substrate surface that nucleate on the substrate at the same time of
the desired c-axis YBa2Cu3O7−δ perpendicular to the substrate surface when the
substrate temperature is below a certain threshold.
In the present case this lower temperature limit is found to be 740 ◦C. Further-
more a significant increase in the size of the outgrowths is also reported to occurs
during long deposition. This size increment is related to precipitates formation
at the edge of the outgrowths as well as to the merging of neighboring c-axis
parallel grains. A similar behaviour is found in figure 5.10 where the size of the
outgrowths on the surface of the thin film deposited at a laser pulses repetition
rate of 1 Hz is much larger than that observed for the thin films deposited at 5
and 10 Hz. Although the quality of these films is reduced by the presence of these
outgrowths, it is significant to make a final comment on the dimension and shape
of the c-axis perpendicular grains that are visible below the outgrowths. These
grains are almost perfectly rectangular, their size is ≈ 300 nm and do not seem
to change with the variation of laser pulses repetition rate. From the images in
figure 5.11 and 5.12 it is evident that the outgrowths are absent.
96
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
Figure 5.13: Average grain size measured from AFM analysis of 5%mol
Ba2YNbO6 doped YBa2Cu3O7−δ thin film deposited at a substrate temperature
of 740, 760 and 780 ◦C as a function of laser pulse rate
Figure 5.13 compares the average grain size measured from the AFM in figure
5.10, 5.11 and 5.12, which shows a well defined trend. The size of the grains and
thus the connectivity increases when the substrate temperature is increased and
when the laser pulses repetition rate is decreased. It is important remembering
that a reduction of the pulses laser repetition rate corresponds to an increase in
the maximum time allowed to the ions migration on the surface. Referring to
a standard model of nucleation and growth, the grain size analysis shows that
an increase in ions mobility as well as the time allowed to their migration shifts
the balance between nucleation and growth towards the latter. In the previous
section it was shown that thin films composed of larger grains have better super-
conducting properties than films composed by smaller grains. Therefore, if this
trend were to be confirmed, the films with the best superconducting properties
should be those deposited with a substrate temperature of 780 ◦C.
A last consideration can be made comparing the surface topographies of the
thin film deposited at 1 Hz and 760 ◦C (figure 5.11a) with the one deposited at 5
Hz and 780 ◦C (figure 5.12b). It is evident that the two images are similar thus
indicating the possibility to achieve similar results with lower mobility (lower
substrate temperature) and longer migration times (lower laser pulse repetition
97
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
rate) or higher mobility and smaller migration times.
5.2.3 The superconducting properties: Tc, Jc(B) and Jc(B, θ)
To conclude the last section of the chapter are shown the results obtained from
the analysis of the superconducting properties. The variation of transition tem-
perature Tc and critical current density as a function of the applied magnetic field
intensity Jc(B) and as a function of the applied magnetic field direction Jc(B, θ)
are reported.
5.2.3.1 Transition temperature, Tc
In figure 5.14 is visualized the Tc variation with the substrate temperature. Each
curve is referred to a laser pulse repetition rate. The Tc values are also summarized
in tabular form in table 5.8.
Figure 5.14: Visualisation of the transition temperature values measured for
5%mol Ba2YNbO6 doped YBa2Cu3O7−δ thin films deposited at a substrate tem-
perature of 740, 760 and 780 ◦C and laser pulses repetition rates of 1, 5 and 10
Hz
The transition temperature values measured for the thin films deposited at a
98
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
substrate temperature of 760 ◦C has the same behaviuor discussed in the previous
section for the thin films deposited on the substrate kept at 780 ◦C. The Tc
reduction is ≈ 3 K as Tc ≈ 88 K are measured for all the samples.
A different Tc trend is observed in thin films deposited on substrates kept at
740 ◦C. In these thin films the Tc values appears to be influenced by the laser
pulses repetition rate and the Tc measured for the thin films deposited at 10 Hz
shows a severe reduction down to ≈ 79 K. The main difference in these films is
the presence of the outgrowths discussed earlier in the section. It is possible that
the YBa2Cu3O7−δ c-axis grains aligned parallel to the substrate surface introduce
high levels of strain and distortion reducing the Tc. The fact that the size of these
outgrowths changes with the variation of the laser pulses repetition rate could
be the reason why only in when these outgrowth are present the Tc values also
changes with the variation of the repetition rate.
A possible explanation to the Tc variation with the laser pulses repetition rate
can be found considering that the strain and distortion are responsible for the
reduced Tc values. These are generated at the interfaces between two phases (or c-
axis parallel YBa2Cu3O7−δ grains and c-axis perpendicular YBa2Cu3O7−δ grains)
thus larger interfaces induce larger Tc reduction. Taking into account that the
an increase of the laser pulses repetition rate decrease the sizes of the outgrowths
but at same time increase their density and thus increase the interfacial area
(figure 5.10), the relation between Tc values and the laser pulses repetition rates
is evident.
1 Hz 5 Hz 10 Hz
740 ◦C 87.8 K 85K 79 K
760 ◦C 87.9 K 89 K 88.2 K
780 ◦C 87.8 K 88.8 K 88.8 K
Table 5.8: Transition temperature values for 5%mol Ba2YNbO6 doped
YBa2Cu3O7−δ thin films deposited at a substrate temperature of 740, 760 and
780 ◦C and laser pulses repetition rates of 1, 5 and 10 Hz.
99
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
5.2.3.2 The critical current density
The critical current values measured on thin films deposited on substrate kept at
temperature values smaller then 780 ◦C are all reduced (figure 5.15). A reduction
of the substrate temperatures appears to reduce the critical current values regard-
less of the repetition rate. As a matter of the fact that the detrimental effects
generated by connectivity issues has been evidenced and discussed in the previous
section and that the atomic force microscopy analysis revealed that larger grains
(better connectivity) where generated at high substrate temperatures a similar
behaviour of the critical current values was expected.
Figure 5.15: Critical current density variation with the applied magnetic field
value measured on 5%mol Ba2YNbO6 doped YBa2Cu3O7−δ thin films deposited
at a substrate temperature of 760 and 780 ◦C and laser pulses repetition rates of
1, 5 and 10 Hz. The field is applied parallel to the c-axis of the YBa2Cu3O7−δ,
T = 77 K.
A self explanatory summary plot is reported in figure 5.16, the critical current
density values measured at 0.5 T and 1 T are plotted as a function of the pulse
rate and for the samples deposited at a substrate temperature of 760 ◦C and 780
◦C.
In figure 5.17 the critical current density variation with direction of the applied
100
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
Figure 5.16: Critical current density measured at 0.5 T and 1 T on 5%mol
Ba2YNbO6 doped YBa2Cu3O7−δ thin films deposited at a substrate tempera-
ture of 760 and 780 ◦C as a function of laser pulse rate. The field is applied
parallel to the c-axis of the YBa2Cu3O7−δ, T = 77 K.
magnetic field is reported. These data do not add new knowledge on the system
analysed and the consideration done in the last section could be repeated. Once
again the relevant results is that the best performances are obtained by the thin
films deposited on substrate kept at 780 ◦C.
An interesting observation is obtained by the comparison of the Jc(B) and
the Jc(B, θ) of the thin film deposited at 760
◦C and 1 Hz with those of the thin
film deposited at 780 ◦C and 5 Hz. The Jc values obtained in these films are
similar and also the topographies were similar. This is an important evidence
that, as stated before, it is possible to reduce the deposition time (increase the
laser pulses repetition rate) without affecting the superconducting properties in-
creasing the ions mobility (increase the substrate temperature). Unfortunately,
even if the above is true, the substrate temperature (mobility) has an upper limit.
When temperatures above 780 ◦C are used new phenomena are introduced and
the superconductive properties are worsened probably by reductions in oxygen
content. In particular Ba2YNbO6 doped YBa2Cu3O7−δ thin films deposited on
substrate kept at temperatures of 800 and 820 ◦C shows Tc values below 80 K.
101
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
Figure 5.17: Critical current density variation with the direction of the applied
magnetic field measured on 5%mol Ba2YNbO6 doped YBa2Cu3O7−δ thin films
deposited at a substrate temperature of 760 and 780 ◦C and laser pulses repetition
rates of 1, 5 and 10 Hz. a) Applied magnetic flux density = 0.5 T, T = 77 K; b)
Applied magnetic flux density = 1 T, T = 77 K.
102
5. Ba2YNbO6 doped YBa2Cu3O7−δ: the deposition parameters
5.2.4 Concluding remarks to the deposition parameters
optimization
In this last section the relevance of grain size on the Ba2YNbO6 doped YBa2Cu3O7−δ
thin films is once again evidenced. The connection between the temperature at
which the substrate is kept during the deposition with the films grains morphol-
ogy is analysed.
Within the parameter range analysed the optimal superconducting properties
for the Ba2YNbO6 doped YBa2Cu3O7−δ system are found in the thin films that
are deposited on substrates kept at 780 ◦C with a laser pulse repetition rate of 1
Hz.
A temperature of ≈ 780 ◦C seems to be an optimal value since further
increments require additional processing to equilibrate the oxygen content. If post
annealing processes are not considered a substrate temperature ≈ 780 ◦C is an
upper limit on the processability parameter variation range. On the other hand,
the laser pulses repetition rate of 1 Hz is the lowest rate analysed and the trend
observed indicated that reducing the repetition rate would lead to larger growth
grains and more ordered nanorods arrays. However, it is important to notice
that even if there is not lower limit of processability for the laser pulse repetition
rate, there is a lower efficacy limit. In other words adopting increasingly smaller
repetition rate at a certain point will cease to be beneficial and from industry
point of view will start to be disadvantageous. In fact, once the process ceases
to be kinetically limited and reaches the thermodynamic equilibrium any further
reductions in the repetition rate would have no influence.
103
Chapter 6
Ba2YNbO6 and Gd3TaO7
simultaneous doping of
YBa2Cu3O7−δ
In the previous chapter an in-depth analysis of the effects of Nb addition to
YBa2Cu3O7−δ is presented. It is proved that niobium additions to YBa2Cu3O7−δ
produce Ba2YNbO6 nanorods and that in general additions to REBa2Cu3O7−δ
produce Ba2RENbO6 nanorods [5,118,138] which are similar to BaZrO3 nanorods,
≈ 10 nm in diameter and ≈ 100 nm in length but with a larger splay than
BaZrO3 around the c-axis.
Another element, the tantallum, also showed similar properties to the nio-
bium. Both niobium and tantallum are highly charged ions and do not substitute
in the YBa2Cu3O7−δ lattice but form non superconducting phases. One of the
phases that has been reported to form in the YBa2Cu3O7−δ doped with tanta-
llum is the Ba2YTaO6 perovskite with a crystalline lattice that is identical to the
Ba2YNbO6 [124]. Referring to the Ba2YNbO6 and Ba2YTaO6 perovskites the Nb
and Ta are present as Nb+4 and Ta+4 coordinating 6 oxygen in an octahedral
lattice site and the ionic radius of both the element is 0.082 nm. As a matter of
the fact the Ba2YNbO6 and Ba2YTaO6 have identical crystalline lattices.
On the other hand Ba2YTaO6 is not the only phase reported to be gener-
ated by the tantallum addition, a defective pyrochlore RE3TaO7 has also been
104
6. Ba2YNbO6 and Gd3TaO7 simultaneous doping of YBa2Cu3O7−δ
reported [67]. Despite the different chemical composition and crystalline struc-
ture attributed to the nanorods formed by the addition of tantallum these have
been described as very fine (≈ 5 nm in diameter), continuous over the entire film
thickness, highly linear and densely distributed thus different from the wider (≈
10 nm in diameter), shorter and less linear nanorods which are formed by the
niobium addition.
Since the niobate and tantalate nanorods in YBa2Cu3O7−δ are of rather dif-
ferent morphology the YBa2Cu3O7−δ doped with niobium and the YBa2Cu3O7−δ
doped with tantallum have different superconducting properties. It is interest-
ing to evaluate whether the addition of both niobium and tantallum of the same
overall doping level results in an averaging effects or in an increased complexity
of the system that could yield to an entirely different and new pinning landscape.
In this chapter is reported a study of simultaneous doping of YBa2Cu3O7−δ
with both Ba2YNbO6 and Gd3TaO7. The reason tantallum was added in the
form of RE3TaO7 is that this was previously studied in our group as a pinning
additive to YBa2Cu3O7−δ thin films.
6.1 Ba2YNbO6 and Gd3TaO7 doped YBa2Cu3O7−δ
target preparation and thin films deposition
The Ba2YNbO6 and Gd3TaO7 doped YBa2Cu3O7−δ target for the pulsed laser
deposition was sintered adopting the process described in the previous chapter to
produce densified target and minimize the particulate amount deposited on film
surfaces (section 5.1.1). As anticipated in chapter 3, unlike the Ba2YNbO6 per-
ovskite, the Gd3TaO7 powders were not prepared prior to the target sintering and
the desired amount of 99.99% Gd2O3 and 99.99% Ta2O5 powders were directly
added to a mixture of YBa2Cu3O7−δ and Ba2YNbO6 powders. The final target
stochiometry is 2.5%mol Ba2YNbO6 + 2.5%mol Gd3TaO7 doped YBa2Cu3O7−δ
doped YBa2Cu3O7−δ.
The pulsed laser deposition parameters chosen for the thin films analysed in
this chapter are those that showed optimised growth for the 5%mol Ba2YNbO6
doped YBa2Cu3O7−δ. In particular, the substrate temperature was kept at 780
105
6. Ba2YNbO6 and Gd3TaO7 simultaneous doping of YBa2Cu3O7−δ
◦C and the laser pulses repetition rate was set at 1 Hz. In table 6.1 is summarized
the complete set of deposition parameters studied.
Parameter Value
Substrate Temperature 780 ◦C
Chamber Pressure 0.3 mbar flowing O2
Laser Fluence 2 Jcm−2
Repetition Rate 1 Hz
Number of Pulses 4500
Annealing Time 1 hr
Annealing Temperature 520 ◦C
Annealing Pressure 500 mbar O2
Table 6.1: Pulsed laser deposition parameters
The Ba2YNbO6 + Gd3TaO7 doped YBa2Cu3O7−δ thin films deposited were
measured to be of ≈ 0.35 µm thickness.
6.2 Crystalline structure analysis: x-ray diffrac-
tion data
Similarly to the crystalline structure characterisation performed on the first
Ba2YNbO6 doped YBa2Cu3O7−δ thin films produced the first results presented
in this chapter are those obtained from the x-ray diffraction analysis. It is funda-
mental to investigate the crystalline structure and the orientation of the phases
introduced in the thin films.
6.2.0.1 Crystalline phases identification
Figure 6.1 shows the x-ray diffraction pattern of a typical thin film.
The (00l) peaks of YBa2Cu3O7−δ (or (Y/Gd)Ba2Cu3O7−δ) as well as the
(002), (004) and (008) peaks of Ba2(Y/Gd)(Nb/Ta)O6 are labeled. A peak at
2θ = 33.3◦ can be identified as the (004) diffraction peak of (Y/Gd)2O3. Further-
more traces of a copper rich cuprate are also present as indicated by the diffraction
peaks (008),(0010) and (0012) at 2θ = 25.9◦, 2θ = 33.2◦ and 2θ = 41.4◦ that can
be associated to (Y/Gd)Ba2Cu4O8.
106
6. Ba2YNbO6 and Gd3TaO7 simultaneous doping of YBa2Cu3O7−δ
Figure 6.1: Bragg-Brentano scan of a YBa2Cu3O7−δ + 2.5%mol Ba2YNbO6 +
2.5%mol Gd3TaO7 thin film deposited on SrTiO3.
It is evident that Gd3TaO7 is absent and that a chemical reaction has gen-
erated new phases. It is possible to express a balanced chemical reaction that
summarize the phases transformation and that is consistent with the phases ob-
served by the x-ray diffraction reported in figure 6.1.
0.95 YBa2Cu3O7−δ + 0.025 Ba2YNbO6 + 0.025 Gd3TaO7
↓
0.85 (Y/Gd)Ba2Cu3O7−δ+ 0.05 Ba2(Y/Gd)(Nb/Ta)O6 + 0.0375 (Y/Gd)2O3
+ 0.075 (Y/Gd)Ba2Cu4O8
Considering the nature of the ions present in the system is easy to understand
the phases evolution observed. Gd can substitute onto the Y site since they
have similar ionic radii. For the same reason Nb and Ta can cross-substitute.
This allows Ta and Gd introduced as Gd3TaO7 together with the Ba2YNbO6
and Ba ions subtracted to the YBa2Cu3O7−δ to form the complex perovskite
Ba2(Y/Gd)(Nb/Ta)O6. Furthermore the Gd3TaO7 → Ba2(Y/Gd)(Nb/Ta)O6
transformation generates a Gd excess and a Ba deficiency that together with
the Gd ↔ Y mutual substitution could be at the origin of the (Y/Gd)2O3 and
107
6. Ba2YNbO6 and Gd3TaO7 simultaneous doping of YBa2Cu3O7−δ
(Y/Gd)Ba2Cu4O8 particles formation.
The previously reported [67] Gd3TaO7 phase does not form and the Ta partic-
ipate to the formation of a perovskite structure similar to the reported Ba2YTaO6
[68, 109]. Ba2YNbO6 and Ba2YTaO6 have identical crystalline structure and the
newly formed Ba2(Y/Gd)(Nb/Ta)O6 do not shows evident differences.
(Y/Gd)Ba2Cu4O8 formation has been previously observed in coated
YBa2Cu3O7−δ superconductors affected by barium deficiency. In particular a bar-
ium depletion by reaction with a CeO2 buffer layer has been reported to generate
(Y/Gd)Ba2Cu4O8 [139]. In the Ba2YNbO6 + Gd3TaO7 doped YBa2Cu3O7−δ
the (Y/Gd)Ba2Cu4O8 formation can be related to the barium deficiency in the
Gd3TaO7 reactant. (Y/Gd)Ba2Cu4O8 has been reported also in films grown by
metal-organic deposition, where it is described to be in the form of stacking fault
defects [140].
Concluding the phase identification it is clear that the Ba2YNbO6 + Gd3TaO7
doped YBa2Cu3O7−δ is a Ba2(Y/Gd)(Nb/Ta)O6 + (Y/Gd)2O3 doped
(Y/Gd)Ba2Cu3O7−δ thin film with additional traces of (Y/Gd)Ba2Cu4O8.
6.2.1 Crystalline phases orientation
The phase orientation analysis is reported in figure 6.2. The SrTiO3 curve is
omitted since the (Y/Gd)Ba2Cu3O7−δ can be taken as an orientation reference.
The cube on cube growth of the Ba2(Y/Gd)(Nb/Ta)O6 in the
(Y/Gd)Ba2Cu3O7−δ is evidenced by the x-ray φ scans: the (202)
Ba2(Y/Gd)(Nb/Ta)O6 peak (χ = 45
◦, 2θ = 30.03◦) matches the (102) (Y/Gd)Ba2Cu3O7−δ
(χ = 57.06◦, 2θ = 27.62◦) peaks. Furthermore considering the out-of-plane c-axis
homogeneous orientation a full heteroepitaxy between the (Y/Gd)Ba2Cu3O7−δ,
the Ba2(Y/Gd)(Nb/Ta)O6 and the SrTiO3 substrate is estabilished.
The φ scan performed on the (404) peak (χ = 45◦, 2θ = 48.518◦) is evidence
that the (Y/Gd)2O3 is rotated 45
◦ in-plane, as previously reported [141]. At a
more accurate analysis a small fraction of the (Y/Gd)2O3 particles appears to
be oriented with a minimum offset from the a and b crystalline directions of
the (Y/Gd)Ba2Cu3O7−δ, a few degrees away from a cube on cube orientation.
These broad (Y/Gd)2O3 φ peaks have been explained as a near coincidence site
108
6. Ba2YNbO6 and Gd3TaO7 simultaneous doping of YBa2Cu3O7−δ
Figure 6.2: X-ray diffraction data from φ scans of (102) YBa2Cu3O7−δ, (202)
Ba2YNbO6 and (404) (Y/Gd)2O3 from a 2.5%mol Ba2YNbO6 + 2.5%mol
Gd3TaO7 doped YBa2Cu3O7−δ thin film.
lattice matching that can possibly accommodate the large strain and structural
differences between (Y/Gd)2O3 and YBa2Cu3O7−δ [142]. However observing the
intensity ratio between the peaks related to the (Y/Gd)2O3 rotated 45
◦ and those
related to the (Y/Gd)2O3 in a near coincidence matching it is evident that only
traces of the latter can be found and that the (Y/Gd)2O3 can be considered as a
rotated 45◦ in plane.
6.3 Cross-section transmission electron microscopy
Transmission electron microscopy cross-sectional images are reported in figure
6.3. At least two different nanostructural features can be clearly observed.
The first and more obvious feature is a set of fine segmented nanorods of ≈
7 nm in diameter parallel to the (Y/Gd)Ba2Cu3O7−δ c-axis. These nanorods are
smaller than the Ba2YNbO6 nanorods which are ≈ 10 nm in diameter [5, 124],
but larger than the nanorods of Gd3TaO7 [67] and Ba2YTaO6 [68] which are ≈
5 nm in diameter. It is important to notice that while in the case of Gd3TaO7
109
6. Ba2YNbO6 and Gd3TaO7 simultaneous doping of YBa2Cu3O7−δ
Figure 6.3: a) TEM image of a YBa2Cu3O7−δ + 2.5%mol Ba2YNbO6 + 2.5%mol
Gd3TaO7 thin film cross section; b) TEM image of plate-like nanoparticleof
(Y/Gd)2O3 nucleated between two Ba2(Y/Gd)(Nb/Ta)O6 nanorods segments;
c) Selected area electron diffraction pattern of the image in (a); d) TEM imafe
of plate-like nanoparticle of (Y/Gd)2O3, inset (d) Fourier transform of the par-
ticle image. (TEM images from Prof. H Wang research group at Texas A&M
University).
110
6. Ba2YNbO6 and Gd3TaO7 simultaneous doping of YBa2Cu3O7−δ
nanorods different nanostructural features could be related to different strains and
lattice mismatch in the case of Ba2YTaO6 this is not possible because Ba2YTaO6
and Ba2YNbO6 have the identical crystalline lattice. However despite having the
same crystalline lattice the Ba2YNbO6 forms nanorods larger in diameter than
the Ba2YTaO6 hence this has to be related to different kinetics of the niobium
and the tantallum. Ba2(Y/Gd)(Nb/Ta)O6 could be described as a mixture of
Ba2YNbO6 and Ba2YTaO6 (with an additional Gd ↔ Y substitution), hence an
averaged growth kinetics of Ba2YNbO6 and Ba2YTaO6 is obtained , as might be
expected. The average nanorod segment length is ≈ 30 nm, this is shorter than
both the continuous Ba2YTaO6 that are long the entire film thickness and the
short Ba2YNbO6 that are ≈ 80-100 nm. An average nanorods spacing of ≈ 28
nm nm can be calculated from image in figure 6.3a. Calculating the matching
field for a triangular and a quadratic rods distribution (see section 4.2.3.2) the
applied magnetic field value in which the pinning potential of the analysed thin
films should be higher is between 2.64 T and 3 T, these values are larger than
those reported for the Ba2YNbO6 inclusion with the same overall doping level.
A second nanostructural feature is clearly shown in figure 6.3b: a plate-like
nanoparticle of RE2O3 (indicated as (Y/Gd)2O3) parallel to the
(Y/Gd)Ba2Cu3O7−δ ab-planes is connecting two adjoining Ba2(Y/Gd)(Nb/Ta)O6
nanorods segments. To confirm the crystalline nature of this plate-like nanopar-
ticles as RE2O3 a Fourier transform confirming the structure of a (Y/Gd)2O3
particle is shown in figure 6.3d. These plate-like particles of (Y/Gd)2O3 are in
the 25-30 nm width range and their extensions along the (Y/Gd)Ba2Cu3O7−δ
c-axis is ≈ 7 nm.
In order to confirm the phases identification obtained with the x-ray diffraction
analysis in figure 6.3c is reported a selected area electron diffraction pattern
of a large cross-sectional region. Diffraction spots that can related to SrTiO3,
(Y/Gd)Ba2Cu3O7−δ, Ba2(Y/Gd)(Nb/Ta)O6, (Y/Gd)Ba2Cu4O8 are labeled and
are consistent with the x-ray diffraction pattern reported in figure 6.1.
To conclude, the nanostructure investigation reported in this section is impor-
tant to emphasize that the nanorods segmentation is a new structural feature, and
that this is the first time it has been observed. The segmentation of the nanorods
yields to an average segment length of ≈ 30 nm. An explanation of the phe-
111
6. Ba2YNbO6 and Gd3TaO7 simultaneous doping of YBa2Cu3O7−δ
nomenon may be in kinetic terms: the tantallum ion is heavier than the niobium
thus it has a slower diffusion, this may generate a shortage in the tantallum sup-
ply as a reactant for the nanorods’ growth thus preventing these from maintaining
continuity as the film growth progresses. A new segment nucleates aligned along
the (Y/Gd)Ba2Cu3O7−δ c-axis with a preexisting segment thanks to the presence
of a nucleation enhancing strain field generated within the (Y/Gd)Ba2Cu3O7−δ
by the nanoinclusion similarly to the BaZrO3 and Ba2YNbO6 [133].
6.4 The superconducting properties: Tc, Jc(B)
and Jc(B, θ)
The complex thin film crystalline structure composition made of three different
nanoinclusion phases ((Y/Gd)2O3, Ba2(Y/Gd)(Nb/Ta)O6, (Y/Gd)Ba2Cu4O8) and
a possible RE (Gd ↔ Y) variation do not influence the transition temperature.
No suppression of the superconducting transition temperature was observed, the
Tc values measured were ≈ 89 K.
Similarly to what observed in Ba2YNbO6 doped YBa2Cu3O7−δ thin films a
small reduction of the transition temperature values from the transition of the
pure YBa2Cu3O7−δ thin films is expected [5,57,68,118,119], when this is limited to
a few Kelvin (≈ 2 K), it is evident that the secondary non superconducting phases
are stable and that the only source of Tc reduction are the induced YBa2Cu3O7−δ
lattice distortions.
6.4.1 The critical current density
In this section are compared the effects on the critical current density of Ba2YNbO6,
BaZrO3, and simultaneous Ba2YNbO6 - Gd3TaO7 doping. The critical cur-
rent density reported are all measured on thin films of ≈ 300 nm thickness.
The BaZrO3 doped YBa2Cu3O7−δ data are taken from one of the best per-
forming films in literature deposited adopting a YBa2Cu3O7−δ target covered
by yttria-stabilized zirconia on 2% of the surface area [6] while a Ba2YNbO6
doped YBa2Cu3O7−δ thin films of ≈ 350 nm thickness was specifically grown
and measured for this comparison.
112
6. Ba2YNbO6 and Gd3TaO7 simultaneous doping of YBa2Cu3O7−δ
6.4.1.1 The critical current density: Jc(B)
Figure 6.4: Critical current density variation with the applied magnetic field value
measured on a pure YBa2Cu3O7−δ, a 5%mol Ba2YNbO6 doped YBa2Cu3O7−δ,
a 2.5%mol Ba2YNbO6 - 2.5%mol Gd3TaO7 doped YBa2Cu3O7−δ and a BaZrO3
doped YBa2Cu3O7−δ thin films [6]. The field is applied parallel to the c-axis of
the YBa2Cu3O7−δ, T = 77 K. the data are on a log-linear plot (a) and a log-log
plot (b).
113
6. Ba2YNbO6 and Gd3TaO7 simultaneous doping of YBa2Cu3O7−δ
In figure 6.4 is reported Jc(B) measurements up to 4 T with the field ap-
plied parallel to the c-axis (B‖c) at 77 K for a pure YBa2Cu3O7−δ, a 5%mol
Ba2YNbO6 doped YBa2Cu3O7−δ, a 2.5%mol Ba2YNbO6 - 2.5%mol Gd3TaO7
doped YBa2Cu3O7−δ and a BaZrO3 doped YBa2Cu3O7−δ thin films. The thin
films with the simultaneous doping of Ba2YNbO6 and Gd3TaO7 show the highest
Jc of all the films studied in the entire field range analysed. An outstanding Jc
value of 1.8 MAcm−2 at 1 T and values above 1 MAcm−2 up to 2.5 T are the best
indications of the excellent potential of this new pinning landscape. Furthermore,
from the data shown in figure 6.4b it is clear that the exponential Jc decay range
(linear on double logaritmic axis) is extended to fields values above 2 T.
Another unusual feature related to this new pinning landscape is the net
change in the slope of Jc(B) that occurs as the field intensity approaches the
value of matching fields (≈ 2.6 T). This could be an indication of the effective
defusing of the low energy depinning mechanisms and of a strong pinning capacity
of the single nanorods. The Jc, in fact, start to decrease only when the fluxons
are about to saturate the pinning sites.
The high pinning capacity of the Ba2YNbO6 + Gd3TaO7 doped YBa2Cu3O7−δ
thin films can also be deduced by the α values reported in table 6.2.
Dopant α
YBa2Cu3O7−δ 0.47
Ba2YNbO6 doped YBa2Cu3O7−δ 0.35
BaZrO3 doped YBa2Cu3O7−δ 0.25
Ba2YNbO6 + Gd3TaO7 doped YBa2Cu3O7−δ 0.14
Table 6.2: α values calculated from the Jc(B) curves shown in figure 6.4.
The α value ≈ 0.14 calculated for the Ba2YNbO6 + Gd3TaO7 doped
YBa2Cu3O7−δ thin film is lower than the lowest previously reported value, an
α ≈ 0.19 reported in a specially grown BaZrO3 doped YBa2Cu3O7−δ where it
was shown that an optimisation of the deposition process could produce a mixture
of BaZrO3 randomly distributed nanoparticles and splayed columnar defects that
could synergetically reduce the depinning mechanism effectiveness [134]. In fact,
the new pinning landscape produced by the Ba2YNbO6 + Gd3TaO7 simultaneous
doping consists of segmented nanorods and plate-like nanoparticles that form a
114
6. Ba2YNbO6 and Gd3TaO7 simultaneous doping of YBa2Cu3O7−δ
synergetic pinning landscape similar to the one produce in the specially grown
BaZrO3 doped YBa2Cu3O7−δ. The two pinning landscapes show two key differ-
ences: the Ba2(Y/Gd)(Nb/Ta)O6 nanorods are smaller in diameter (≈ 7 nm)
than the BaZrO3 nanorods (≈ 10-15 nm); the Ba2(Y/Gd)(Nb/Ta)O6 nanorods
are segmented while the BaZrO3, similarly to the Ba2YNbO6, while not being
continuous have an average length of ≈ 100 nm. The segmentation could be
the key factor to explain the lower α, segmented rods could in fact gave rise to
strongly pinned staircase vortices.
6.4.1.2 The critical current density angular dependence: Jc(B, θ)
Investigating the angular dependence of Jc (figure 6.5) it is evident that the
Ba2YNbO6 + Gd3TaO7 doped YBa2Cu3O7−δ thin film shows superior pinning
properties over a large angle ranges and not only when the field is applied parallel
the YBa2Cu3O7−δ c-axis.
The Ba2YNbO6 + Gd3TaO7 doped YBa2Cu3O7−δ films show at low field val-
ues (figure 6.5a) a strong, narrow c-axis pinning peak and no variation of critical
current respect the pure YBa2Cu3O7−δ thin film is measured when the field is
applied parallel to the YBa2Cu3O7−δ ab-planes. The Jc values measured for this
new pinning landscape with the field applied parallel to the YBa2Cu3O7−δ c-
axis of 1.8 MAcm−2 is 3 times higher then the one measured for the Ba2YNbO6
doped YBa2Cu3O7−δ and 6 times higher then the one measured for the undoped
YBa2Cu3O7−δ.
The Ba2YNbO6 doped YBa2Cu3O7−δ thin film shows results that are identical
to those reported in the previous chapter: a noticeable Jc increase when the field
is applied parallel to the YBa2Cu3O7−δ c-axis together with a small Jc increase
with fields applied parallel to the YBa2Cu3O7−δ ab-planes.
Increasing the field intensity the strong, narrow c-axis pinning peaks of the
Ba2YNbO6 + Gd3TaO7 doped YBa2Cu3O7−δ films changes to a strong and broad
peak. The segmented nanorods provide strong pinning to the vortices over a
wide range of angular directions. The Jc values measured on the Ba2YNbO6
+ Gd3TaO7 doped YBa2Cu3O7−δ are higher then those measured on Ba2YNbO6
doped YBa2Cu3O7−δ indicating that the pinning of the finer, segmented
115
6. Ba2YNbO6 and Gd3TaO7 simultaneous doping of YBa2Cu3O7−δ
Figure 6.5: Critical current density variation with the direction of the applied
magnetic field measured on a pure YBa2Cu3O7−δ, a 5%mol Ba2YNbO6 doped
YBa2Cu3O7−δ, a 2.5%mol Ba2YNbO6 - 2.5%mol Gd3TaO7 doped YBa2Cu3O7−δ
and a BaZrO3 doped YBa2Cu3O7−δ thin films [6]. The field is applied parallel to
the c-axis of the YBa2Cu3O7−δ. a) Applied magnetic flux density = 1 T, T = 77
K; b) Applied magnetic flux density = 3 T, T = 77 K.
116
6. Ba2YNbO6 and Gd3TaO7 simultaneous doping of YBa2Cu3O7−δ
Ba2(Y/Gd)(Nb/Ta)O6 nanorods together with (Y/Gd)2O3 plate-like nanopar-
ticles is more effective compared to the coarser, non segmented Ba2YNbO6 rods.
6.4.1.3 A new pinning feature
A last interesting piece of information is the evolution of the angular Jc increasing
the field (figure 6.6a). If a low field the most prominent feature is the strong
c-axis pinning peak and substantially unmodified ab-planes peaks (figure 6.5a)
increasing the field values a new feature start to become visible. This new pinning
feature can be initially observed at 3 T and it appears as shoulders on the broad
c-axis peak and increasing the field value these shoulders resolves into distinct
peaks around ± 60◦. The magnitude of this two peaks increases with increasing
the field values and for field values of 5 T and above they become the dominant
pinning peaks.
This additional pinning preferential direction is a new feature that is directly
related to the new pinning landscape generated in the Ba2YNbO6 + Gd3TaO7
doped YBa2Cu3O7−δ thin films. The existence of this additional preferential pin-
ning direction can be explained using the vortex path model [143]. In this model
the combination of c-axis pinning structures (segmented Ba2(Y/Gd)(Nb/Ta)O6
nanorods) and ab-planes structures (plate-like (Y/Gd)2O3 nanoparticles and
YBa2Cu3O7−δ ab-planes) results in staircase vortices that are pinned simulta-
neously by the different structures. These vortices are subjected to the strongest
pinning at a characteristic angle determined by the distribution of defects and
defects length (figure 6.6b). Since in this type of pinning are involved both the
available species of additional pinning structure (Ba2(Y/Gd)(Nb/Ta)O6 nanorods
and (Y/Gd)2O3 nanoparticles) the hypothetical matching field would have an
higher value thus this pinning is more effective at high field values.
6.4.2 Concluding remarks to Ba2YNbO6 and Gd3TaO7 si-
multaneous doping of YBa2Cu3O7−δ
The study reported in this chapter is on the Ba2YNbO6 and Gd3TaO7 simul-
taneous doping of YBa2Cu3O7−δ. The expected mixture of Ba2YNbO6 coarse
splayed nanorods and Gd3TaO7 fine dense linear nanorods was not generated.
117
6. Ba2YNbO6 and Gd3TaO7 simultaneous doping of YBa2Cu3O7−δ
Figure 6.6: a) High field critical current density variation with the direction of the
applied magnetic field measured on a 2.5%mol Ba2YNbO6 - 2.5%mol Gd3TaO7
doped YBa2Cu3O7−δ. Applied magnetic flux density = 3 T to 6 T, T = 77 K;
b) sketch of vortices interacting simultaneously with nanorod segments along the
c-axis and intrinsic and extrinsic defects along ab-planes.
118
6. Ba2YNbO6 and Gd3TaO7 simultaneous doping of YBa2Cu3O7−δ
Combining Ba2YNbO6 and Gd3TaO7, two well studied phases, a completely new
self-assembled pinning landscape was generated instead. This pinning landscape
is not a mere mixture of the landscape generated by Ba2YNbO6 and Gd3TaO7
but presents an optimal nanorods’ architecture of fine, straight, segmented rods.
Furthermore (Y/Gd)2O3 nanoparticles were fortuitously generated because the
nanorods of Ba2(Y/Gd)(Nb/Ta)O6 formed were of different composition (poorer
in rare earth, richer in barium) than the reactants Ba2YNbO6 and Gd3TaO7 that
were added to the YBa2Cu3O7−δ pulsed laser deposition target.
The critical current densities measured at 77 K are greatly enhanced com-
pared to previously studied pinning additives. In particular Jc values above 1
MAcm−2 for fields up to 2.5 T were achieved and this is a new benchmark in the
YBa2Cu3O7−δ thin films properties. Furthermore new pinning features around
60◦ were observed for the first time proving the possibility of tuning the pinning
landscape in order to generate thin films with an arbitrary preferential pinning
direction. Applications where the magnetic field is applied obliquely to the su-
perconductor could greatly benefit from the use of this new pinning landascape.
119
Chapter 7
Conclusions and further work
The study presented in this thesis investigated the use of novel secondary non
superconducting phases to generate nanostructured YBa2Cu3O7−δ thin films pro-
duced by pulsed laser deposition from a single composite target with enhanced
flux pinning and improved performance in applied magnetic fields.
First an intensive study on the addition of a niobium based non superconduct-
ing phase was completed. At the time the pinning additive known to be effective
for the production of nanorods and the creation of a c-axis pinning peak were
BaZrO3 (a zirconium based perovskite) [57], RE3TaO7 (tantallum based pyro-
chore) [67] and BaSnO3 (a tin based perovskite) [64]. With the exception of the
RE3TaO7 all the other phases are barium based perovskites. These perovskites
form a different pinning landscape thanks to different ions kinetics and differ-
ent lattice mismatches. Thus the YBa2Cu3O7−δ thin films doped with different
perovskites have different properties.
It was known that in bulk samples adding niobium or tantallum to
YBa2Cu3O7−δ forms Ba2YNbO6 or Ba2YTaO6 [109] but at the time of that study
RE3TaO7 was the only phases reported in pulsed laser deposited YBa2Cu3O7−δ
thin films doped with tantallum while an early study reported the successfull
production of Ba2NbErO6 perovskite in ErBa2Cu3O(7−δ) [118,119].
The first step of the research work reported in this thesis is the demonstration
of the possibility to introduce Ba2YNbO6 nanorods in YBa2Cu3O7−δ together
with providing the experimental evidence of the potential of broad c-axis pinning
peaks. The broad c-axis pinning peak was related to the morphology of the novel
120
7. Conclusions and further work
Ba2YNbO6 nanorods, that was found to be different from those reported before.
The nanorods resulted to be shorter, wider and with a larger c-axis splay than
the those generated by the other known pinning additives [5].
Once discovered, the potential of Ba2YNbO6 as pinning additive, a study
of the effects of the deposition rate and substrate temperatures on Ba2YNbO6
doped YBa2Cu3O7−δ thin films was performed. The study demonstrated that fast
deposition rates induce smaller growth grains and the formation of Ba2YNbO6
nanoparticles togheter with the ordered Ba2YNbO6 nanorods. The thin films
produced adopting high deposition rates while proved to be highly effective at
high field values show a severe reduction of the intrinsic YBa2Cu3O7−δ ab-planes
pinning. This reduction is related to the large amount of discontinuity intro-
duced to the YBa2Cu3O7−δ crystall ab-planes. On the other hand Ba2YNbO6
doped YBa2Cu3O7−δ thin films deposited at low deposition rate showed that it is
possible to introduce Ba2YNbO6 and increase the pinning potential along the c-
axis without reducing the connectivity or the YBa2Cu3O7−δ crystall quality thus
without reducing the ab-plane pinning potential. In conclusion Ba2YNbO6 doped
YBa2Cu3O7−δ thin films with enanched c-axis togheter with good ab-planes pin-
ning potential were succesfully deposited. These Ba2YNbO6 doped YBa2Cu3O7−δ
thin films stand out because they show a low anisotropy of the angular depen-
dence of the critical current greatly reducing the design difficulties encountered
in many coated conductors applications.
The last part of the thesis is devoted at a new research that arises from
the desire to investigate the effects on the pinning of a synergistic combination
of morphologically different nanoinclusions. In particular the combination of the
wide, short splayed Ba2YNbO6 nanorods and the fine, long, higly linear Gd3TaO7
nanorods was investigated. This work demonstrated that in a niobium and tan-
tallum combined scenario the tantallum partecipate to the formation of a per-
ovskyte and do not form Gd3TaO7. A tantallum based perovskite Ba2YTaO6
was also recently reported in tantallum doped YBa2Cu3O7−δ [68], however also
the Ba2YTaO6 perovskite is morphologically different from the Ba2YNbO6 being
fine, long and highly linear like the previously reported Gd3TaO7 nanorods.
A completely new pinning landscape was produced by depositing thin films
from a Ba2YNbO6 + Gd3TaO7 + YBa2Cu3O7−δ PLD target. The structural
121
7. Conclusions and further work
and morphological characterization of these films showed that the these films
are mainly made of (Y/Gd)Ba2Cu3O7−δ with fine, linear, segmented nanorods of
Ba2(Y/Gd)(Nb/Ta)O6 and (Y/Gd)2O3 plate-like nanoparticles inclusions. This
new defect landscape produced in the thin films is capable of generating state-of-
the-art pinning properties. Astonishingly high critical currents are not the only
effect: for the first time possibility to achieve self segmentation of nanorods was
demostrated, and for the first time additional new pinning features around 60◦
were observed.
Owing to the very high critical currents obtained, exceeding the industry
standard BaZrO3, it is likely that the results obtained in this work will at-
tract the interest of many research groups from different area of the field. Fur-
thermore, modifying the ratio of niobium and tantallum and the ratio between
Ba2(Y/Gd)(Nb/Ta)O6 and (Y/Gd)2O3 will almost certanly demonstrate the tun-
ability of the pinning landscape, and a new era will begin in which it will be
possible to produce superconductors specifically designed for pinning in a prede-
termined direction.
A direct continuation of this work should investigate the effects of the simul-
taneous niobium tantallum doping in a simplified system. At first YBa2Cu3O7−δ
films doped only with Ba2YNbO6 and Ba2YTaO6 produced with different of
Nb:Ta ratios could be grown in order to investigate the tunability of the seg-
mentation. In a second stage (Y/Gd)2O3 could be added to the Ba2YNbO6 and
Ba2YTaO6 doping to evaluate the importance of the synergistic combination,
and find the optimal balance between plate like nanoparticles and segmented
nanorods. In a last stage Gd, as well as other rare earths, could be used to
partially substitue the Y riproducing the otpimal pinning landscape obtained in
this work but with an increased knoledge of the single constituent effects and
increased freedom on the composition of the films.
In addition a study of the pinning properties at temperatures below 77 K of
the system demonstrated in this dissertation could provide additional evidence
of the effective possibility of tuning nanoengineered YBa2Cu3O7−δ in order to
achieve optimal pinning along arbitrary directions.
122
Appendix
Publications related to the research described in this dissertation are presented
in this section. The publications order is the following:
• Ercolano G et al. “Enhanced flux pinning in YBa2Cu3O7−δ thin films
using Nb-based double perovskite additions”. Supercond. Sci. Technol.,
23:022003, 2010.
• MacManus-Driscoll J L et al. “High current, low cost YBCO conductors-
what’s next?”. Supercond. Sci. Technol., 23:034009, 2010.
• Ercolano G et al. “State-of-the-art flux pinning in YBa2Cu3O7−delta by
the creation of highly linear, segmented nanorods of Ba2(Y/Gd)(Nb/Ta)O6
together with nanoparticles of (Y/Gd)2O3 and (Y/Gd)Ba2Cu4O8”. Sub-
mitted to Supercond. Sci. Technol.
A general table of the crystallographic parameters of non-superconductive
second phase materials used as pinning phases in YBa2Cu3O7−δ films is also
reported.
123
IOP PUBLISHING SUPERCONDUCTOR SCIENCE AND TECHNOLOGY
Supercond. Sci. Technol. 23 (2010) 022003 (6pp) doi:10.1088/0953-2048/23/2/022003
RAPID COMMUNICATION
Enhanced flux pinning in YBa2Cu3O7−δ
thin films using Nb-based double
perovskite additions
G Ercolano1, S A Harrington1, H Wang2, C F Tsai2 and
J L MacManus-Driscoll1
1 Department of Materials Science and Metallurgy, University of Cambridge, Pembroke Street,
Cambridge CB2 3QZ, UK
2 Department of Electrical and Computer Engineering, Texas A&M University,
College Station, TX 77843, USA
E-mail: ge228@cam.ac.uk
Received 31 August 2009, in final form 11 December 2009
Published 13 January 2010
Online at stacks.iop.org/SUST/23/022003
Abstract
The addition of a new niobate double perovskite pinning phase to YBa2Cu3O7−δ thin films
grown by pulsed laser deposition is reported. The YBa2NbO6 phase self-assembles into stacks
of ∼10 nm second phase particles, aligned with the c-axis of the YBa2Cu3O7−δ. The
YBa2Cu3O7−δ/YBa2NbO6 composite thin films have enhanced critical current, by a factor of 2,
at 1 T (H ‖ c) over the pure YBa2Cu3O7−δ, whilst maintaining a high transition temperature.
Niobium does not substitute in the YBa2Cu3O7−δ matrix. This, together with the high stability
of the second phase formed, makes it an ideal pinning additive.
1. Introduction
Improved performance of high temperature superconductors
(HTS) in magnetic fields is a highly sought after goal to
make HTS conductors commercially viable. There are
two routes to achieve higher current carrying performance,
(a) controlled nano-engineering of superconducting materials
to enhance flux pinning [1–8] and (b) growth of thicker
highly epitaxial material in a simple and cost-effective manner
through applying novel processing routes. With industrial
applications in mind, the current focus should now be to
produce efficient, reproducible and cost-effective nanopinning
defect arrays. To achieve these end goals it is necessary to:
(1) select an ion or phase addition which produces a secondary
pinning phase which forms in a wide processing window,
and (2) to understand and hence control the second phase
assemblage. The group IV and group V ion additions, more
specifically Zr4+, Hf4, Ta5+ and, as shown in this work, Nb5+,
appear to be the best to meet the first goal [9] and we are now
at the stage of beginning to understanding the second one.
YBa2Cu3O7−δ (YBCO) thin films containing such
additions with improved Jc (up to a factor of 5) in magnetic
field was first demonstrated by the introduction of BaZrO3
(BZO) [10]. Through careful process optimization, the results
with BZO have been further substantially improved over the
last 5 years [11, 12]. More recently, with the aim of producing
minimally chemically and structural perturbative effects, rare
earth tantalates (RE3TaO7, RTO) have been studied. Compared
to BZO, these have been found to yield greater tunability of
particle assemblage, and no Tc reduction. In a short space of
time, this has led to a Jc increase by up to a factor of 10 [13].
Further studies on the RTO addition effects and on the process
optimization have been performed and will be published in the
future.
It is important to note that for the PLD technique, when
deciding on new pinning additions to YBCO the phase which
forms is not necessarily the one which is added to the
YBCO, but is the one which is the most thermodynamically
and epitaxially stable. Hence, if Zr4+ is added to YBCO,
Ba(Zr, Y)O3 forms [10]. In this paper we present the results
of addition of another large, highly charged ion, Nb5+, (see
table 1) which does not substitute for the Cu site in YBCO
and hence which has great potential to form a benign second
0953-2048/10/022003+06$30.00 © 2010 IOP Publishing Ltd Printed in the UK1
Supercond. Sci. Technol. 23 (2010) 022003 Rapid Communication
Table 1. Ionic charge and radius of critical ion additions for second
phase formation, and percentage lattice mismatch with respect to
YBCO (along ab and c). Critical pinning ion addition indicates that
the ion is not a component of the YBCO phase and that addition of
the pure element to YBCO will, for reasons of thermodynamic
stability and epitaxial stabilisation, lead to the pinning addition phase
in column 1.
Critical ion pinning
addition/ionic radius (A˚)
for 6 fold co-ordination
% Misfit to
YBCO ab
% Misfit to
YBCO c
YBa2NbO6 Nb5+/0.78 9.42 8.34
BaZrO3 Zr4+/0.86 8.11 7.03
BaSnO3 Sn4+/0.83 6.77 5.72
Sm3TaO7 Ta5+/0.78 −2.33 −7.99
Yb3TaO7 Ta5+/0.78 −4.74 −11.04
phase addition. The charge and size of the ion are important for
minimizing the level of substitution in the YBCO lattice. Zr4+
and Sn4+ (table 1) can substitute for Y3+ (whose ion radius is
1.159 A˚ for 8 fold co-ordination), whereas Ta5+ and Nb5+ are
less likely to substitute for Y3+ because of the more dissimilar
ion sizes and ion valences.
We study the addition of a double perovskite Nb-based
phase (YBa2NbO6) to YBCO. This phase is the most stable
Nb-related compound when in equilibrium with YBCO,
explaining why when we add Nb2O5 or BaNbO3 to YBCO, the
double perovskite YBa2NbO6 still forms [14–16]. We find that
nanoparticles of the phase self-assemble in a manner similar
to the tantalates [13] and BZO [17]. Addition of Nb2O5 to
bulk YBCO has been shown to enhance Jc [18]. However this
enhanced material was not well understood until later [14–16]
when segregation into YBCO + YBa2NbO6 was observed.
The stability and compatibility of the YBa2NbO6 compound
with YBCO made it a candidate for novel substrates [19],
as well as for new buffer layers for high quality HTS thin
films and electronic devices [20, 21]. Double perovskites,
of general formula A2BB′O6, have attracted much interest in
other research fields because their high flexibility in lattice
parameter and magnetic behaviour [22–25]. YBa2NbO6 has
the advantage over many of the double perovskite phases,
e.g. the well studied FeSr2MoO6 phase, in that there is less
propensity for the cations to have a mixed valence [26],
which would then cause unwanted local oxygen stoichiometry
changes in the YBCO lattice.
The use of YBa2NbO6 as a secondary phase to improve
pinning in YBCO epitaxial films on ZrO2 single crystalline
substrates has been reported with negative results [27].
The physical properties and chemical compatibility make
YBa2NbO6 an ideal candidate as a pinning addition to YBCO.
However, difficulty in obtaining adequate oxygenation of
composite bulk material [28] resulted in depressed transport
properties. This is not an obstacle to YBa2NbO6 doped
thin film growth as high surface to volume ratio and
small grain size allows full oxygenation in a rapid manner.
Here we present success in producing pinning engineered
YBa2Cu3O7−δ/YBa2NbO6 thin films with high critical currents
grown by pulsed laser deposition.
2. Experimental methods
PLD targets were made by mixing and grinding pure YBCO
powder (SCI Engineered Materials) (99.999%) with the
desired amount of YBa2NbO6 powder in an agate mortar, the
latter were produced as single phase by the solid state reaction
method, namely mixing, grinding of 99.99% Y2O3, Ba(NO3)2,
and Nb2O5 powders followed by reaction at 1450 ◦C for 24 h.
The mixed powders were pressed in the form of a cylindrical
target and sintered at 950 ◦C for 12 h in flowing O2 in a
dedicated tubular furnace. Films were grown by PLD on (001)
SrTiO3 (5 mm × 10 mm) single crystal substrate (Pi-Kem
Ltd). Deposition were performed with a Lambda Physik KrF
excimer laser (λ = 248 nm) in 30 Pa flowing O2, and the
substrate temperature was kept at 770 ◦C. A repetition rate
of 5 Hz and 4500 pulses were used for all films, which were
measured to be of 0.5 μm thickness.
The transport critical current density was measured using
a conventional four-point probe method and a ∼1 μV cm−1
criterion on photolithographically patterned bridges of 250 μm
width. The angular dependence of Jc at 77 K and 0.5 T was
measured with the applied magnetic field rotated in a plane
perpendicular to the current flow direction to an angle θ with
the c-axis of the film.
The crystallographic structure and orientation of the
films were studied using x-ray diffraction in Bragg–Brentano
geometry. Lattice parameters were refined using eva software,
using the positions of high angle peaks to minimize errors.
Cross-sectional transmission electron microscopy (TEM) and
selected area electron diffraction patterns (SADP) were also
undertaken to determine the crystal structure and sizes of the
nanoparticles formed within the YBCO matrix.
3. Results and discussion
Figure 1(a) displays the results of x-ray analysis of the
YBa2NbO6 powder produced by solid state reaction. These
results agree with the data recorded in the JCPDS or in more
recent studies [29]. No impurities or secondary phases peaks
are found. The lattice parameter calculated from the x-ray data
is 0.844 nm ± 0.001 nm.
The x-ray diffraction data from the deposited films
(figure 1(b)) shows peaks related to (00l) planes from the
YBCO and the SrTiO3 for both undoped and doped samples.
This means that the films are c-axis oriented, which is the
standard orientation for YBCO thin films on (001) SrTiO3.
Additionally, the data from the doped sample only shows the
peaks related to the (00l) planes of YBa2NbO6, indicating that
the additive phase is aligned out-of-plane with the YBCO. A
small peak at 34◦, possibly arising from the (111) planes of
YBCO is present in both the undoped and doped sample data.
No other secondary phases were present.
The (00l) orientations of both phases is expected from the
crystallographic matching of these double and triple perovskite
cells (shown schematically in figure 2(a)). Looking along the
[010]YBCO direction, 3 unit cells of the cubic YBa2NbO6 match
2 unit cells of YBCO giving a lattice mismatch strain along the
c-axis of YBCO, ((3aYBa2NbO6 −2cYBCO)/2cYBCO), of +8.34%.
2
Supercond. Sci. Technol. 23 (2010) 022003 Rapid Communication
Figure 1. X-ray diffraction data for YBa2NbO6 in the form of pure
precursor powder and when embedded within YBCO films. (a) X-ray
diffraction spectrum of YBa2NbO6 powder reacted at 1450 ◦C. (b)
X-ray diffraction data for undoped YBa2Cu3O7−δ and a 5 mol%
YBa2NbO6 doped film grown on SrTiO3; (c) φ scan of (101) SrTiO3,
(102) YBCO, (202) YBa2NbO6 from a 5 mol% YBa2NbO6 doped
film.
In the doped film all the (00l) peaks from the YBa2NbO6
are shifted to higher angles compared to the bulk values giving
a lattice parameter of 0.830 nm ± 0.001 nm compared to
the bulk value of 0.844 nm ± 0.001 nm. Hence, there is a
compressive strain of 1.7% along the c-axis, much lower than
the theoretical value of 8.34% assuming the YBCO lattice is
not distorted (table 1). The YBCO (00l) peaks are shifted to
in the opposite direction to the YBa2NbO6 peaks, i.e. to lower
angles, and the ‘c’-axis is extended to a value of 11.73 A˚ giving
Figure 2. Crystallographic matching of YBCO with YBa2NbO6. (a)
b-axis view and (b) c-axis view.
a tensile strain of 0.4%. The extended c-axis is not likely to
arise from reduced oxygen content since the sample Tc is not
reduced, as shown later. The overall strain is lower than the
theoretical value and indicates partial strain relief by formation
of misfit dislocations. Dislocation formation has previously
been reported in BZO doped YBCO [17]. Nevertheless, the
fact that the YBCO is locally partially strained in tension gives
the possibility of additional pinning from strain fields in the
vicinity of the YBCO/YBa2NbO6 interfacial regions.
To assess the in-plane texture, phi scans were recorded
as shown in figure 1(c). The (202) YBa2NbO6 peaks
reveal in-plane alignment and match both the (101) STO and
(102) YBCO peaks, confirming that the perovskite particles
have grown aligned ‘cube on cube’ with the YBCO. These
results combined with the out-of-plane x-ray data suggest
full heteroepitaxy between the YBCO, YBa2NbO6 and the
substrate.
The cube-on-cube in-plane orientation determined from
the phi scan is represented in the crystallographic model of
figure 2(b). Looking down the [001]YBCO direction, 1 unit
cell of YBa2NbO6 matches 2 unit cells of YBCO giving
an in-plane lattice mismatch (a and b average) strain of
+9.42%. The lattice mismatch of YBa2NbO6 with YBCO
is larger than other widely studied pinning phases additions
(table 1). This large mismatch should be beneficial for low field
(<1 T, 77 K) pinning due the localized strain fields that will
3
Supercond. Sci. Technol. 23 (2010) 022003 Rapid Communication
arise [30] but likely produces shorter and wider self-assembled
columnar array, contrasting with tantalate columns where the
lattice mismatch is very small and long, narrow columns are
produced [13]. They are more similar to BaZrO3 nanorods
which have similar lattice mismatch with YBCO [17, 31].
The next basic research challenges to be addressed for niobate
pinning additions to YBCO are (a) determining the optimum
level of niobate addition, (b) tuning the column sizes and
distributions.
Cross-section transmission electron micrographs of a
5 mol% YBa2NbO6 doped sample (figures 3(a) and (b)) show
nanorods aligned along the c-axis of the YBCO. The rods are
10–15 nm in diameter and their spacing is roughly 40 nm,
giving a matching field of 1.3 T. Some random particles were
also observed as indicated by the white arrows drawn along
the ab planes. The inset to figure 3(a) shows a selected area
diffraction pattern of a region around the nanoparticles. A set
of streaky diffraction dots for (004) YBa2NbO6 is observed.
Streaky dots rather than clearly distinguished spots has been
reported in similar systems previously [33] and is an indication
of a large YBCO c-plane distortion created by the YBa2NbO6
nanorod additions. d(004) estimated from the diffraction
pattern is ∼0.214 nm, indicating a ∼ 0.856 nm, broadly
consistent with the x-ray measurements.
Figure 3(b) shows a higher magnification image revealing
fringes arising from the lattice mismatch between the
YBa2NbO6 and YBCO lattices. Since a relatively slow growth
rate, 5 Hz, was used, the appearance of both self-assembled
columns and random particles are anticipated, similar to the
case of RETa3O7 additions [32]. The rods are wider (up 15 nm
as opposed to 5 nm) and shorter (∼100 nm as opposed to the
whole film thickness) than tantalate rods, most likely because
of the large mismatch strain which enhances the self-assembly
kinetics [13].
At the interface between the film and substrate, a dark
layer (about 10–20 nm thick) is observed which clearly shows
different contrast than the rest of the film. This is in agreement
with previous observations of asymmetric in-plane Jc seen in
YBCO + RETa3O7 films arising from interface disorder [34].
It is not clear if the nanorods nucleated within this darker layer
or on the substrate surface.
An atomic force micrograph of the film surface
(figure 3(c)) shows superficial particles ∼15–20 nm in
diameter uniformly distributed over the surface. These
particles are consistent with the nanorods terminating at the
sample surface. The size of the particles observed with the
AFM is larger than the rods observed with the TEM. This
could be due to the fact that we are observing the surface where
there is higher mobility and possible agglomeration of surface
particles immediately after the growth is terminated.
Despite the 5 mol% doping level and the large lattice
mismatch between the phases, only a minimal reduction of
the transition temperature is observed (figure 4(a)), with a
transition temperature of 89 K. This value indicates that the
Nb does not poison the YBCO, but remains ‘locked’ in the
second phase. In addition, the 0.4% strain level measured in the
YBCO does not lead to any significant Tc reduction. Similar
Tc reductions were also observed in ∼5 mol% BZO doped
Figure 3. Microstructures of YBCO + 5 mol%YBa2NbO6 films.
(a) and (b) TEM cross-sectional images (lower and higher
magnifications, respectively). White arrows in the direction of the
c-axis mark the positions of self-assembled nanorods along c. White
arrows perpendicular to the c-axis indicate the positions of random
nanoparticles as well as the interface between the substrate and the
film. Inset to (a) shows selected area diffraction pattern of region in
vicinity of nanoparticles. Moire fringes arising from lattice mismatch
between YBa2NbO6 and YBCO are observed in (b); (c) AFM image
showing nanoparticles at the surface.
systems [13]. However, in this case, since we are forming a
double perovskite instead of a single perovskite as for BZO,
the same mol% addition of double perovskite yields twice the
volume of inclusions as the single perovskite. Hence, the
small Tc reduction is, in fact, less than might be expected.
Jc versus magnetic field measurements up to 1 T (H ‖ c) at
4
Supercond. Sci. Technol. 23 (2010) 022003 Rapid Communication
Figure 4. Transport properties of 0.5 μm micron thick
YBCO + 5 mol%YBa2NbO6 films on STO. (a) Resistance versus
temperature for a thin film doped with 5 mol% YBa2NbO6. Inset
shows magnified region of the superconducting transition; (b) critical
current density versus applied magnetic field at 77 K, H ‖ c. Inset
shows data on log–log axis; (c) angular dependence of the critical
current density at 77 K, 0.5 T.
77 K (figure 4(b), with log–log plot shown inset) indicates that
the doped sample has better performance over the whole field
regime analysed, with Jc of the doped sample being by a factor
of ∼2 higher at 1 T.
The improved angular dependence of Jc in the YBa2NbO6
doped sample compared with pure YBCO, measured at 0.5 T
and 77 K, is shown in figure 4(c). The doped sample
shows a reduced anisotropy compared with the pure YBCO.
This lower anisotropy results from the increased Jc when
the field is aligned with the YBa2NbO6 nanorods. The ab
peak is depressed somewhat presumably by the interruption
of the intrinsic layer pinning by the highly continuous niobate
nanorods. Compared with the typical shape of the angular
dependence of pinning enhanced YBCO, for example in
YBCO + BZO and YBCO + RTO, the YBa2NbO6 nanorods
produce a very broad ‘c’-axis peak which make the particles
effective over a relatively large angular range. The large
breadth of the peak is possibly related to the large mismatch
strain which produces relatively short and wide rods (figure 3).
A significant fraction of random isotropic pinning which would
improve Jc for all angles but have little effect on the shape of
the angular dependence of Jc with magnetic field does not seem
to be present as the ab peak is lower in the doped sample than
in the pure YBCO.
In conclusion a new pinning additive, the double
perovskite YBa2NbO6, was used to produce self-assembled
non-superconducting nanorods within YBCO thin films.
Additional flux pinning was observed not only when the
magnetic field was aligned along the c-axis but also over
a relatively large angular range. Improvements in Jc of a
factor ∼2 at 1 T at 77 K were measured. The minimal
poisoning and potential for high tunability of YBa2NbO6
puts it in a similar category to RE3TaO7 pinning additions
(which has distinct advantages to BZO). The higher lattice
mismatch of YBa2NbO6 with YBCO also gives potential to
induce extra pinning from strain effects. Further optimization
of growth conditions starting from our preliminary results will
almost certainly lead to further enhancements of critical current
density.
Acknowledgments
The work was supported by the European Commission as
part of NESPA, Nano-Engineered Superconductors for Power
Application, a framework of the Marie Curie Research
Training Network, funded within the EU’s 6th framework
programme. JLM-D acknowledges support from the
Marie Curie Excellence Grant ‘NanoFen’, (MEXT-CT-2004-
014156), and HW acknowledge support from the US NSF
(DMR-0709831).
References
[1] Civale L et al 1990 Phys. Rev. Lett. 65 1164
[2] Civale L et al 1991 Phys. Rev. Lett. 67 648
[3] Konczykowski M et al 1991 Phys. Rev. 44 7167
[4] Kim D H et al 1997 IEEE Trans. Appl. Supercond. 7 1997
[5] Puig T et al 2008 Supercond. Sci. Technol. 21 034008
[6] Sparing M et al 2007 Supercond. Sci. Technol. 20 S239–46
[7] Maiorov B et al 2007 Supercond. Sci. Technol. 20 S223–9
[8] Kursumovic A et al 2009 Supercond. Sci. Technol. 22 015009
[9] Driscoll J et al 2006 US Patent Specification 0025310
[10] MacManus-Driscoll J L et al 2004 Nat. Mater. 3 439
[11] Mele P et al 2008 Supercond. Sci. Technol. 21 015019
5
Supercond. Sci. Technol. 23 (2010) 022003 Rapid Communication
[12] Mairov B et al 2009 Nat. Mater. 8 398
[13] Harrington S A et al 2009 Supercond. Sci. Technol. 22 022001
[14] Strukova G K et al 1993 Supercond. Sci. Technol. 6 589
[15] Strukova G K et al 1996 Physica C 267 67
[16] Pillai C G S and George A M 1992 J. Mater. Sci. Lett. 11 1639
[17] Goyal A et al 2005 Supercond. Sci. Technol. 18 1533
[18] Kuwabara M and Kusaka N 1988 Japan. J. Appl. Phys.
27 L1504
[19] Paulose K V et al 1992 Physica C 193 273
[20] Sathiraju S et al 2005 IEEE Trans. Appl. Supercond. 15 3009
[21] Grekhov I et al 1997 Physica C 276 18
[22] Brandle C D and Fratello V J 1990 J. Mater. Res. 5 2160
[23] Koshy J et al 1999 Bull. Mater. Sci. 22 243
[24] Aguiar J A et al 1998 Physica C 307 189
[25] Aguiar J A et al 1997 Phys. Rev. 58 2454
[26] Rager J et al 2004 J. Am. Ceram. Soc. 87 1330
[27] Jia J-H et al 1995 Mod. Phys. Lett. 9 439
[28] Vlakhov E et al 1997 J. Mater. Sci. Lett. 16 763
[29] Barnes P W et al 2006 Acta Crystallogr. B 62 384
[30] MacManus-Driscoll J L et al 2004 Appl. Phys. Lett. 84 5329
[31] Li J et al 2006 Microsc. Microanal. 12 566
[32] Harrington S et al 2010 Nanotechnology accepted
[33] Yamada K et al 2008 Appl. Phys. Lett. 92 112503
[34] Harrington S et al 2009 Appl. Phys. Lett. 95 022518
6
IOP PUBLISHING SUPERCONDUCTOR SCIENCE AND TECHNOLOGY
Supercond. Sci. Technol. 23 (2010) 034009 (5pp) doi:10.1088/0953-2048/23/3/034009
High current, low cost YBCO
conductors—what’s next?
J L MacManus-Driscoll1, S A Harrington1, J H Durrell1,
G Ercolano1, H Wang2, J H Lee2, C F Tsai2, B Maiorov3,
A Kursumovic1 and S C Wimbush1
1 Department of Materials Science and Metallurgy, University of Cambridge,
Pembroke Street, Cambridge CB2 3QZ, UK
2 Department of Electrical and Computer Engineering, Texas A&M University,
College Station, TX 77843, USA
3 MPA-STC, Los Alamos National Laboratory, Los Alamos, NM 87545, USA
E-mail: jld35@cam.ac.uk
Received 12 November 2009, in final form 1 December 2009
Published 22 February 2010
Online at stacks.iop.org/SUST/23/034009
Abstract
The Holy Grail for high temperature superconducting conductors is achieving high current
material in a simple and cost-effective way. The current status is encouraging but even after
more than twenty years of intense worldwide research, there are still many new avenues to be
explored. Innovative functional oxide materials science is central to future progress. This paper
discusses three key areas of our research focusing on new directions: highly tailored flux
pinning using the new core pinning additives R3TaO7 and RBa2NbO6 for control of
nanostructure formation; pinning using magnetic phase additives such as RFeO3 with the
potential for a magnetic contribution to the flux pinning; and the use of liquid assisted growth
enabling very high growth rates leading to thick films with no critical current degradation.
1. Introduction
Over the last decade, progress in the manufacture of high
quality, long length, high current coated conductors based on
the second generation superconducting material YBa2Cu3O7−δ
(YBCO) for a variety of transport and magnetic applications
has been impressive [1]. In particular, in just the past few
years, great progress has been made with practical flux pinning
approaches [2, 3]. There has also been some success in
minimizing the complicated series of buffer and seed layers
required between metal substrate and YBCO coating [4]. Less
attention has been paid to exploring radical new routes to
forming conductors by fast, more cost-effective routes. Of
course, the weak link problem of YBCO grains [5] means
that high angle grain boundaries need to be circumvented and
this is no mean feat. The future for coated conductors will
involve attaining the required performance at minimal cost.
Approaches involving both optimization and further scaling
up of the current processing routes as well as exploring new
horizons for achieving higher performance by more scalable
routes are required.
The areas which we believe represent new frontiers and
could ultimately lead to a step-change include highly tailored
core pinning, practical magnetic pinning, and the use of rapid
liquid assisted growth [6].
2. Experimental details
Composite deposition targets were prepared in-house from
commercial YBa2Cu3Ox powder (SCI Engineered Materials,
99.99%) together with the appropriate additives of R2O3 and
Ta2O5 (for R3TaO7 where R is a rare earth) or Y2O3, BaCO3
and Nb2O5 (for YBa2NbO6) or Y2O3 and FeO (for YFeO3).
Targets of varying compositions from 0.5 to 10 mol% were
prepared by mixing, pressing and sintering at 985 ◦C in flowing
oxygen for 12–24 h. A pure YBCO target was also prepared
under the same conditions for control purposes.
The substrates used were single crystal SrTiO3(001)
or buffered metallic tapes (∼150 nm La2Zr2O7 grown on
RABiTS Ni–W by dip coating and ex situ processing [7]).
Samples were grown in a deposition atmosphere of 30 Pa
of flowing oxygen at either 760–795 ◦C by standard pulsed
laser deposition (PLD) or at 815 ◦C by PLD incorporating
hybrid liquid phase epitaxy (HLPE). In the HLPE process, a
sub-micron thick BaO–CuO liquid layer is first deposited and
0953-2048/10/034009+05$30.00 © 2010 IOP Publishing Ltd Printed in the UK1
Supercond. Sci. Technol. 23 (2010) 034009 J L MacManus-Driscoll et al
Figure 1. Cross-sectional TEM of films of YBCO + 5 mol% Gd3TaO7 and YBCO + 5 mol% YBa2NbO6 pinning additions grown under
similar conditions, namely 770 ◦C and 790 ◦C respectively at 5 Hz on SrTiO3. Self-assembled nanocolumns of the respective pinning addition
are indicated with arrows.
growth of the YBCO occurs under this layer [8]. Subsequent to
deposition, all samples were annealed in situ in a static oxygen
pressure of 50 kPa to achieve optimum oxygenation. Film
thicknesses were typically in the range 0.5–2.5 μm.
Films were structurally characterized by x-ray diffraction
(XRD) and transmission electron microscopy (TEM). Mag-
netic properties were measured in a cryogenic vibrating sample
magnetometer (VSM) equipped with a 1 T electromagnet.
Electrical properties were measured by a standard four-
probe technique in an 8 T superconducting magnet following
photolithographic patterning and ion beam etching to form
bridge structures of typically 50–125 μm width and 1 mm
length. For angular measurements of the critical current
density, Jc, the field was applied in the maximum Lorentz
force configuration and rotated from normal to the film surface
(θ = 0◦) to parallel to it (θ = 90◦).
3. Results
3.1. Core pinning
Several new approaches to the enhancement of core pinning
in YBCO through the incorporation of novel nanostructural
inclusions have been tested and developed. The most widely
studied and successful method so far involves incorporation
of BaZrO3 nanoparticles but there are many other additions
which are also effective [9]. The current research focus
should now be on tailoring nanopinning additions to achieve
desired and controlled arrays, the optimal array (whether it
be random, correlated or a combination of both geometries)
being dependent on the application field and temperature, and
the growth method to be used. In order to nanoengineer the
materials, the kinetics and thermodynamics of phase formation
of the particular phase(s) formed when additions are made
to YBCO films needs to be understood. The mechanism of
formation of the nanostructured inclusions depends on the
lattice mismatch with the YBCO and the kinetics of self-
assembly (mobility of the constituent ions) which depend on
the growth temperature relative to the melting point of the
phase, but is also readily controllable by PLD by altering the
growth rate.
Table 1. Pinning ion additions made to YBCO films, stable phases
formed, and lattice misfits to YBCO.
Pinning ion
addition
Stable phase in
YBCO film
Misfit to
YBCO ab
Misfit to
YBCO c
Ta5+ Yb3TaO7 −4.8%a −10.9%
Ta5+ Gd3TaO7 −2.4%a −8.7%
Zr4+ BaZrO3 +8.4%b +7.5%
Nb5+ YBa2NbO6 +9.3%b +8.5%
a Rotated 45◦ to cube-on-cube.
b Cube-on-cube.
With the knowledge that Group IV and Group V ions do
not readily substitute into the YBCO lattice, but instead form
stable heteroepitaxial second phases [10], we have investigated
Ta5+ and Nb5+ additions to YBCO films [11–13] in the form
of Yb3TaO7 (a = 1.039 nm), Gd3TaO7 (a = 1.065 nm) and
YBa2NbO6 (a = 0.8436 nm), as alternatives to the common
Zr4+ ion addition in the form of BaZrO3 (a = 0.4181 nm).
The lattice mismatches of each of these phases to YBCO
(a = 0.3828 nm, b = 0.3887 nm, c = 1.1665 nm), calculated
from the bulk lattice parameters, for the most promising core
pinning additions we have developed are shown in table 1.
Figure 1 compares TEM images of nanocolumns of
Gd3TaO7 and YBa2NbO6, showing Gd3TaO7 nanocolumns
to be straighter and more continuous than YBa2NbO6
nanocolumns. If we also consider the known form of BaZrO3
nanocolumns [14], we find that lattice misfit and nanocolumn
size and perfection are interrelated, the finest and most
continuous columns arising from lower in-plane lattice misfit
to YBCO [11].
Angular critical current measurements of so far optimized
R3TaO7-added samples show the superior flux pinning
obtained with these additives compared to pure films. By
growing under conditions which produce a composite random
nanoparticle/correlated nanocolumn microstructure, Jc values
of 1.1 MA cm−2 at 1 T (figure 2(a)) and 0.7 MA cm−2 at 3 T
were obtained compared to 0.2 MA cm−2 and 0.04 MA cm−2
at the same fields for pure YBCO (all at 77 K, H ‖ c).
The YBa2NbO6-added samples generally show broader c-
axis peaks which are consistent with the more discontinuous
columns obtained (figure 2(b)).
2
Supercond. Sci. Technol. 23 (2010) 034009 J L MacManus-Driscoll et al
Figure 2. Angular Jc data at 77 K for 0.5 μm thick YBCO films with additions of Yb3TaO7 or YBa2NbO6. (a) Data at 1 T for 5 mol%
Yb3TaO7 additions compared to pure YBCO showing a large enhancement in Jc, particularly for fields applied in the c direction. (b) Data at
0.5 T comparing 1.5 mol% Yb3TaO7 and 5 mol% YBa2NbO6 additions normalized to Jc ‖ c in order to emphasize the broader c-axis peak
obtained for YBa2NbO6.
3.2. Pinning by magnetic phase additions
In practical high temperature superconductors, even with the
introduction of artificial pinning centres, the self-field Jc
is still 5–10 times lower than the theoretical limit. The
realizable advantages of magnetic pinning over ‘standard’
non-magnetic pinning originate from the additional magnetic
interaction that this provides [15]. Whereas normal materials
interact with the superconducting order parameter via the
coherence length ξ (on the order of nanometres in the high-
Tc materials), magnetic materials have a potential interaction
over length scales of the penetration depth λ which is much
larger (hundreds of nanometres). This larger interaction
volume implies that a lower density of pinning centres will
be required to achieve effective pinning and that consequently
the disruption to the superconducting phase (and its resultant
properties) will be minimized. Beyond the basic studies of
magnetic pinning [16–18], the main hurdle to overcome is
to introduce magnetic material in such a way as to avoid
significant ‘poisoning’ of the superconducting phase. This
requires forming very stable second phases that effectively
‘lock away’ the magnetic ion addition while remaining
magnetic themselves. Our exploratory studies [19] have
indicated YFeO3 to be a suitable addition. YFeO3 belongs
to the family of rare earth orthoferrites, RFeO3, which
exhibit weak ferromagnetism as a result of a slight canting
of their predominantly antiferromagnetically coupled Fe3+
moments [20]. Their structure is that of four distorted
perovskite units assembled into an orthorhombic unit cell (a =
0.5282 nm, b = 0.5596 nm, c = 0.7605 nm) [21], thereby
suggesting a high likelihood of structural compatibility with
YBCO.
There is no measurable reduction in the superconducting
critical temperature Tc for a 1 mol% YFeO3 addition [22, 23],
indicating an effective segregation of the magnetic species.
Randomly dispersed magnetic nanoparticles of sub-5 nm size
(upper right panel of figure 3) were observed leading to a 2–
3 times increase in Jc at self-field and a 10 times increase at
6 T compared to pure YBCO. Increasing the dopant amount
leads to a slight suppression of Tc and consequent reduction in
Figure 3. Angular Jc measurements at 77 K, 1 T of a 1 μm thick
film of YBCO with 1 mol% YFeO3 additions compared with a
control sample of pure YBCO. Upper right panel shows a
cross-sectional TEM image of a 1 mol% magnetically doped film,
with arrows indicating the positions of several sub-5 nm dopant
nanoparticles. Lower right panel is the room temperature (RT)
magnetic hysteresis loop (magnetic moment m versus applied field
H ) of a sample with 5 mol% YFeO3 additions.
the Jc enhancement. The particles have a good in-plane lattice
parameter match to YBCO resulting in a strain of just 0.12%
in-plane (for a 45◦ in-plane rotation of the YFeO3 unit cell
relative to the YBCO) and 2.4% out-of-plane. The coercivity
μ0 Hc is found to be rather small at around 6 mT (lower right
panel of figure 3), due to the small particle size causing a
superparamagnetic response at the temperatures of interest,
and the measured saturation field is around 0.2 T. The random
particle arrangement results in random pinning producing an
upwards shift in Jc for all field angles (figure 3).
3.3. Rapid, innovative growth of coated conductors
Rapid liquid assisted growth of YBCO + 5 mol% BaZrO3
coated conductors was achieved on metallic substrates by
HLPE. Samples of two different thicknesses were grown in
4 min (1 μm film) and 10 min (2.5 μm film). The XRD
rocking curves of the (004) peak of the La2Zr2O7 buffer
gave a rolling direction FWHM of 5.4◦ and a perpendicular
direction FWHM of 9.1◦. The XRD rocking curves for the
3
Supercond. Sci. Technol. 23 (2010) 034009 J L MacManus-Driscoll et al
Figure 4. Characterization of coated conductors of YBCO + 5 mol% BaZrO3 grown by HLPE on ∼100 nm dip-coated La2Zr2O7-buffered
RABiTS Ni-W. (a) X-ray rocking curves of YBCO layer and La2Zr2O7 buffer layer, and (b) Jc as a function of field applied parallel to c and
(inset) as a function of applied field angle θ at 1 T.
(005) peak of the YBCO layer grown on the buffer by HLPE
are improved over the buffer by approximately 0.7◦ in each
case, to 4.7◦ and 8.3◦, respectively (figure 4(a)). Transport
measurements (figure 4(b)) show that high Jc material can be
achieved on a rapidly grown, dip-coated buffer and that there is
no degradation in the Jc of the superconductor with thickness.
Hence, it is relatively easy to grow 5 μm thick material in
situ in under 20 min and still achieve self-field Jcs of around
1 MA cm−2. The several advantages of the HLPE method can
be summarized as follows: (a) the film texture is improved
over the substrate and so the method allows for the use of
much cheaper, rapidly grown substrate-buffer materials; (b) the
growth rate is several times faster than standard methods;
(c) there is no degradation in Jc with thickness. Finally, we
note that for a given Ic, a potential advantage in being able to
grow thick material having a moderate Jc over thinner material
with higher Jc is that preliminary investigations indicate that
flux creep rates are lower in thicker material [24].
4. Conclusions
There are several potential new avenues to be explored
for achieving low cost, high performance superconducting
conductors. Highly tailored core pinning, practical magnetic
pinning, and rapid liquid assisted growth are key areas. Phases
formed from Group IV and V ion additions are ideal for
achieving strong, tunable core pinning. The strain between
the particles and the YBCO matrix influences the form of their
distribution which directly impacts on the field and temperature
dependent pinning behaviour. The main challenge for practical
magnetic pinning is to form stable, non-poisoning phases.
YFeO3 is one such phase and has been shown to enhance
pinning at low (1 mol%) levels. Finally, rapid, easy growth
of thick YBCO by HLPE on ∼100 nm dip-coated La2Zr2O7-
buffered RABiTS Ni–W gives conductors with Ic > 250 A
grown in a matter of minutes.
Acknowledgments
This work has been supported by the European Commission
(Marie Curie Excellence Grant ‘NanoFen’, MEXT-CT-2004-
014156), the UK EPSRC (‘Novel materials nanoengineering
for enhanced performance of superconducting coated conduc-
tors and functional oxides’), and the EU RTN ‘NESPA’. SAH is
supported by the Mays Wild Research Fellowship of Downing
College, Cambridge. SCW is supported by The Leverhulme
Trust with supplementary funding from The Isaac Newton
Trust.
References
[1] American Superconductor, Superpower, Southwire 2009
USDOE High Temperature Superconductivity Program Peer
Review (Alexandria, VA, Aug. 2009) available at http://www.
htspeerreview.com
[2] Yamada Y et al 2005 Appl. Phys. Lett. 87 132502
[3] Kang S et al 2006 Science 311 1911
[4] Stan L, Feldmann D M, Usov I O, Holesinger T G, Maiorov B,
Civale L, DePaula R F, Dowden P C and Jia Q X 2009 IEEE
Trans. Appl. Supercond. 19 3459
[5] Dimos D, Chaudhari P and Mannhart J 1990 Phys. Rev. B
41 4038
[6] Ohnishi T, Huh J-U, Hammond R H and Jo W 2004 J. Mater.
Res. 19 977
[7] University of Cambridge and Nexans SuperConductors GmbH,
Chemiepark Knapsack, 50351 Hu¨rth, Germany 2008
European Patent Application 201373
[8] Kursumovic A, Maiorov B, Durrell J H, Wang H, Zhou H,
Stan L, Harrington S, Wimbush S, Holesinger T G and
MacManus-Driscoll J L 2009 Supercond. Sci. Technol.
22 015009
[9] Foltyn S R, Civale L, MacManus-Driscoll J L, Jia Q X,
Maiorov B, Wang H and Maley M 2007 Nat. Mater. 6 631
[10] Driscoll J L and Foltyn S R 2004 US Patent Specifiation
20060025310
[11] Harrington S A, Durrell J H, Maiorov B, Wang H,
Wimbush S C, Kursumovic A, Lee J H and
MacManus-Driscoll J L 2009 Supercond. Sci. Technol.
22 022001
[12] Harrington S A, Durrell J H, Wang H, Wimbush S C, Tsai C F
and MacManus-Driscoll J L 2009 Nanotechnology at press
[13] Ercolano G, Harrington S A, Wang H, Tsai C F and
MacManus-Driscoll J L 2009 Supercond. Sci. Technol.
at press
[14] Mele P, Matsumoto K, Horide T, Ichinose A, Mukaida M,
Yoshida Y, Horii S and Kita R 2008 Supercond. Sci.
Technol. 21 032002
[15] Blamire M G, Dinner R B, Wimbush S C and
MacManus-Driscoll J L 2009 Supercond. Sci. Technol.
22 025017
4
Supercond. Sci. Technol. 23 (2010) 034009 J L MacManus-Driscoll et al
[16] Jan D B, Coulter J Y, Hawley M E, Bulaevskii L N,
Maley M P, Jia Q X, Maranville B B, Hellman F and
Pan X Q 2003 Appl. Phys. Lett. 82 778
[17] Lange M, Van Bael M J, Moshchalkov V V and
Bruynseraede Y 2002 Appl. Phys. Lett. 81 322
[18] Yang Z, Lange M, Volodin A, Szymczak R and
Moshchalkov V V 2004 Nat. Mater. 3 793
[19] Wimbush S C, Yu R, Bali R, Durrell J H and
MacManus-Driscoll J L 2009 Physica C at press
doi:10.1016/j.physc.2009.10.117
[20] Treves D 1962 Phys. Rev. 125 1843
[21] Coppens P and Eibschu¨tz M 1965 Acta Crystallogr.
19 524 PDF 73-1345
[22] Wimbush S C, Durrell J H, Bali R, Yu R, Wang H,
Harrington S A and MacManus-Driscoll J L 2009 IEEE
Trans. Appl. Supercond. 19 3148
[23] Wimbush S C, Durrell J H, Tsai C F, Wang H, Jia Q X,
Blamire M G and MacManus-Driscoll J L 2009 Supercond.
Sci. Technol. submitted
[24] Feldmann M, Holesinger T, Maiorov B and Civale L 2009
Results presented at the USDOE High Temperature
Superconductivity Program Peer Review (Alexandria, VA,
Aug. 2009) available at http://www.htspeerreview.com
5
State-of-the-art flux pinning in YBa2Cu3O7-δ by the creation of
highly linear, segmented nanorods of Ba2(Y/Gd)(Nb/Ta)O6
together with nanoparticles of (Y/Gd)2O3 and (Y/Gd)Ba2Cu4O8
G. Ercolano1, M. Bianchetti1, S. C. Wimbush1, S. A. Harrington1, H. Wang2, J. H. 
Lee2 and J. L. MacManus-Driscoll1
1Department of Materials Science and Metallurgy, University of Cambridge, Pembroke
Street, Cambridge CB2 3QZ, UK
2Department of Electrical and Computer Engineering, Texas A&M University, College
Station, TX 77843, USA
E-mail: ge228@cam.ac.uk
Abstract. Self-assembled, segmented nanorods of c-axis aligned Ba2(Y/Gd)(Nb/Ta)O6 as well as
randomly-distributed nanoparticles of (Y/Gd)2O3 and (Y/Gd)Ba2Cu4O8 were grown into
YBa2Cu3O7-δ (YBCO) thin films by pulsed laser deposition. The complex pinning landscape
proves to be extremely effective, particularly at higher fields where the segmented vortices yield a
plateau in critical current density (Jc) with field angle around 60°. In 0.3 µm thick films, the Jc
values are higher than 1 MAcm-2 at 2.5 T (H||c-axis). Owing to the combined interactions of the
vortices with the different pinning centres, interesting new features are observed at high fields in
the angle-dependence of Jc.
1. Introduction
Even though conductors based on YBa2Cu3O7-δ (YBCO) have reached an extremely sophisticated stage of
technological development with exemplary critical current performance arising from nanoengineering of
pinning centres [1-8], since higher current conductors mean lower production costs, achieving even stronger
flux pinning by simple means is still a highly sought-after goal.
Probably the most successful method for engineering flux pinning in YBCO is by the introduction of an
epitaxial non-superconducting secondary phase, e.g. BaZrO3 [9-11]. Ta and Nb additions to YBCO have
more recently been studied based on the fact that they are highly charged (5+) ions that should perform in a
similar way to Zr and not substitute into the YBCO lattice. Tantalate additions to YBCO produce excellent
pinning performance via the formation of very fine, dense nanorods whose composition has been ascribed to
either a defective pyrochlore, RE3TaO7 (RE = rare earth, Gd being the most widely studied) [12], or to the
double perovskite, Ba2YTaO6 [13]. Niobate additions produce Ba2RENbO6 nanorods [14-17] which are
similar to BaZrO3 nanorods — wider, shorter and less linear than the tantalate rods [14] (diameter ~ 10–15
nm, splayed around the c-axis). The niobate rods remain highly effective pinning centres at low fields
aligned along the c-axis, although considering the entire field range tantalate rods are superior overall.
Since the niobate and tantalate nanorods in YBCO are of rather different morphology and hence give
different performance characteristics, it is interesting to consider whether addition of both of these second
phases of the same overall level results in an averaging effect or whether further complexity is induced in the
system, thereby yielding an entirely different pinning landscape. In fact, the situation is a mix of the former
and the latter with the final result that the nanostructure created fortuitously yields an overall superior
pinning performance, particularly at higher fields.
2. Experimental methods
Films were grown by pulsed laser ablation of a composite target of appropriate composition. Targets were
prepared by mixing and grinding pure YBa2Cu3O7-δ powder (SCI Engineered Materials 99.99%) with 2.5
mol% of Ba2YNbO6 powder and 2.5 mol% of Gd3TaO7. Ba2YNbO6 powder was produced by mixing and
grinding stoichiometric quantities of 99.99% Y2O3, Ba(NO3)2 and Nb2O5 followed by solid state reaction at
Confidential: not for distribution. Submitted to IOP Publishing for peer review 27 January 2011
1450°C for 24 h in flowing O2. Gd3TaO7 was introduced in the target by adding the desired amount of
99.99% Gd2O3 and 99.99% Ta2O5 powders to the YBa2Cu3O7-δ powder. The target mixture was pressed in a
cylindrical die (diameter = 2 cm) and sintered at 950°C for 12 h in flowing O2. Pulsed laser ablation was
performed using a Lambda Physik KrF excimer laser (λ = 248 nm). The substrates were single crystal
SrTiO3 (100) held at 780°C, and the substrate temperature was monitored using a pyrometer. 4500 laser
pulses at a 1 Hz repetition rate were used to produce 0.3 µm thick films. Films where post-annealed in situ at
520°C for 1 h in 500 mbar O2 atmosphere.
Phase and orientation analysis were performed using x-ray diffraction. Cross sectional transmission electron
microscopy (TEM) was used to determine the size, distribution and structure of the nanoparticles and
nanorods. A conventional four-point electrical measurement on photolithographically patterned bridges of 50
µm width was used to determine the critical current density, Jc, using a 1 µVcm-1 criterion. The angular
dependence of Jc was measured by rotating the applied magnetic field in a plane perpendicular to the current
direction (maximum Lorentz force configuration).
3. Results and discussion
Figure 1. Bragg-Brentano scan of a YBa2Cu3O7-δ + 2.5 mol% Gd3TaO7 + 2.5 mol% Ba2YNbO6 sample.
Figure 1 shows the x-ray diffraction pattern of a typical film. The (00l) peaks of YBCO as well as the (002)
and (004) peaks of Ba2(Y/Gd)(Nb/Ta)O6 (2θ = 21.5°, 2θ = 43.2°) are labelled. A peak at 2θ = 33.3° is
identified as the (004) peak of (Y/Gd)2O3. (Y/Gd)Ba2Cu4O8 is also present as evidenced by the (008), (0010),
and (0012) peaks (2θ = 25.9°, 2θ = 33.2°, 2θ = 41.4°).
A balanced chemical reaction consistent with the phases observed by XRD is shown below:
0.95 YBa2Cu3O7-δ + 0.025 Ba2YNbO6 + 0.025 Gd3TaO7 
 0.85 (Y/Gd)Ba2Cu3O7-δ + 0.05
Ba2(Y/Gd)(Nb/Ta)O6 + 0.0375 (Y/Gd)2O3 + 0.075 (Y/Gd)Ba2Cu4O8
Gd and Y can easily cross-substitute for one another since they have similar ionic radii. For the same reason,
Nb and Ta can also cross-substitute. This allows the double perovskite Ba2(Y/Gd)(Nb/Ta)O6 to grow within
the (Y/Gd)Ba2Cu3O7-δ lattice. Furthermore the excess Gd and the Ba deficiency produced by the Gd3TaO7 

Ba2(Y/Gd)(Nb/Ta)O6 transformation together with the Gd  Y cross-substitution is the origin of the
(Y/Gd)2O3 and (Y/Gd)Ba2Cu4O8 formation. The previously reported [12] Gd3TaO7 phase does not form here
because in the presence of Nb the Ba2(Y/Gd)(Nb/Ta)O6 phase is more stable. (Y/Gd)Ba2Cu4O8 forms
because of the barium deficiency in the Gd3TaO7 reactant. YBa2Cu4O8 has previously been observed
in coated conductors where there is barium deficiency owing to its depletion by reaction with CeO2 [18].
Also YBa2Cu4O8 has been reported in films grown by metal-organic deposition, where it is in the form of
stacking fault defects [19].
Figure 2. X-ray φ scans of the (202) Ba2(Y/Gd)(Nb/Ta)O6 and the (404) (Y/Gd)2O3 peaks from a
YBa2Cu3O7-δ + 2.5 mol% Gd3TaO7 + 2.5 mol% Ba2YNbO6 sample.
Ba2(Y/Gd)(Nb/Ta)O6 is aligned cube-on-cube with the (Y/Gd)Ba2Cu3O(7-δ) as shown in the x-ray φ scans of
figure 2. The (Y/Gd)2O3 is rotated both 45° in-plane, as previously reported [20] but also a few degrees away
(with a broad range of angles) from the cube-on-cube orientation. These broader rotations are consistent with
near coincidence site lattice matching to accommodate the very large strains and structural difference
between (Y/Gd)2O3 and YBCO [21]. 
 
Figure 3. a) TEM image of the YBa2Cu3O7-δ + 2.5 mol% Gd3TaO7 + 2.5 mol% Ba2YNbO6 sample cross
section; b) TEM image of a plate-like nanoparticle nucleated between two rod segments; c) Selected area
electron diffraction pattern of the image in (a); d) TEM image of a plate-like nanoparticles of (Y/Gd)2O3.
Fourier transform of the image of particle.
TEM cross-sections show that there are at least two different nanoscale structural features present. Firstly,
there is a set of c-axis aligned fine rods of ~7 nm in diameter, significantly smaller than the pure niobate rods
of Ba2YNbO6 which are ~15 nm in diameter [14, 15], but larger than the pure tantalate rods of Gd3TaO7
which are ~5 nm in diameter [12]. Hence, the averaged growth kinetics of pure niobate and pure tantalate
rods is obtained, as might be expected. Secondly, figure 3b shows a plate-like nanoparticle of RE2O3
nucleated along the ab-plane between two Ba2(Y/Gd)(Nb/Ta)O6 nanorods. Fig. 3c shows a selected area
diffraction pattern of a region of the matrix and inclusions. Ba2(Y/Gd)(Nb/Ta)O6, (Y/Gd)2O3 and YBa2Cu4O8
are present, consistent with the x-ray diffraction pattern of figure 1. A Fourier transform confirming the
structure of a (Y/Gd)2O3 particle of 25 nm width is shown in figure 3d inset. Other particles of (Y/Gd)2O3
were observed to be in the 25 – 30 nm width range.
A fundamentally new structural feature, which has not been observed before, is the self-segmentation of the
nanorods, yielding an average rod segment length of 30 nm. Kinetic effects of slower diffusion of the heavy
Ta5+ ion compared to the Bb5+ ion may prevent rods from maintaining continuity as the film growth
progresses. Similar to BaZrO3 and Ba2YNbO6 rods and because of the nucleation enhancing strain field
generated within the YBCO, a new segment nucleates aligned along the c-axis with a segment below it [22].
The average rod length (30 nm) is shorter than both the continuous tantalate rod length (hundreds of nm) and
the short niobate rods (~80-100 nm).
Despite the complexity of sample composition, the possibility of varying RE concentration across samples,
as well as the three different nanoinclusion phases observed, no suppression of the superconducting
transition temperature was observed. The Tc of ~ 89 K for all the samples indicates no substitution of Nb or
Ta into the YBCO matrix, as might be expected for these higher valence ions.
Figure 4. Jc versus applied magnetic field at 77 K, H || c for samples with Nb, Ta, Nb-Ta and Zr doping
compared to a pure YBCO sample. Data on a log-linear plot (a) and a log-log plot (b). The (Zr) YBCO data
is taken from [23].
Jc(H) for fields applied parallel to the c-axis for the films of this study compared with the best reported films
[23] of similar thickness (0.3µm) with other additions (Ta, Nb or Zr) is shown in figure 4. The sample with
simultaneous addition of Nb and Ta has the highest Jc across the entire field range measured with a value of
1.8 MAcm-2 at 1 T and values above 1 MAcm-2 up to 2.5 T. Furthermore, as shown in figure 4b, the linear Jc
decay range is increased from less than 1 T for pure YBCO, YBCO + 5 mol% Ba2YNbO6 ((Nb) YBCO) and
YBCO + 5 mol% Gd3TaO7 ((Ta) YBCO) to ~2 T for combined Nb+Ta ((Nb-Ta) YBCO) samples, exceeding
even that of YBCO + BaZrO3 ((Zr) YBCO) (2% of the YBCO target surface area made of YSZ [23]). The
matching field value calculated from the nanorod spacing of ~28 nm observed in the TEM image of figure 3a
is ~2.6 T, consistent with the Jc plateau up to 2.5 T of figure 4. The α values calculated for the different
additions are reported in table 1. The value of 0.14 for the (Nb-Ta) YBCO samples is lower than the best
previously reported value which was 0.19 in a specially grown YBCO + BaZrO3 sample where it was shown
that optimisation of the pinning landscape could produce both randomly distributed nanoparticles and
splayed columnar defects which together strongly reduce the depinning [24]. In fact, segmented rods were
produced which gave rise to staircase vortices.
Table 1. Calculated α values for the different additions.
Sample α
YBCO
(Nb) YBCO
(Ta) YBCO
(Nb-Ta) YBCO
(Zr) YBCO
0.47
0.35
0.26
0.14
0.25
Figure 5. Angular dependence of Jc measured at 77 K for samples with Nb, Ta, and Nb-Ta doping compared
to a pure YBCO sample at 1 T (a) and 3 T (b).
Superior properties were also found in the angular dependence of Jc of the (Nb-Ta) YBCO samples of this
study. A strong, narrow c-axis pinning peak at low fields (1 T, figure 5a) changed to a strong, broad c-axis
peak as the applied field increased (3 T, figure 5b). The broad peak is a feature already observed in (Nb)
YBCO [14]. However the Jc values of the samples prepared in this work are higher than previously observed
indicating more effective pinning by the finer, segmented rods compared to the coarser, shorter, non-
segmented rods observed in the (Nb) YBCO samples.
Figure 6. a) High field angular dependence of Jc measured at 77 K for a sample with (Nb-Ta) doping,
applied fields ranging from 3 T to 6 T; b) A sketch of vortices interacting simultaneously with rod segments
along the c-axis and intrinsic and extrinsic defects along the ab-planes.
Figure 6a shows the evolution of the angular Jc with increasing field. At low fields, the most prominent
feature is the strong c-axis pinning peak contrasted with relatively weak ab-plane peaks (figure 5a). A
shoulder initially observed on the c-axis peak resolves itself at fields beyond 3 T into distinct peaks at around
+/- 60 degrees, that grow in magnitude with increasing field to dominate all other types of pinning at fields
above 5 T. The existence of this additional preferential pinning direction is an interesting new feature arising
from the novel pinning landscape that has been established in these samples. It can be explained in terms of a
vortex path model [25] in which the combination of c-axis pinning structures (segmented
Ba2(Y/Gd)(Nb/Ta)O6 nanorods) and ab-plane pinning structures (plate-like (Y/Gd)2O3 particles) results in
staircase vortices that experience the strongest pinning at a characteristic angle determined by the particular
distribution of defect lengths present (figure 6b). This type of pinning is most effective at high fields because
it utilises both available species of pinning structure and therefore has the capacity to strongly pin a greater
overall vortex length.
4. Conclusion
A completely new self-assembled pinning landscape was produced in laser deposited YBCO films by
combining the synergistic effects from both Ta and Nb additions which together yield an optimal nanorod
architecture of fine, straight, segmented rods. (Y/Gd)2O3 nanoparticles were also generated in the matrix for
reasons of cation compensation, namely because the rods of Ba2(Y/Gd)(Nb/Ta)O6 formed were of different
average composition (poorer in rare earth, richer in barium) than the reactants of Ba2YNbO6 and Gd3TaO7
which were added to the YBCO PLD target. Greatly enhanced current densities were achieved at 77 K for
applied fields higher than 2 T compared to previously studied pinning additives. A Jc of 1 MAcm-2 was
achieved at 2.5 T and 77 K in 0.3 µm thick films which is a new performance benchmark. In addition, new
pinning features around 60° field angle were observed at > 3 T with the potential to extend the angular range
of operation of YBCO-based coated conductors.
5. Acknowledgements
Work carried out at the University of Cambridge was performed under the auspices of NESPA, Nano-
Engineered Superconductors for Power Application, a framework of the Marie Curie Research Training
Network, funded within the EUs 6th framework program. The effort at Texas A&M University was
supported by NSF-1007969. JLM-D acknowledges ERC Grant No. ERC-2009-adG 247276. SCW is funded
by The Leverhulme Trust, with supplementary funding from The Isaac Newton Trust.
6. References
[1] Haugan T, Barnes P N, Wheeler R, Meisenkothen F, Sumption M. Addition of nanoparticle dispersions
to enhance flux pinning of the YBa2Cu3O7-x superconductor. 2004, Nature, 430, 867-870.
[2] Civale L, Marwick A D, Worthington T K, Kirk M A, Thompson J R, Krusinelbaum L, Sun Y, Clem J R,
Holtzberg F. Vortex confinement by columnar defects in YBa2Cu3O7 crystals: Enhanced pinning at
high fields and temperatures. 1991, Phys. Rev. Lett., 67, 648-651.
[3] Gutierrez J, Llordes A, Gazquez J, Gibert M, Roma N, Ricart S, Pomar A, Sandiumenge F, Mestres N,
Puig T, Obradors X. Strong isotropic flux pinning in solution-derived YBa2Cu3O7-x nanocomposites
superconductor films. 2007, Nature Materials, 6, 367-373.
[4] Puig T, Gutierrez J, Pomar A, Llordes A, Gazquez J, Ricart S, Sandiumenge F, Obradors X. Vortex
pinning in chemical solution nanostructured YBCO films. 2008 Supercond. Sci. Technol., 21, 034008.
[5] Sparing M, Backen E, Freudenberg T, Huhne R, Rellinghaus B, Schultz L, Holzapfel B. Artificial
pinning centres in YBCO thin films induced by substrate decoration with gas-phase-prepared Y2O3
nanoparticles. 2007, Supercond. Sci. Technol., 20, S239-S246.
[6] Mele P, Matsumoto K, Ichinose A, Mukaida M, Yoshida Y, Horii S, Kita R. Systematic study of the
BaSnO3 insertion effect on the properties of YBa2Cu3O7-x films prepared by pulsed laser ablation.
2008, Supercond. Sci. Technol., 21, 125017.
[7] Wimbush S C, Durrell J H, Bali R, Yu R, Wang H , Harrington S A, MacManus-Driscoll J L. Practical
magnetic pinning in YBCO. 2009, IEEE Trans. Appl. Supercond., 19, 3148-3151.
[8] Reich E, Thersleff T, Huhne R, Iida K, Schultz L, Holzapfel B. Structural and pinning properties of
Y2Ba4CuMOy (M = Nb, Zr)/ YBa2Cu3O7-δ quasi-multilayers fabricated by off-axis pulsed laser
deposition. 2009, Supercond. Sci. Technol., 22, 105004.
[9] Macmanus-Driscoll J L, Foltyn S R, Jia Q X, Wang H, Serquis A, Civale L, Maiorov B, Hawley M E,
Maley M P, Peterson D E. Strongly enhanced current densities in superconducting coated conductors
of YBa2Cu3O7-x + BaZrO3. 2004, Nature Materials, 3, 439-443.
[10] Mele P, Matsumoto K, Horide T, Ichinose A, Mukaida M, Yoshida Y, Horii S, Kita R. Incorporation of
double artificial pinning centers in YBa2Cu3O7-δ films. 2008, Supercond. Sci. Technol., 21, 015019.
[11] Mele P, Matsumoto K, Horide T, Ichinose A, Mukaida M, Yoshida Y, Horii S, Kita R. Ultra-high flux
pinning properties of BaMO3-doped YBa2Cu3O7-x thin films (M = Zr, Sn). 2008, Supercond. Sci.
Technol., 21, 032002.
[12] Harrington S A, Durrell J H, Maiorov B, Wang H, Wimbush S C, Kursumovic A, Lee J H, MacManus-
Driscoll J L. Self-assembled, rare earth tantalate pyrochlore nanoparticles for superior flux pinning in
YBa2Cu3O7-δ films. 2009, Supercond. Sci. Technol., 22, 022001.
[13] Wee S H, Goyal A, Specht E D, Cantoni C, Zuev Y L, Selvamanickam V, Cook S. Enhanced flux
pinning and critical current density via incorporation of self-assembled rare-earth barium tantalate
nanocolumns within YBa2Cu3O7-δ films. 2010, Phys. Rev. B, 81, 14050.
[14] Ercolano G, Harrington S A, Wang H, Tsai C F, MacManus-Driscoll J L. Enhanced flux pinning in
YBa2Cu3O7-δ thin films using Nb-based double perovskite additions. 2010, Supercond. Sci. Technol.,
23, 022003.
[15] Wee S H, Goyal A, Zuev Y L, Cantoni C, Selvamanickam V, Specht E D. Formation of Self-
Assembled, Double-Perovskite, Ba2YNbO6 Nanocolumns and Their Contribution to Flux-Pinning and
Jc in Nb-Doped YBa2Cu3O7-δ Films. 2010, Appl. Phys. Express, 3, 023101.
[16] Kai H, Horii S, Ichinose A, Kita R, Matsumoto K, Yoshida Y, Fujiyoshi T, Teranishi R, Mori N,
Mukaida M. The effects of growth temperature on c-axis-correlated pinning centers in PLD
ErBa2Cu3O7-δ films with Ba(Er0.5Nb0.5)O3. 2010, Supercond. Sci. Technol., 23, 025017.
[17] Horii S, Yamada K, Kai H, Ichinose A, Mukaida M, Teranishi R, Kita R, Matsumoto K, Yoshida Y,
Shimoyama J, Kishio K. Introduction of c-axis-correlated 1D pinning centers and vortex Bose glass in
BaNbO doped ErBa2Cu3Oy films. 2007, Supercond. Sci. Technol., 20, 1115-1119.
[18] Holesinger T G, Foltyn S R, Arendt P N, Kung H, Jia Q X, Dickerson R M, Dowden P C, DePaula R F,
Groves J R, Coulter J Y. The microstructure of continuosly processed YBa2Cu3Oy coated conductors
with underlying CeO2 and ion 'beam' assisted yttria stabilized zirconia buffer layers. 2000, Jour. of
Mat. Res., 15, 1110-1119.
[19] Specht E D, Goyal A, Li J, Martin P M, Li X, Rupich M W. Stacking faults in YBa2Cu3O7-x:
Measurement using x-ray diffraction and effects on critical current. 2006, Appl. Phys. Lett., 89,
162510.
[20] Gapud A A, Kumar D, Viswanathan S K, Cantoni C, Varela M, Abiade J, Pennycook S J, Christen D K.
Enhancement of flux pinning in YBa2Cu3O7-δ thin films embedded with epitaxially grown Y2O3
nanostructures using a multi-layering process. 2005, Supercond. Sci. Technol., 18, 1502-1505.
[21] Garrison S M, Newman N, Cole B F, Char K, Barton R W. Observation of 2 in-plane epitaxial states in
YBa2Cu3O7-δ films on yttria-stabilized ZrO2. 1991, Appl. Phys. Lett., 58, 2168-2170.
[22] Harrington S A, Durrell J H, Wang H, Wimbush S C, Tsai C F, MacManus-Driscoll J L. Understanding
nanoparticle self-assembly for a strong improvement in functionality in thin film nanocomposites.
2010, Nanotechnology, 21, 095604.
[23] Mele P, Matsumoto K, Horide T, Ichinose A, Mukaida M, Yoshida Y, Horii S. Enhanced high-field
performance in PLD films fabricated by ablation of YSZ-added YBa2Cu3O7-x target. 2007, Supercond.
Sci. Technol., 20, 244-250.
[24] Maiorov B, Baily S A, Zhou H, Ugurlu O, Kennison J A, Dowden P C, Holesinger T G, Foltyn S R,
Civale L. Synergetic combination of different types of defect to optimize pinning landscape using
BaZrO3-doped YBa2Cu3O7. 2009, Nature Materials, 8, 398-404.
[25] Long N J. Model for the angular dependence of critical currents in technical superconductors. 2008,
Supercond. Sci. Technol., 21, 025007.
Pinning Stable Lattice Lattice Lattice Ref.
ion Phase Parameter misfit misfit
addition in YBa2Cu3O7−δ [nm] to abY BCO to cY BCO
Nb+4 Ba2YNbO6 0.422 +9.4%
a +8.4% [5,74]
[123,124]
Ta+4 Ba2YTaO6 0.422 +9.4%
a +8.4% [68]
Zr+4 BaZrO3 0.419 +8.8%
a +7.6% [57–63]
Sn+4 BaSnO3 0.411 +6.8%
a +5.6% [64–66]
Hf+4 BaHfO3 0.417 +8.3%
a +7.1% [144]
Ir+4 BaIrO3 0.410 +6.5%
a +5.3% [145]
Ti+4 BaTiO3 0.402 +4.4%
a +3.2% [146,147]
Y+3 Y2O3 1.060 -2.8%
b -9.2% [78,142]
[141,148]
Ta+5 Gd3TaO7 1.065 -2.4%
b -8.7% [67]
Ta+5 Yb3TaO7 1.039 -4.8%
b -10.9% [67]
Table 1: Crystallographic parameters of non-superconductive second phase ma-
terials used as pinning phases in YBa2Cu3O7−δ films (a cube-on-cube, b rotated
45◦ to cube-on-cube).
145
References
[1] Tinkham M. Introduction to Superconductivity 2nd ed. McGraw-Hill, New
York, 1996. xi, 6
[2] Cava R J et al. Structural anomalies, oxygen ordering and superconductiv-
ity in oxygen deficient Ba2YCu3Ox. Physica C, 165:419–433, 1990. xi, 12,
13
[3] Singh R K et al. Pulsed laser deposition and charcterization on high-Tc
Y Ba2Cu3O7−x superconducting thin films. Mat. Sci. and Eng. RR, 22:113–
185, 1998. xi, 23, 25
[4] Paulose K V et al. Ba2YNbO6: synthesis, properties and compatibility with
YBa2Cu3O7−δ. Physica C, 193:273–276, 1992. xii, 41, 42, 43, 44, 46
[5] Ercolano G et al. Enhanced flux pinning in YBa2Cu3O7−δ thin films using
Nb-based double perovskite additions. Supercond. Sci. Technol., 23:022003,
2010. xii, 43, 44, 46, 48, 49, 104, 109, 112, 121, 145
[6] Mele P et al. Enhanced high-field performance in PLD films fabricated by
ablation of YSZ-added YBa2Cu3O7−x target. Supercond. Sci. Technol., 20,
2007. xvii, 89, 112, 113, 116
[7] Onnes H K. Investigation into the properties of sustances at low temper-
atures, which have led, amongst other things to the preparation of liquid
helium. Nobel lectures, pages 306–336, 1913. 1
[8] Onnes H K. Liquid helium. Comptes Rendus Hebdomadaires Des Seances
de l’Academie Des Sciences, 147:421–424, 1908. 1
146
REFERENCES
[9] Onnes H K. Further experiments with liquid helium D - On the change of
the electrical resistance of pure metals at very low temperatures, etc V The
disappearance of the resistance of mercury. Proceedings of the Koninklijke
Akademie Van Wetenshappen te Amsterdam, 14:113–115, 1911. 1
[10] Ochsenfeld R Meissner W. Ein neuer Effekt bei Eintritt der
Supraleitfhigkeit. Naturwissenschaften, 21:787–788, 1933. 2
[11] London H London F. The electromagnetic equations of the supraconductor.
Proc. Roy. Soc., 149:71–88, 1935. 2
[12] Landau L D Ginzburg V L. To the Theory of Superconductivity. Zh. Eksp.
Teor. Fiz., 20:1064, 1950. 4
[13] Abrikosov A A. On the Magnetic Properties of Superconductors of the
Second Group. Soviet Phys. jetp-USSR, 5:1174–1183, 1957. 5
[14] Abrikosov A A. Type II Superconductors and the vortex lattice. Nobel
Lecture, 2003. 5
[15] Kleiner W H et al. Bulk solution of ginzburg-landau equations for type II
superconductors: Upper critical field region. Phys. Rev. A, 133:1226, 1964.
5
[16] Kramer L. Thermodynamical behaviour of type-II superconductors with
small k near the lower critical field. Phys. Lett., 23:619, 1966. 5
[17] Schrieffer J R Bardeen J, Cooper L N. Theory of Superconductivity. Phys-
ical Review, 108:1175–1204, 1957. 7
[18] Cooper L N. Bound Electron Pairs in a Degenerate Fermi Gas. Physical
Review, 104:1189–1190, 1956. 7
[19] Mu¨ller K A Bednorz J G. Possible highTc superconductivity in the Ba-La-
Cu-O system. Zeitschrift fr Physik B Condensed Matter, 64:189–193, 1986.
8
147
REFERENCES
[20] Wu K M et al. Superconductivity at 93 K in a new mixed-phase Yb-Ba-
Cu-O compound system at ambient pressure. Phys. Rev. Lett., 58:908–910,
1987. 8
[21] Maeda H et al. A New High-Tc Oxide Superconductor without a Rare Earth
Element. Jpn. J. Appl. Phys, 27:209–210, 1988. 8
[22] Hermann A M Sheng Z Z. Bulk Superconductivity at 120 K in the Tl-
Ca/Ba-Cu-O system. Nature, 332:138–139, 1988. 8
[23] Nagamatsu J et al. Superconductivity at 39 K in magnesium diboride.
Nature, 410:63–64, 2001. 8
[24] Larbalestier D C et al. Strongly linked current flow in polycrystalline forms
of the superconductor MgB2. Nature, 410:186–189, 2001. 8
[25] Buzea C and Yamashita T. Review of the superconducting properties of
MgB2. Supercond. Sci. Technol., 14:R115–R146, 2001. 8
[26] Eom C B et al. High critical current density and enhanced irreversibility
field in superconducting MgB2 thin films. Nature, 411:558–560, 2001. 8
[27] Jin S et al. High critical currents in iron-clad superconducting MgB2 wires.
Nature, 411:553–565, 2001. 8
[28] Kamihara Y et al. Iron-Based Layered Superconductor: LaOFeP. J. Am.
Chem. Soc, 128:10012–10013, 2006. 8
[29] Takahashi H et al. Superconductivity at 43 K in an iron-based layered
compound LaO1−xFxFeAs. Nature, 453:376–378, 2008. 8
[30] Kamihara Y et al. Iron-based layered superconductor La[O1−xFx]FeAs
(x=0.05-0.12) with Tc=26 K. J. Ame. Chem. Soc., 130:3296, 2008. 8
[31] Rotter M et al. Superconductivity at 38 K in the iron arsenide
(Ba(1−x)K(x))Fe(2)As(2). Phys. Rev. Lett., 101:107006, 2008. 8
[32] Chen X H et al. Superconductivity at 43 K in SmFeAsO1−xFx. Nature,
453:761–762, 2008. 8
148
REFERENCES
[33] Wu G et al. Superconductivity at 56 K in samarium-doped SrFeAsF. J. of
Phys:Cond. mat., 21, 2009. 8
[34] Dinger T R et al. Direct Observation of Electronic Anisotropy in Single-
Crystal YBa2Cu3O7−x. Phys. Rev. Lett., 58:2687–2690, 1987. 10
[35] Chaudhari P et al. Critical Currents Measurements in Epitaxial Films of
YBa2Cu3O7−x Compound. Phys. Rev. Lett., 58:2684–2686, 1987. 10
[36] Larbalestier D C et al. Experiments concerning the connective nature of
superconductivity in YBa2Cu3O7. J. Appl. Phys., 62:3308–3313, 1987. 10
[37] Dimos D et al. Orientation Dependence of Grain-Boundary Critical Cur-
rents in YBa2Cu3O7−δ Bicrystals. Phys. Rev. Lett., 61:219–222, 1988. 10
[38] Mayer B et al. Superconducting transport properties of Bi2Sr2CaCu2O8+x
bicrystal grain boundery junctions. Appl. Phys. Lett., 63:996–998, 1993. 10
[39] Heine K et al. High-field critical current densities inBi2Sr2Ca1Cu2O8+x/Ag
wires. Appl. Phys. Lett., 55:2441–2443, 1989. 10
[40] Yamada Y et al. Critical Current Density of Wire Type Y-Ba-Cu Oxide
Superconductor. Jpn. J. Appl. Phys., 26:865–866, 1987. 10
[41] Flu¨kiger R et al. Metallurgy and critical current in YBa2Cu3O7 wires.
Physica C, 153:1574–1579, 1988. 10
[42] Wu X D et al. High-current YBa2Cu3O7−δ thick films on flexible nickel
substrates with textured buffer layers. Appl. Phys. Lett., 65:1961–1963,
1994. 10
[43] Goyal A et al. Epitaxial superconductors on rolling-assisted biaxially-
textured substrates (RABiTS): A route towards high critical current density
wire. Appl. Superconductivity, 4:403–427, 1996. 10
[44] Larbalestier D et al. High-Tc superconducting materials for electric power
applications. Nature, 414:368–377, 2001. 11
149
REFERENCES
[45] Cava R J et al. Bulk Superconductivity at 91 K in Single-Phase Oxygen-
Deficient Perovskite Ba2YCu3O9−δ. Phys. Rev. Lett., 58:1676–1679, 1987.
12
[46] Beno M A et al. Structure of the single-phase high-temperature supercon-
ductor YBa2Cu3O7−δ. Appl. Phys. Lett., 51:57–59, 1987. 12
[47] David W I F et al. Structure and crystal chemistry of the high-Tc super-
conductor YBa2Cu3O7−x. Nature, 327:310–312, 1987. 12
[48] Tietz L A and Carter C B. Special grain boundaries in YBa2Cu3O7−x thin
films. Physica C Supercond., 182:241251, 1991. 12
[49] Linfeng Mei and Siu-Wai Chan. Enthalpy and entropy of twin boundaries
in superconducting YBa2Cu3O7−x. J. Appl. Phys., 98:033908, 2005. 12
[50] Jorgenson J D et al. Structural and superconducting properties of or-
thorhombic and tetragonal YBa2Cu3O7−x: The effect of oxygen stoichiom-
etry and ordering on superconductivity. Phys. Rev. B, 36:5731–5734, 1987.
12
[51] Jorgenson J D et al. Oxygen ordering and the orthorhombic-to-tetragonal
phase transition in YBa2Cu3O7−x. Phys. Rev. B, 36:36083616, 1987. 12
[52] Blatter G et al. Vortices in high-temperature superconductors. Rev. Mod.
Phys., 66:1125–1338, 1994. 14
[53] Evetts J E. Flux Pinning by Phase Boundaries in Type-II Superconductors.
Phys. Rev. B, 2:95–96, 1970. 16
[54] Coote R I et al. Flux Line Pinning by Large Normal Particles in Type-II
Superconductors. Can. J. Phys., 50:421–427, 1972. 16
[55] Campbell A M and Evetts J E. Flux vortices and transport currents in
type II superconductors. Adv. Phys., 21:199–428, 1972. 16
[56] Larkin A I and Ovchinnikov Y N. Pinning in type II superconductors. J.
Low Temp. Phys., 34:409–428, 1979. 16
150
REFERENCES
[57] MacManus-Driscoll J L et al. Strongly enhanched current densities in su-
perconducting coated conductors of Y Ba2Cu3O7−x+BaZrO3. Nature Ma-
terials, 3:439–443, 2004. 17, 39, 50, 59, 89, 112, 120, 145
[58] Foltyn S R et al. Materials science challenges for high-temperature super-
conducting wire. Nature Materials, 6:631, 2007. 18, 145
[59] Galluzzi V et al. YBa2Cu3O7−δ films with BaZrO3 inclusions for strong-
pinning in superconducting films on single crystal substrate. IEEE Trans.
Appl. Supercond., 17:3628, 2007. 18, 145
[60] Peurla M et al. YBCO films prepared by PLD using nanocrystalline targets
doped with BaZrO3 or Y211. IEEE Trans. Appl. Supercond., 15:3050–3053,
2005. 18, 145
[61] Ichinose A et al. Microstructures and critical current densities of YBCO
films containing structure-controlled BaZrO3 nanorods. Supercond. Sci.
Technol., 20:1144, 2007. 18, 145
[62] Ito M et al. Microstructure of ErBa2Cu3O7−δ films with BaZrO3 dispersion
pinning centers for high Jc applications. Physica C, 426:1415, 2005. 18, 145
[63] Feldmann D M et al. 1000 Acm−1 in a 2 µm thick YBa2Cu3O7−x film with
BaZrO3 and Y2O3 additions. Supercond. Sci. Technol., 23:115016, 2010.
18, 145
[64] Varanasi C V et al. Enhancement and angular dependence of transport
critical current density in pulsed laser deposited Y Ba2Cu3O7−x +BaSnO3
films in applied magnetic fields. Jour. of Appl. Phys., 102, 2007. 18, 39,
50, 120, 145
[65] Mele P et al. Systematic study of the BaSnO3 insertion effect on the prop-
erties of YBa2Cu3O7−x films prepared by pulsed laser ablation. Supercond.
Sci. Technol., 21, 2008. 18, 39, 89, 145
[66] Mele P et al. Ultra-high flux pinning properties of BaMO3-doped
YBa2Cu3O7−x thin films (M = Zr, Sn). Supercond. Sci. Technol., 21:032002,
2008. 18, 39, 145
151
REFERENCES
[67] Harrington S A et al. Self-assembled, rare earth tantalate pyrochlore
nanoparticles for superior flux pinning in Y Ba2Cu3O7−δ films. Supercond.
Sci. Technol., 22, 2009. 18, 39, 59, 105, 108, 109, 120, 145
[68] Hun Wee S et al. Enhanced flux pinning and critical current density via
incorporation of self-assembled rare-earth barium tantalate nanocolumns
within Y Ba2Cu3O7−δ films. Phys. Rev. B, 81, 2010. 18, 39, 59, 108, 109,
112, 121, 145
[69] Gutierrez J et al. Strong isotropic flux pinning in solution-derived
Y Ba2Cu3O7−x nanocomposite superconductor films. Nature Materials,
6:367–373, 2007. 18, 63
[70] Hougan T et al. Addition of nanoparticle dispersions to enhance flux pin-
ning of the Y Ba2Cu3O7−x superconductor. Nature, 430:367–373, 2004. 18
[71] Campbell T A et al. Flux pinnig effect of Y2O3 nanoparticulate dispersion
in multilayered YBCO thin films. Physica C: Supercond., 423:1–8, 2005.
18, 19
[72] Kang S et al. Strong enhancement of flux pinning in YBa2Cu3O7−δ multi-
layers with columnar defects comprised of self-assembled BaZrO3 nanodots.
Supercond. Sci. Technol., 20, 2007. 18
[73] Kang S et al. Flux pinning characteristic as a function of density of colum-
nar defects comprised of self-assembled nanodots and nanorods in epitaxial
YBa2Cu3O7−δ films for coated conductor applications. Physica C, 457:41–
46, 2007. 18
[74] Reich E et al. Structural and pinning properties of Y2Ba4CuMOy (M =
Nb, Zr) / YBa2Cu3O7−δ quasi-multilayers fabricated by off-axis pulsed laser
deposition. Supercond. Sci. Technol., 22, 2009. 19, 43, 89, 145
[75] Matsumoto K et al. Enhancement of critical current density of YBCO films
by introduction of artificial pinning centers due to the distributed nano-
scaled Y2O3 islands on substrates. Physica C Supercond., 412:1267–1271,
2004. 19
152
REFERENCES
[76] Matsumoto K et al. Effects of artificial pinning centers on vortex pinning
in high-temperature superconducting films. Physica C Supercond., 426-
431:1091–1095, 2005. 19
[77] Mele P et al. Critical current enhancement in PLD YBa2Cu3O7−x films
using artificial pinning centers. Physica C Supercond., 445-448:648–651,
2006. 19
[78] Mele P et al. Tuning of the critical current in YBa2Cu3O7−x thin films by
controlling the size and density of Y2O3 nanoislands on annealed SrTiO3
substrates. Supercond. Sci. Technol., 19:44, 2006. 19, 145
[79] Li A H et al. Effect of substrate surface modification using Ag nano-dots
on the improvement of Jc and microstructures in YBa2Cu3O7 thin films
grown on LaAlO3 (100) by pulsed laser deposition. Journ. Electroceram.,
16:605–609, 2006. 19
[80] Dang V S et al. Combination of Ag Substrate Decoration with Introduc-
tion of BaZrO3 Nano-Inclusions for Enhancing Critical Current Density of
YBa2Cu3O7 Films. Journ. Electroceram., 24:505–509, 2011. 19
[81] Aytug T et al. Enhancement of flux pinning and critical currents in
YBa2Cu3O7−δ films by nanoscale iridium pretreatment of substrate sur-
faces. Jour. of Appl. Phys., 98, 2005. 19
[82] Aytug T et al. Analysis of flux pinning in YBa2Cu3O7−δ films by
nanoparticle-modified substrate surfaces. Phys. Rev. B, 74:184505, 2006.
19
[83] Aytug T et al. Substrate Surface Decoration With CeO2 Nanoparticles:
An Effective Method for Improving Flux Pinning in YBa2Cu3O7−δ Films.
IEEE Trans. Appl. Supercond., 17:3720–3723, 2007. 19
[84] Crisan A et al. Sputtered nanodots: A costless method for inducing effective
pinning centers in superconducting thin films. Appl. Phys. Lett., 79:4547–
4549, 2001. 19
153
REFERENCES
[85] Sparing M et al. Artificial pinning centres in YBCO thin films induced by
substrate decoration with gas-phase-prepared Y2O3 nanoparticles. Super-
cond. Sci. Technol., 20:S239–S246, 2007. 19
[86] Cui X M et al. Enhancement of critical current density of YBa2Cu3O7−δ
thin films by nanoscale CeO2 pretreatment of substrate surfaces. Physica
C Supercond., 466:1–4, 2007. 19
[87] MacManus-Driscoll J L et al. Guidelines for optimizing random and cor-
related pinning in rare-earth based superconducting films. Supercond. Sci.
Technol., 19:S55–S59, 2006. 19
[88] Civale L et al. Vortex confinement by columnar defects in YBa2Cu3O7
crystals: Enhanced pinning at high fields and temperatures. Phys. Rev.
Lett., 67, 1991. 19
[89] Weber H W et al. Critical currents in neutron-irradiated YBCO and BIS-
CCO single-crystals. Supercond. Sci. Technol., 4:S103–S105, 1991. 19
[90] Jirong J et al. Irradiation effect of slow neutrons on melt-textured
YBa2Cu3O7−δ. Phys. Stat. Sol., 125:289–294, 1991. 19
[91] Wisniewski A et al. Comparison of neutron-irradiation effects in the 90-K
and 60-K phases of YBCO ceramics. Physica C, 197:365–370, 1992. 19
[92] Kato T et al. Radiation effect of YBa2Cu3O7−y irradiated by γ-Rays and
14 MeV neutrons. Jpn. J. Appl. Phys., 27:L2097–L2099, 1988. 19
[93] Vasek P et al. Gamma-irradiation of YBa2Cu3O7−x ceramics. Solid State
Com., 69:23–25, 1989. 19
[94] Kim D H et al. Effect of Heavy-Ion Irradiation on Transport Properties of
Y Ba2Cu3Ox Films. Trans. on appl. Supercond., 7, 1997. 20
[95] Nakashima K et al. Effect of ion-irradiation and annealing on supercon-
ductive property of PLD prepared YBCO tapes. Physica C Supercond.,
463:665–668, 2007. 20
154
REFERENCES
[96] Eisterer M et al. Neutron irradiation of coated conductors. Supercond. Sci.
Technol., 23:014009, 2010. 20
[97] Chudy M et al. Asymmetric angular dependence of J(c) in coated con-
ductors prior to and after fast neutron irradiation. Physica C Supercond.,
470:1300–1303, 2010. 20
[98] Chudy M et al. Characterization of Commercial YBCO Coated Conductors
After Neutron Irradiation. IEEE Trans. Appl. Supercond., 21:3162–3165,
2011. 20
[99] Fuger R et al. Influence of neutron irradiation on high temperature super-
conducting coated conductors. Physica C Supercond., 468:1647–1651, 2008.
20
[100] Yamada H et al. Flux pinning centres correlated along the c-axis in PLD-
YBCO films. Supercond. Sci. Technol., 17:58, 2004. 22
[101] Civale L et al. Influence of crystalline texture on vortex pinning near the
ab-plane in YBa2Cu3O7 thin films and coated conductors. Physica C Su-
percond., 412:976–982, 2004. 22
[102] Takahashi K et al. Investigation of thick PLD-GdBCO and ZrO2 doped
GdBCO coated conductors with high critical current on PLD-CeO2 capped
IBAD-GZO substrate tapes. Supercond. Sci. Technol., 19:924–929, 2006.
22
[103] Traito K et al. Magnetic field dependence of the critical current and the
flux pinning mechanism in YBa2Cu3O6+x films doped with BaZrO3. Phys.
Rev. B, 73:224522, 2006. 22
[104] Horide T et al. Flux pinning properties of YBCO thin films deposited on
SrTiO3(100) and MgO(100) substrates. Physica C Supercond., 412:1291–
1295, 2004. 22
[105] Douglas B C and Graham K H. Pulsed Laser Deposition of Thin Films.
Jhon Wiley & Sons, New York, 1994. 22, 51
155
REFERENCES
[106] Kwok H S et al. Laser induced deposition of YBa2Cu3O7−x thin films. In
High-Temperature Superconductors, pages 273–277, 1988. 22
[107] Horcas I et al. WSXM: A software for scanning probe microscopy and a
tool for nanotechnology. Rev. Sci. Instrum., 78:013705, 2007. 35
[108] Kuwabara M et al. Microstructure and Superconducting Properties in
YBa2Cu3O7−x ceramics Doped with Nb2O5 and WO3. Jap. Jour. of Appl.
Phys., 27:L1504–L1506, 1988. 40
[109] Greaves C and Slater PR. Nb and Ta Substitutions in YBa2Cu3O7−x
and related phases − Structural Characterization of La1.1Ba1.9Cu2.1M0.9O8
(M=Nb, Ta). Physica C, 161:245–251, 1989. 40, 108, 120
[110] Strukova G K et al. Effect of Nb doping on properties of Y-Ba-Cu ceramics.
Supercond. Sci. Technol., 6:589–592, 1993. 40
[111] Strukova G K et al. A cubic phase in ceramic Y-Ba-Cu-O doped with Nb.
Physica C, 267:67–73, 1996. 40
[112] Babu N H et al. New chemically stable, nano-size artificial flux pinning cen-
tres in (RE)-Ba-Cu-O superconductors. Supercond. Sci. Technol., 16:L44–
L45, 2003. 41
[113] Babu N H et al. Artificial flux pinning centers in large, single-grain (RE)-
Ba-Cu-O superconductors. Appl. Phys. Lett., 83:4806–4808, 2003. 41
[114] Yeoh W K et al. Improved Flux Pinning in Y-Ba-Cu-O Superconductors
Containing Niobium Oxide. IEEE Trans. Appl. Supercond., 19:2970–2973,
2009. 41
[115] Satiraju S et al. Studies on Ba2YNbO6 buffer layers for subsequent
YBa2Cu3O7−x film growth. IEEE Trans. Appl. Supercond., 15:3009–3012,
2005. 41
[116] Grekhov I et al. A new buffer layer for high-quality HTSC ultrathin film
fabrication. Physica C, 276:18–24, 1997. 41
156
REFERENCES
[117] Jia J H et al. Effect of the Second Phase on Critical Current Density in
Nb-Added YBCO Films. Mod. Phys. Lett., 9:439–443, 1995. 41
[118] Hori S et al. Introduction of c-axis-correlated 1D pinning centers and vortex
Bose glass in BaNbO-doped ErBa2Cu3Oy films. Supercond. Sci. Technol.,
20:1115–1119, 2007. 41, 59, 89, 104, 112, 120
[119] Kai H et al. Superconducting properties and microstructure of PLD-
ErBa2Cu3O7−δ film with BaNb2O6. Physica C: Supercond., 463-465:895–
899, 2007. 41, 59, 112, 120
[120] Kai H et al. Influence of Growth Temperature on Microstructure and Su-
perconducting Properties of ErB2Cu3O7−δ Films With Ba(Er0.5Nb0.5)O3
Nanorods. IEEE Trans. Appl. Supercond., 19:3435–3438, 2009. 41
[121] Kai H et al. The effects of growth temperature on c-axis-correlated pinning
centers in PLD-ErB2Cu3O7−δ films with Ba(Er0.5Nb0.5)O3. Supercond. Sci.
Technol., 23:025017, 2010. 41
[122] Yamada K et al. Transmission electron microscopy characterization of
nanorods in BaNb2O6-doped ErB2Cu3O7−δ films. Appl. Phys. Lett.,
92:112503, 2008. 41
[123] Feldmann D M et al. Improved flux pinning in YBa2Cu3O7 with nanorods
of the double perovskite Ba2YNbO6. Supercond. Sci. Technol., 23:095004,
2010. 43, 145
[124] Hun Wee S et al. Formation of Self-Assembled, Double-Perovskite,
Ba2YNbO6 Nanocolumns and Their Contribution to Flux-Pinning and Jc
in Nb-Doped YBa2Cu3O7−δ Films. Appl. Phys. Express, 3:023101, 2010.
43, 104, 109, 145
[125] Barnes P W et al. Structure determination of A2M
3+TaO6 and A2M
3+NbO6
ordered perovskites: octahedral tilting and pseudosymmetry. Acta Crystal-
logr., B62:384–396, 2006. 45
157
REFERENCES
[126] Goyal A et al. Irradiation-free, columnar defects comprised of self-assembled
nanodots and nanorods resulting in strongly enhanced flux-pinning in
YBa2Cu3O7−δ films. Supercond. Sci. Technol., 18:1533–1538, 2005. 50
[127] MacManus-Driscoll et al. Systematic enhancement of in-field critical cur-
rent density with rare-earth ion size variance in superconducting rare-earth
barium cuprate films. Appl. Phys. Lett., 84:5329–5331, 2004. 50
[128] Singh R K et al. Control of surface particle density in pulsed laser deposition
of superconducting YBa2Cu307 and diamondlike carbon thin films. Appl.
Phys. Lett., 61:483–485, 1992. 51
[129] Wu W B et al. Particulate-free a-YBa2Cu307−x/La0.5Sr0.5CoO3/a-
YBa2Cu307−x sandwiched thin film on (100)SrTiO3 fabricated by pulsed
laser deposition with a shadow mask. Physica C, 282:701–702, 1997. 51
[130] Boffa V et al. High-quality surface YBCO thin films prepared by off-axis
pulsed laser deposition technique. Physica C, 276:218–224, 1997. 51
[131] Backen E et al. Improved pinning in YBCO based quasi-multilayers pre-
pared by on- and off-axis pulsed laser deposition. IEEE Trans. Appl. Su-
percond., 17:3733–3736, 2007. 51
[132] Cillessen J F M et al. Improved uniformity of multielement thin films
prepared by off-axis pulsed laser deposition using a new heater design. Rev.
Sci. Instr., 67:3229–3237, 1996. 51
[133] Harrington S A et al. Understandig nanoparticle self-assembly for a strong
impovement in functionality in thin film nanocomposites. Nanotechnology,
21, 2010. 56, 112
[134] Maiorov B et al. Synergetic combination of different types of defect to
optimize pinning landscape using BaZrO3-doped YBa2Cu3O7−δ. Nature
Materials, 8, 2009. 56, 63, 89, 90, 114
[135] Liu L F et al. Effect of Deposition Temperature on the Epitaxial Growth
of YBCO Thin Films on RABiTS Substrates by Pulsed Laser Deposition
Method. IEEE trans. on appl. supercond., 20:1553–1553, 2010. 96
158
REFERENCES
[136] Norton M G et al. Surface outgrowths on laser-deposited YBa2Cu3O7 thin
films. Physica C, 233:321–326, 1994. 96
[137] Guo L P et al. Microstructure of outgrowths on the surface of laser-ablated
YBa2Cu3O7 thin films. Physica C, 241:30–36, 1995. 96
[138] Kai H et al. The effects of growth temperature on c-axis-correlated pinning
centers in PLD-ErBa2Cu3O7−δ films with Ba(Er0.5Nb0.5)O3. Supercond.
Sci. Technol., 23, 2010. 104
[139] Holesinger T G et al. The microstructure of continuosly processed
YBa2Cu3Oy coated conductors with undelying CeO2 and ion’beam’assisted
yttria’stabilized zirconia buffer layers. Jour. of Mat. Res., 15:1110–1119,
2000. 108
[140] Specht E D et al. Stacking faults in YBa2Cu3O7−x: Measurement using x-
ray diffraction and effects on critical current. Appl. Phys. Lett., 89:162510,
2006. 108
[141] Gapud A A et al. Enhancement of flux pinning in YBa2Cu3O7−δ thin
films embedded with epitaxially grown Y2O3 nanostructures using a multi-
layering process. Supercond. Sci. Technol., 18:1502–1505, 2005. 108, 145
[142] Garrison S M et al. Observation of 2 in-plane epitaxial states in
YBa2Cu3O7−δ films on yttria-stabilized ZrO2. Appl. Phys. Lett., 58:2168–
2170, 1991. 109, 145
[143] Long N J. Model for the angular dependence of critical currents in technical
superconductors. Supercond. Sci. Technol., 21, 2008. 117
[144] Zhang J L and Evetts J E. BaZrO3 and BaHfO3 - preparation, properties
and compatibility with YB2Cu3O7−x. J. Mater. Sci., 29:778–781, 1994. 145
[145] Hanish J et al. Formation of BaIrO3 precipitates and their contribution to
flux pinning in Ir-doped YB2Cu3O7−x quasi-multilayers. Appl. Phys. Lett.,
86:122508, 2005. 145
159
REFERENCES
[146] Alok K J and Khare N. Strongly enhanced pinning force density in
YBCOBaTiO3 nanocomposite superconductor. Physica C Supercond.,
469:810–813, 2009. 145
[147] Shingai Y et al. Improvement of superconducting properties by BaTiO3
doping into ErB2Cu3O7−δ films. Physica C Supercond., 445-448:841–844,
2006. 145
[148] Mele P et al. Incorporation of double artificial pinning centers in
YB2Cu3O7−δ films. Physica C Supercond., 468:1631–1634, 2008. 145
160