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Microstructure/electrical property correlations for YBa2Cu3O7-x/barrier layer films deposited on Al2O3, silicon, and yittria-stabilized zirconia substrates

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Microstructure/electrical property correlations for YBa2Cu3O7-x/barrier layer films deposited on Al2O3, silicon, and yittria-stabilized zirconia substrates
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Microstructure/electrical property correlations for YBa2Cu3O7-x/barrier layer films deposited on Al2O3, silicon, and yittria-stabilized zirconia subst
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Mueller, Carl Henry,
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Barrier layers ( jstor )
Conceptual lattices ( jstor )
Electrons ( jstor )
Film criticism ( jstor )
Grain boundaries ( jstor )
Oxygen ( jstor )
Temperature resistance ( jstor )
Thermal stress ( jstor )
X ray diffraction ( jstor )
X ray film ( jstor )

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MICROSTRUCTURE/ELECTRICAL PROPERTY CORRELATIONS FOR
YBa2Cu3O7_JBARRIER LAYER FILMS DEPOSITED ON A1203, SILICON, AND
YITRIA-STABILIZED ZIRCONIA SUBSTRATES













By

CARL HENRY MUELLER


A DISSERTATION PRESENTED TO THE GRADUATE SCHOOL
OF THE UNIVERSITY OF FLORIDA IN PARTIAL FULFILLMENT
OF THE REQUIREMENTS FOR THE DEGREE OF
DOCTOR OF PHILOSOPHY

UNIVERSITY OF FLORIDA


1992














ACKNOWLEDGEMENTS


Several people have helped me complete this thesis. I would like to thank my

advisor, Professor Paul Holloway for being an excellent role model and showing me

how a scientist solves problems. I am especially grateful for his guidance during the

early stages of the project, and for letting me choose my own directions as the project

matured. I was fortunate to have very talented people on my committee, and

Professors Abbaschian, Anderson, Connell, and DeHoff each provided insights into

materials properties which were important to this project and will be valuable

throughout my career.

I am grateful for the help I recieved from co-workers within our group, especially

Kelly Truman and Ludie Hampton. Much of the data was collected at the Major

Analytical Instrumentational Center at the University of Florida, and Eric Lambers,

Wayne Acree, and Richard Crockett did a superb job of keeping the instruments in

top condition, which made the data collection and interpretation possible.

I would like to thank Drs. Kul Bhasin, Felix Miranda, Mark Stan, and Crystal

Cubbage of NASA Lewis Research Center for their help. A large portion of this

thesis would not have been possible without their help.








My family provided a strong emotional base which allowed my to complete the

program. I would like to thank my parents, Gerhard and Lillian Mueller, for

stressing the importance of finishing a project. I would also like to thank my

brothers, Don, Lloyd, and Keith, my sister Jan, and their families for their

encouragement and for helping to keep things in perspective.

Finally, I would like to thank the friends who made my stay in Gainesville so

enjoyable. I will miss the Sunday afternoon soccer games.



















TABLE OF CONTENTS

ACKNOWLEDGEMENTS ..................................... ii

ABSTRACT ................................................ vi

CHAPTER

1. INTRODUCTION ......................................... 1

2. LITERATURE REVIEW ................................... 4

D evices ............................................... 4
Microstructure/Superconductor Relations ..................... 18
W eak Link Behavior .................................... 27
Nucleation and Epitaxial Growth ........................... 35
Thermally Induced Stresses ................................ 46
YBa2Cu3O7. Film Growth ................................ 53
Barrier Layer Technology ................................. 67

3. EXPERIMENTAL TECHNIQUES ........................... 81

Film Growth by Laser Deposition ........................... 80
X-ray Diffraction ....................................... 84
Scanning Electron Microscopy ............................. 86
Scanning Auger Electron Spectroscopy ....................... 86
Electrical Resistance Measurements ......................... 90
Critical Current Density .................................. 94
Raman Spectroscopy .................................... 96
Millimeter-Wave Transmission Measurements .................. 99

4. RESULTS ......................... ................ 101

YBa2Cu3O. on (1i02) LaA10 ........................... 101








YBa2Cu307.x/Y-ZrO2 Films on Si, Y-ZrO2, and LaAO03 Substrates .108
YBa2Cu307./SrTiO3 Films on A1203 Substrates ................ 137
YBa2Cu307.x/LaA103 Films on Si and A1203 Substrates .......... 154
YBa2Cu307x/YA103 Films on Si and A1203 ................... 160
YBa2CuO37./Y203 Films on Si, Y-ZrO2, and SrTiO3 Substrates .... 173
YBa2Cu3O7./(YA103, LaA1O3, or Y203)/Al6Si2013 Films on
Si and A1203 ..................... ................... 180



5. DISCUSSION .......................................... 196

Intergranular Versus Intragranular Effects ................... 196
Effect of Y-ZrO2, Y203, and ZrO2 Barrier Layers on Jc Values .... 201
Dependence of To and J. on A1203 Substrate Orientation ........ 209
Estimate of Stress and Cracking Due to Differential
Thermal Expansion ................................... 211
Effects of Texture and In-Plane Alignment ................... 212
Millimeter-Wave Properties .............................. 214
Effects of Surface Energy ................................ 215
Effects of Lattice Matching ......................... ...... 217
Effects of Oxygen Pressure on SrTiO3 Growth ................ 219
LaA103 Barrier Layers .................... .... ........... 225
Y203 Barrier Layers ................... ............... 228
YA103 Barrier Layers ................................... 229
Al6Si2013 Barrier Layers .......................... ........ 231

6. SUMMARY AND CONCLUSIONS .......................... 236

APPENDIX A. CALCULATION OF X-RAY ABSORPTION DEPTH
FOR YBaCu30 ............................... 241

APPENDIX B. FLOW CHART OF THE Y Al Si Cu O SYSTEM 243

REFERENCES ........................................... 266

BIOGRAPHICAL SKETCH ................................... 276













Abstract of Thesis Presented to the Graduate School
of the University of Florida in Partial Fulfillment of the
Requirements for the Degree of Doctor of Philosophy

MICROSTRUCTURE/ELECTRICAL PROPERTY CORRELATIONS FOR
YBa2Cu307./BARRIER LAYER FILMS DEPOSITED ON A1203, SILICON, AND
YITTRIA-STABILIZED ZIRCONIA SUBSTRATES

By

CARL HENRY MUELLER

December, 1992


Chairperson: Professor Paul Holloway
Major Department: Materials Science and Engineering

YBa2Cu3O7x and barrier layer films were deposited on single-crystal silicon (Si),

A1203, yittria-stabilized zirconia (Y-ZrO2), SrTiO3, and LaA103 substrates. A pulsed

laser deposition process was used to deposit the films at a substrate temperature of

730 750 *C, and the films were cooled in an oxygen ambient. The films were

characterized using resistance versus temperature, critical current density (Je), x-ray

diffraction (XRD), scanning electron microscopy (SEM), Auger electron spectroscopy

(AES), and Raman spectroscopy.

Growth of barrier layers on Si and A1203 substrates prior to the superconductor

suppressed chemical interdiffusion between the superconductor and substrate. For

(1102) A1203, the best barrier layer was a SrTiO3 film deposited at 200 mTorr of

oxygen. The YBa2Cu307. film had a zero resistance temperature of 83 K, and the








Jc was 2.5x106 amps/cm2 at 4.5 K. The surface resistance was 10-2 ohms at 36

gigahertz.

On silicon substrates, YBa2Cu3O07 degradation is aggravated by thermal stresses

created by the difference in thermal expansion coefficients between YBa2Cu307- and

Si (13.2 versus 3.8x10-6/*C, respectively), which causes microcracking in the

YBa2Cu307.x films. Cracking and interdiffusion were minimized by depositing a

YA103 barrier layer prior to YBa2Cu3O7-.. The thermal stresses were relieved by

viscoelastic relaxation in the YBa2Cu3O7. film, and the To was 78 oK.

The Jc values of YBa2C3Cu307 films on Y-ZrO2 substrates were increased by

depositing Y-ZrO2 or Y203 barrier layers. YBa2Cu3O7-x/Y2O3 films on Y-ZrO2

substrates had Jc values of 9 10 and 1xl10 amps/cm2 at 77 and 4.5 K. The Jc of

YBa2Cu3O7. films deposited on a Y-ZrO2 substrate without a barrier layer was

6.8 x 103 amps/cm2 at 4.5 *K. The higher Jc values were attributed to pinning of the

magnetic flux by excess Y203 at high-grain boundaries.













CHAPTER 1
INTRODUCTION



There are several commercial and military application for thin films which are

superconducting at temperatures above 77 "K. Since most semiconducting

devices perform optimally near 77 K, the potential for high-speed devices with

low-attenuation superconducting interconnects is promising'. In another

electronics area, passive microwave and millimeter-wave devices patterned into

superconducting thin films dramatically outperform the resonators and filters

presently being used.2 As the technology for depositing superconducting films on

A1203 substrates improves, this performance will be further enhanced. Potentially,

the largest applications for superconducting films are in power applications.3

Transmission cables, motors, and high field magnets which are too costly to

operate at 4.5 K would be economically feasible at 77 *K. Progress in

fabricating superconducting wires has been difficult because the wires currently

being fabricated suffer from brittleness and low critical current density (J,) values

(3 x104 amps/cm2 at 77 K). A better technique for making superconducting

cables may be to deposit superconducting films on fibers such as yittria-stabilized

zirconia (Y-ZrOz) which are mechanically strong and have a similar thermal

expansion coefficient as the superconducting film.










For each of these applications, Jc values greater than 1l amps/cm2 are

required.3 The sensitivity of J, to interfacial phases and high-angle grain

boundaries in YBa2Cu3O7. is one of the most significant problems, and it has

inhibited commercialization of YBa2Cu3O7.x films. To date, YBa2Cu3O7. films

with JC values above 10s amps/cm2 at 77 K have only been deposited on single-

crystal substrates such as SrTiO3, LaA103, and Y-ZrO2 which nearly lattice match

YBazCu3O7x. The ability to deposit films with high Jc values on randomly

oriented or polycrystalline substrates would be a tremendous boost towards

commercialization, since this would allow a variety of substrates with different

shapes and sizes to be used.

In this thesis, the literature is reviewed in chapter 2 in order to provide a

background for this work. The experimental techniques used to deposit and

characterize the films are described in chapter 3. The data original to this thesis

is presented in chapter 4, and is discussed in detail in chapter 5.

Chapter 2 shows how superconductivity evolves. The temperature at which

the resistance disappears (To) and Jc are closely tied to the microstructure, and an

overview of the microstructural phenomena which most critically affects the

superconducting properties is introduced. The laser deposition process is

described, and the mechanisms which enable multicomponent films to grow with

the same stoichiometry as the target are presented. Chapter 2 concludes with a

survey of the work directed towards depositing YBa2Cu3O7 films on various

substrate materials.










Chapter 3 describes the film growth technique, and the methods used to

characterize the films. Each of the characterization tools relies on different

physical phenomena to probe the film microstructures, so different aspects of the

microstructures were uncovered. A brief description of the operating principles of

each of the probes is given in this chapter.

The experimental results are presented in chapter 4. The sections are

arranged so as to compare the effectiveness of each barrier layer to induce growth

of optimal quality YBa2Cu3O7. films on different substrates. The text explains

what information the data is providing, and points out the most notable features

in each figure.

Chapter 5 compares and contrasts the data in order to explain the observed

phenomena. The primary objective of this thesis is to uncover the microstructural

features which were primarily responsible for the normal state and

superconducting properties. By using a variety of experimental techniques, we

arrived at a more clear understanding of film microstructures than if only one or

two techniques were used; thus we were able to correlate the film microstructures

with the electrical properties.













CHAPTER 2
LITERATURE REVIEW



Devices


Much of the interest in high temperature superconductivity is due to the large

number of applications which could benefit from replacing normal metals and

conventional solid state electronic devices with superconducting materials.

Basically, applications which could utilize high temperature superconducting thin

films can be divided into two categories: active devices which require one or

more weak link junctions, and passive devices which utilize the intrinsic properties

of superconductors for enhanced performance.

In superconductors, zero resistance and macroscopic quantization occur

because the wave functions of the electron or hole pairs are coherently coupled to

each other. This coherence enables calculations of the phase difference between

two different points of the superconductor to be made.4 Weak-link junctions are

thin insulating or normal metal regions which separate two superconducting

volumes, and the lower density of superconducting charge carriers causes a

gradient in the phase of the superconducting electron (or hole) wave function

across the weak link.








5
A Josephson junction is a special type of weak link junction in which the

relationship between the supercurrent, I,, and the gradient of the phase of the

superconducting wave function (AL ) is given by:s'6

1, = Isin (A Q) (2-1)

where I, is the maximum superconducting current which can travel through the

junction. For the rest of the discussion on devices, it will be assumed that the

weak links are Josephson junctions. The current vs. voltage behavior of a

Josephson junction is characterized by a zero resistance supercurrent which

persists when A~Q is constant, and a non-linear voltage which appears when A4Q
varies with time. The voltage, V, across a Josephson junction in this state is given

by:


2eV= h a(A ) (2-2)
27r &t

where e is the carrier charge and h is Planck's constant.

The I-V plot for a Josephson junction is given in figure 2-1,7 and virtually

all of the active superconducting devices utilize the behavior predicted by

equations 2-1 and 2-2. A brief overview of some of the devices which use

Josephson junctions, and their operating principles, is presented below.

Oscillators are devices which generate a repeating voltage vs. time waveform.

Superconducting oscillators, in which the output voltage is controlled by the

frequency of the supercurrent travelling through the junction (equation 2-2), can









6

generate electromagnetic radiation at milliwatt-level powers,8 at frequencies up to

1012 Hz.














Critical Current

M ------ -------------
0o
0r


VOLTAGE
VOLTAGE


Figure 2-1. The current versus voltage characteristic of a Josephson junction.
Reference 7.



Radiation detectors are based on the principle that electronic radiation will

induce AC currents in the Josephson junction.7 Regardless of the orgin, currents

add algebraically, hence the DC critical current will be suppressed by the

radiation induced AC currents. In the finite voltage regime, the mixture of AC










Josephson supercurrent with induced normal state AC currents will produce sum

and difference harmonic currents within the junction. When the ac Josephson

frequency and radiation frequency are harmonics of each other, there is a mixing

harmonic, and a step in the current vs. voltage spectrum appears at zero

frequency and at the mixing frequency. By varying the voltage and thereby the

supercurrent frequency across the junction, a large range of frequencies can be

sampled via the conversion factor:


1 microvolt = 484 MHz. (2-3)

Surface Quantum Interference Devices (SQUID's) are used to detect weak

magnetic fields, and operate on the principle that the number of flux lines which

pass through a Josephson junction must be quantized.4'9 If the magnetic field

being detected is not a quantized number of fluxons, the J. of the junction will

change so that the magnetic flux from the magnetic field plus the flux generated

by the current through the junction is quantized (figure 2-2). The magnetic

intensity of a fluxon, o0, is given by:


Q = h = 2.07 x 10-15 Webers (2-4)

Thus the I V profile of a Josephson junction is modulated by the magnetic field.

In practice, the magnetic sensitivity of a de SQUID is significantly improved by

placing two Josephson junctions in parallel (figure 2-3), which makes the loop

defined by the film plus the two junctions the area through which the modulating










magnetic flux travels. This dramatically increases the area over which the flux is

measured. For two junctions placed in parallel,


(2-5)


1,(10) = I, o) cos0' )

The high sensitivity of SQUID devices make them attractive for a variety of

biomedical, geological, and military applications.












0


-J


F:
o


0 o 20
9oA 0oA
APPLIED MAGNETIC FIELD STRENGTH, Ha



Figure 2-2. Periodic variations in the critical current density with increasing
magnetic field strength. Reference 4.




























1/21X1 W 1/2 1 V












Figure 2-3. Schematic of a two-junction SQUID. Reference 4.



Because Josephson junctions can be changed from superconducting to

semiconducting and vice-versa by varying the current or magnetic field, they have

been used as switching devices in digital electronics. The intrinsic switching

speed of a Josephson junction is:











Switching speed = h (2-6)
2 rA (7)

where A(T) is the temperature dependent superconducting energy gap. The
intrinsic switching speed for niobium-based Josephson junctions is 0.22

picoseconds. As an example of the increased performance which can be achieved

by replacing semiconducting switching devices with superconducting Josephson

junctions,10 a four-bit data processor made with GaAs transistors had a clock

speed of 72 MHz, and a power dissipation of 2.2 watts. A processor which

performed the same functions using Josephson junctions had a clock speed of 770

MHz and dissipated 5 milliwatts. More recently, superconducting electronics were

used to fabricate a four-bit shift register using 3 um linewidths, which operated at
9.6 GHz and dissipated 40 microwatts. By comparison, devices made from GaAs

or Si which used 0.5 pm linewidths, dissipated approximately 100 milliwatts. The
lower power dissipation associated with Josephson switching devices is an

important advantage, since heat generation and removal limit the density and

bandwidth of circuits based on semiconducting electronics.

Superconducting devices employing Josephson junctions have great potential

and, in theory, could completely change the materials and operating principles on

which high speed electronics are currently based. However, reproducible and

reliable Josephson junctions are difficult to fabricate. Efforts to create high

speed, commercially acceptable switching devices based on low temperature

superconducting electronics have not been successful. Because the high

temperature superconductors are much more sensitive to microstructural defects










than their low temperature counterparts, it is unlikely that high temperature

superconductors will be used as switching devices in the near future. While

devices which utilize fewer Josephson junctions, such as electromagnetic detectors

and SQUID's are more feasible, difficulties in creating reliable Josephson

junctions remain formidable.

The materials challenges presented by the second group of devices based on

superconducting materials, passive devices, are more likely to be surmounted in

the near future. The first commercial applications for high temperature

superconductors will probably emerge from this group. The largest applications

of passive superconductors will be as transmission lines, delay lines, and filters for

receiving and transmitting electromagnetic signals.

The advantages gained by replacing normal metals with superconducting

materials are derived from the magnetic field penetration depth:"


(T) -- x(0)
[ 4-( ] (2-7)

where A(T) is the temperature dependent magnetic penetration depth, A(0) is the
penetration depth near 0 *K, T is the operating temperature, and To is the

highest temperature at which zero resistance is observed. The response of a

superconducting material to an applied magnetic field is shown in figure 2-4. For

a superconductor, A is a material property, so it is independent of the device
operating frequency.







































x
,Penetratlon:
Depth




Figure 2-4. Variation of magnetic flux density at the boundary of a
superconductor. Reference 4.










In microwave and millimeter-wave transmission lines, there is an electric field

induced in the superconducting line because of the inertia of superconducting

electron pairs to the electromagnetic field. A surface resistance results because

the normal-state electrons are excited by the electric field, and the surface

resistance is qualitatively a measure of how much of the electromagnetic signal is

lost as heat in the transmission line:6


Joule losses
Surface resistance = R= Joul losses, (2-8)

where Hsufa is the magnetic field intensity inside the superconductor. In

superconductors, the depth to which a magnetic or electric field can extend is

limited by the penetration depth, and the surface resistance (R,) is given by:6'12


R,(T,) = A2 exp-A (2-9)
T kT

where A is a constant, o is the frequency of the electromagnetic signal, T is the
temperature, kB is Boltzmann's constant, and A(T) is the superconducting energy

gap.

By contrast, the surface resistance of normal metals is given by:13


R,= ( o) (2-10)
2oo

where po is the magnetic permeability of the metal, and o = a1 + iO2 is the

complex electrical conductivity, with,











= 0 2 (2-11)
1 + M2r2

and


o= o (2-12)
1 + o22

a, and a2 are the conduction and displacement currents, and r is the mean time

between electron-phonon interactions. In normal metals, the skin depth (6) is the
distance into the metal which an electric field can penetrate, and is a function of

the electrical conductivity:


6 = (2-13)
(27rOowr)o.5

where c is the speed of light (3x1010 cm/sec) and p is the magnetic permeability.

The dependence of the skin depth on electrical conductivity and hence

frequency means the dielectric constant of the metal changes with frequency. The

parameters pertinent to microwave signal transmission are contained within the

propagation constant, y,14 where


y=a+rjp (2-14)

a is the attenuation constant, and f is the wavenumber of the microstrip line. f
is given by:











S2 (2-15)


where Ag is the wavelength of the electromagnetic signal as it propagates in the

microstrip line. The speed at which electrical signals propagate through the metal

is given by the phase velocity, vp, where


v 1 (2-16)
p (pe).

p = PoUr and E = EoE, where Po and E, are the magnetic permeability and
dielectic permittivity of free space, and (uo0o)"5 is equal to the speed of light.

Given that /r and Er are the relative permeability and dielectric constant of the

microstrip material, the phase velocity of an electromagnetic signal propogating in

the metal is given by:


S c (assuming P, = 1) (2-17)
(e,)"o

Changes in the electrical conductivity, and hence dielectric constant of the metal

as a function of frequency, will also result in phase velocities which vary with

frequency. This phenomena, termed dispersion, causes electrical pulses composed

of various Fourier components to become spread out as they propagate along a

transmission line, and is the primary reason why normal metal conductors are

inadequate for long transmission lines. Dispersion in metallic interconnect lines

makes them inadequate for delay times greater than 1 microsecond, or for

wideband lines carrying short pulses (figure 2-5). For example, assuming typical










dimensions for a VLSI microstrip line (thickness = 0.1 0.3 um, width = 1 3

jum, and length = 1 10 pm) dispersion necessitates that pulse lengths longer than
100 picoseconds must be used. However, a superconducting line of the same

dimensions would have negligible dispersion for pulses longer than 1

picosecond.1516

1011

1010
9 : Superconducting Line
>- 109

S108
-j
S107 Aluminum Line

U) 106

S105 I
0 2 4 6 8 10 12
LOG (Frequency)

Figure 2-5. Calculated phase velocities as a function of frequency for
superconducting and aluminum microstrip lines, at 77 K. Reference
15.



The real part of the propagation constant, a, is the attenuation and results

from Joule losses within the microstrip line. In a superconducting material, the

penetration depth is constant, so a (decibel/cm) is linearly proportional to the

number of wavelengths/cm, and thus the frequency. In normal metals, a is

dependent on the skin depth, which is proportional to o0"5, so a increases more

slowly with increasing frequencies in metallic microstrip lines than it does in

superconducting lines. However, at microwave and millimeter wave frequencies,










the superconducting penetration depth is much less than the normal metal skin

depth, and it is only at w > 1012 Hz that attenuation in a superconducting

stripline approaches a for normal metals (figure 2-6).16'17

In semiconductor devices, the reduced attenuation and Joule heating provided

by superconducting interconnect lines would permit interconnect lines to be scaled

down considerably, thereby reducing chip to chip propagation delays.


10-1



U,
E
. 10-2
0
LLF



CO
LU
LU
0C

L 10-4


10-5

10-5


1 10
FREQUENCY, GHz


100


Figure 2-6. Surface resistances of YBa2Cu30., (dotted line) and copper (solid
line) films deposited on LaAlO, substrates. Reference 17.













Microstructure/Superconductor Relations


The microstructure of a superconducting film has a large influence on the

electrical and magnetic properties, and understanding how the superconducting

properties are impacted by film microstructure will dictate the processing

procedures used to fabricate the films. In YBa2Cu3O7., the superconducting and

normal-state properties are anisotropic with respect to crystallographic direction,1

so film orientation is a critical parameter. The To and Jc of the film are

dependent on superconducting properties such as coherence length (4), and
penetration depth (A), which, in turn, are influenced by orientation, grain size,
and grain boundaries. Clearly, microstructural considerations play a large role in

defining the electrical and magnetic properties of the film. To understand the

relationship between microstructure and superconductive behavior, it is instructive

to review how superconductivity evolves. Although the physical processes

responsible for superconductivity in YBa2Cu30.x have not been uncovered, most

of the properties are well described by the microscopic Bardeen-Cooper-Schrieffer

(BCS) theory and the phenomenological equations which predate BCS, so the

properties of YBa2Cu30.x will be described in terms of conventional

superconductors.

In a normal metal, there is a repulsive energy between electrons, and the

energy levels are described by Fermi-Dirac statistics:19,2











( (E E) (2-18)
e +1

where f(E) is the probability that a given electron state is occupied, E is the

energy, Ep is the Fermi energy, k is Boltzman's constant, and T is the

temperature. In normal metals, electrical conductivity is possible because there

are empty electronic states at energies greater than Ep, into which electrons can

hop and thus move through the lattice. Current is transported through a material

because the applied electric field raises the electron energy level of one terminal

relative to the other, thus increasing the probability that electrons will hop from

the high potential to the lower potential region of the sample. Electrical

resistance arises from processes which transfer kinetic energy to the crystal lattice

by electron-electron and electron-phonon collisions.

The mechanisms for current transport in superconductors are different from

those observed in normal metals. Superconductivity is a cooperative phenomena

involving many electron or hole pairs, and is possible because of electron (hole)-

phonon coupling, which creates an attractive force between electrons (holes). The

potential energy caused by an electronic transition from an initial state k, to

another state k', is given by:6










2
V(q,0) [1+ q ] (2-19)
q2+2 k 02 2

where V (q,m)is the electron-electron potential, q is the wave-vector difference
between the k and k' states, wq is the phonon frequency at wave-vector q, and k is

the wave number of superconducting electrons at the Fermi surface. During an

electron transition between two states (k to k'), a phonon may be absorbed. If w
< Oq, and the transition lowers the potential energy of the system, the electron-

phonon interaction creates an energy gap in the energy vs. momentum spectrum,

and the most favorable way for two electrons with p > pf (p = momentum at the

Fermi level) to lower their energy below 2E, is to form a bound state, in which

electrons with equal and opposite moment combine to form Cooper pairs. The

wavefunction for a Cooper pair is given by:4


(iPp=e (2-20)

where is the amplitude of the wavefunction (also known as the "order
parameter"), multiplied by the travelling wave expression in which P = electron

momentum, r = position, h is Planck's constant, and I' 2 is the density of

superconducting electrons. Cooper pairs obey Bose-Einstein statistics, hence it is

energetically favorable for all the pairs can have the same momentum. Because

all the superconducting pairs have the same momentum, and thus the same

wavelength, superposition of the coherent waves results in another wave with the

same wavelength. Superconductors possess macroscopic quantization because all










the superconducting electrons have the same momentum. When a transport

current is impressed on the system, the quantum mechanical relationship


v,2m h(k +k2) (2-21)
2x

(where v, = velocity of superconducting electrons, and k1 = k2 = k. for the

superconducting electrons) produces long range order in the momentum. There is

no resistance because the superconducting electrons are coupled together, and the

energy required to break a Cooper pair into excited electrons (quasiparticles) is

greater than twice the energy gap between superconducting and normal state

electrons (= 2A(T)).
The size of the superconducting energy gap is dependent on the fraction of

electrons which are superconducting, and this fraction varies with temperature.

Gorter and Casimer derived:


.=1-( 4, (2-22)
n To

where n/n is the fraction of superconducting electrons, To is the temperature at

which resistance vanishes in weak magnetic fields, and T is the temperature. The

variation of the superconducting energy gap (A(T)) with temperature is given by:2

1
A (7)TT=aTo(1-I) (2-23)


where a is constant equal to 3.06 if weak electron-phonon coupling is assumed.

The variation of the superconducting gap with temperature is shown in figure 2-7,










1.0 -...e A- 0o indium

SA A Tin
0.8 o *Lead I
0** BCS Theory oA
o 0.6 o

60.4-

0.2-

0 I I I
0 0.2 0.4 0.6 0.8 1.0
T/Tc

Figure 2-7. The superconducting energy gaps of lead, tin, and indium versus
temperature. Reference 4.


and shows the gap is a maximum at T = O *K. To illustrate the fraction of the

ideal superconducting bandgap which commercial high temperature

superconducting devices will be expected to operate at, we assume To to be 90

K, and the device operating temperature to be 77 K. These values correspond

to the temperature at which zero resistance is achieved in YBa2Cu3O7x, and the

boiling point of liquid nitrogen. Using equation 2-23, we see that A(T)/A(0)
drops to 0.38 at T/Tc = 0.86, and disappears at T = To. Increasing the

superconducting energy gap at 77 K is one of the prime motivations for studying

higher T, materials such as the bismuth and thallium-based superconductors.

A parameter which critically affects film and device quality is the coherence

length (i), which is the decay length for the wave function created by the

formation of Cooper pairs,7
W '. (2-24)










where vp is the velocity of electrons at the Fermi level and A is the size of the

superconducting energy gap. An equivalent definition is that the coherence length

is the minimum distance over which the density of superconducting electrons can

vary, and hence the minimum distance between superconducting and

nonsuperconducting regions. The coherence length (4) is much shorter in

YBa2Cu37O.x than it is in conventional superconductors such as Nb3Sn, and is

highly dependent on crystallographic direction (table 2-1). The short coherence

length in YBa2CU3O7-. (table 2-2) is primarily due to the small Fermi velocity of

the superconducting electrons, and an important consequence of the short

coherence length is that microstructural defects, such as grain boundaries,

impurity atoms, dislocations, or chemically unstable surfaces which create

imperfect or disordered regions of similar size to the coherence length, can

significantly alter the superconducting wave function, especially in the <001>

direction. By contrast, microstructural defects are considerably less degrading to

the order parameters of conventional superconductors.



Table 2-1. Normal state parameters of conventional metals (such as Nb3Sn) and
YBa2Cu3O7x. Reference 7.

Parameters Conventional YBa2zCU30OT
Metals
Metals (001) 1 (001)
m* 1-1.5 m, 5 me 25 mg
Ep (eV) 5-10 0.3 0.3
kF (cm)- 108 5x107 5x107
vp (cm/sec) 1-2x108 107 2x106










A final parameter which is important for understanding the magnetic and

electrical transport properties of superconductors is the magnetic penetration

depth, A, which is a measure of how deeply a magnetic field penetrates into a
superconductor. The depth to which a magnetic field can penetrate into a

superconductor is limited because the superconductor will generate circular eddy

currents, which create an internal magnetic field that nullifies the applied field. A
is defined as4


f B(x)d= Bo (2-25)

where Bo is the applied magnetic field, and B(x) decays exponentially as it enters

the superconductor (figure 2-4). By virtue of the Maxwell equation,21


VxB=E +4nJ, (2-26)
9t

A is the maximum depth at which a transport current (J) can flow, hence it is
analogous to the skin depth in normal metals.



Table 2-2. Superconducting parameters of conventional metals (such as Nb3Sn)
andYBa2Cu307., Reference 7.


Parameter Conventional YBa2Cu307-
Metal
Metal (001) 1. (001)

To (OK) < 23 95 95
2A/kBTo < 4.4 5-8 2-3.5
A/E, 104 2x10-1 1xlO-1
o (A) 10'-10 15 7








25
The penetration depth varies with superconducting electron density, and hence

temperature. The Gorter-Casimir expression for A is:"


[1( 0o 1 (2-7)
T[-( ) 4 2


where A is the penetration depth at 0 K, and varies with film quality. For a
highly-(001)-oriented, 3000 A thick YBa2Cu3O0.. film grown on LaAIlO, 10 = 1800

A was measured.17 Films with larger fractions of non-(001) oriented grains, or

which contain high-angle grain boundaries, had larger penetration depths.

The way in which a superconductor responds to a magnetic field is dependent

on the Ginzburg-Landau parameter K, defined as4


K =0.966- (2-27)
to

The magnetic field may be applied, or may be induced by the transport current (a

self-field). Superconductors with K < j 2 are classified as Type 1, while those
with K > 2 are Type 2. In both types, there is a magnetic contribution which
increases the free energy density, AG, while the electron ordering associated with
the formation of Cooper pairs lowers AG. If the AG associated with electron
ordering occurs over a shorter distance from the surface than the magnetic

contribution ( as in Type 2 superconductors), there will be a minimum in the free

energy density at the surface, and it becomes energetically favorable to form an

interface between the normal and superconducting regions (figure 2-8). Thus

there are isolated circular regions in the sample which are normal, while the rest









26

of the sample is superconducting. The coexistance of normal and superconducting

regions (i.e. the mixed state) enables superconductivity to be maintained in much

higher magnetic fields in type 2 materials than are allowed in type 1. All of the

commercially useful superconductors are type 2 materials.



Normal Superconducting
&Number of
SSuperelectrons
Magnetic
Flux
Density


(a) Penetration depth and coherence range


Magnetic
Contribution
Free
Energy
Density


S- -,-_-_ --- Electron-ordering
Contribution


(b) Contributions to free energy


Free
Energy
Density


(c) Total free energy


Figure 2-8. Orgin of negative surface energy. Reference 4.












Weak Link Behavior


Because of the short coherence length, microstructural defects significantly

affect the electrical properties of YBa2CU3O7-.. Defects are regions in which the

amplitude of the superconducting wave function is depressed, and the density of

superconducting electrons is lowered. Since these defects reduce the

superconducting order parameter, reductions in To and J, result if the transport

current is forced to travel through the defects.6

Defects are classified as either superconductor-insulator-superconductor (SIS)

or superconductor-normal-superconductor (SNS), depending on the electrical and

magnetic properties of the junction.4 For the case in which two superconductors

are separated by a thin insulating layer, there is a finite probability that

superconducting electron pairs will tunnel through the insulator and a

superconducting current will be maintained. Because of the insulating layer, the

wave function is no longer continuous between superconductors, but there is a

phase difference, A,, between the wavefunction at each of the interfaces which
determines the maximum supercurrent, I,, which can pass through the insulator:5


,= i Ain A (2-1)

If the transport current exceeds i, A4 is no longer constant with time, and a

voltage appears across the junction. The currents which result from the time

varying A~Q and subsequent voltage across the insulator are given by:4











Ch o2(A) + h c(Ai) (2-28)
4xe at2 R44xe at

where C is the capacitance, h is Planck's constant, and R, is the normal state

resistance of the junction. Ambegaokatar and Baratoff derived an expression

which shows that JC is inversely proportional to the normal state resistance,2


J(T) = ()tanhT A()]. (2-29)
2eRn 2kT

where A(T) is the superconducting energy gap.
Coupling between the superconductor wavefunctions on each side of a

junction, and thus tunneling current density, are rapidly diminished by increasingly

thick insulator regions:2


S c "c e "c (2-30)


where s is the insulator thickness (nm), a = 0.1 (eV)'-/nm is the effective barrier

height, and 4D the electron energy gap between valence and conduction bands in
the insulator. Many devices, such as detectors and oscillators, could be made

using superconductors if SIS junctions were more reproducible. However, since

superconducting areas within a distance smaller than the coherence length, , of
the junction are often degraded, A(T) is not the energy gap for a perfect material
but is depressed because of defects near the junction. Thus Jc for a given junction

is difficult to control.

When two superconductors are separated by a normal metal (SNS structure),

superconductivity is induced in the normal metal via the proximity effect. The










proximity effect entails diffusion of superconducting electron pairs and normal

electrons across the interface to create a weak superconducting layer in the

normal metal. There is a coherence length for superconducting pairs in the

normal metal which is given by:2

hvP
wasorIM a v, (2-31)

where vF is the Fermi velocity of the normal metal, and the electron mean free

path is assumed to be longer than the coherence length ("clean limit").

Conversely, if the electron mean free path is shorter than the coherence length

("dirty limit"), the coherence length is given by:


S-_ hvl I. (2-32)
12%kT)

where 1 is the electron mean-free path. The critical current across an SNS

junction is given by:



To

where B is a constant, and a is the normal metal thickness.

The short coherence length in YBa2Cu3O7- is the primary reason why

microstructural defects have a significant effect on the superconducting properties.

Dimos et al. conducted a series of experiments designed to determine the

dependence of Jc and To on temperature and magnetic field, and reveal the types

of weak links responsible for reduced superconducting properties.2 For epitaxial










YBa2Cu30.x films grown on (100) SrTiO3 substrates, where two substrates were

sintered together so as to deliberately produce a misorientation angle O in the
film yet insure the <001> directions in the YBa2Cu3O7. film on both sides of the

sintered SrTiO3 junction were parallel, variations in J, as a function of grain

misorientation showed a dramatic drop as the grain boundary angle was increased.

Inside the grains, Jc values of 4x106 amps/cm2 at 5 K were measured. However,

across the grain boundaries Jc values dropped from 4 x 106 amps/cm2 at 00

misorientation to 0.16x106 amps/cm2 at 12.50 misorientation.

Further increases in resulted in similar or lower J, values, and did not

follow a systematic trend. The authors observed that Jbw/JC was approximately

proportional to 1/0 for O < 200, and proposed a model in which the dislocation
spacing was the predominant factor in determining J,. The dislocation array acts

as a partial barrier to superconducting electrons, or perhaps as an easy path for

flux flow, since the superconducting order parameter is depressed at the

dislocation cores. Although various studies have observed significantly larger Jc

values (> 5x106 amps/cm2 at 77 K) in epitaxial (001)-oriented films as opposed

to polycrystalline films, this study was the first to indicate that larger J, would be

observed in (001) films with in-plane epitaxy and low grain boundary angles,

compared to to (001) oriented films with random in-plane orientation.

Later experiments by Mannhart et al. on epitaxial, (001)-oriented YBa2Cu3O7.

films grown on (100) SrTiO3 substrates determined that the values of

intergranular J. as a function of temperature could be fitted by the

Ambegtaokatar-Baratoff equation for a SIS junction in which an energy gap A(0)









31
of 5 meV was assumed.25 However, the Ambegaokatar-Baratoff expression is also

valid for dirty SNS junctions, provided that significant changes in the phase of the

wavelength only occur near or in the normal layer; hence the observed gap could

represent the reduced order parameter inside the normal layer. The authors

concluded that grain boundaries were dirty SNS junctions because Jc(0) was a

factor of 10 less than values predicted by the SIS model. On the other hand, J, vs

B(T) for different orientations of the magnetic field indicated intragranular J,

values were limited by flux creep across the grains, and intragranular Jc values

were determined by the density and pinning energies of the flux-pinning sites.

Because the sensitivity of electrical properties in the superconducting and

normal states to film microstructure is very pronounced, a great deal of

information about the film microstructure can be obtained from the resistance vs.

temperature data. Zero resistance occurs when there is a continuous pathway in

which the wave function of the superconductor is phase ordered along the entire

path. If the wave function is strongly coupled between grains, the transition from

the normal to the superconducting state is abrupt, and the film To is close to the

transition temperature observed in bulk samples. However, if the film

superconductivity is localized, meaning there are isolated superconductor volumes

such as grains separated by barriers at which the intergranular coupling is weak,

or if the coupling strength varies randomly at the different grain boundaries, the

transition to the superconducting state will be percolative and there will be a

finite resistance at temperatures below Tot. Deviations from the ideal resistivity










and superconducting transition of a YBa2Cu307. film can be attributed to

microstructural defects such as high angle grain boundaries, interfacial phases at

the grain boundaries, and randomly oriented grains. An expression for the normal

state resistivity, p, of a single crystal YBa2Cu3O7. film free of microstructural

defects was proposed by Halbritter:26


p (7) = UcT+ p0, (2-34)

where the superscript i denotes intragranular effects, a1 is a parameter which

contains the temperature dependence of intragranular normal state resistivity, and

PoL is the intragranular resistivity exptrapolated to T = O OK. For single crystal
YBa2CU30.x films, Halbritter reported poi = 0 and a' = 0.5 ohmxcm/K.

However, many YBa2Cu3O7.x films, especially those grown on substrates which are

reactive or do not lattice match YBa2Cu3O7-, are polycrystalline and contain grain

boundaries or other types of weak links which increase the percolation distance of

the electrical conduction network. Using a technique similar to Eom et al.,7

deviations from the ideal resistance vs. temperature curve for YBa2Cu370x can be

attributed to specific types of microstructural defects. Halbritter expanded the

equation to separate the normal state resistivity into intergranular (e.g. grain

boundary or cracks) and intragranular contributions caused by electron scattering

from other electrons, lattice ions, or defects within the grain:











P (7) = (p, +p(a 'T+ pL (2-35)
where z(pg,) is the sum of the resitivities associated with all grain boundaries, and

(aiT + POLi) describes the intragranular resistivity. The percolation parameter, p,

accounts for increases in the percolation distance of the transport current resulting

from microstructural defects including grain boundaries and cracks. When pg is

small compared to p(a' + PoL'), intergranular defects have little effect on p, and

separation of inter and intragranular defects is difficult. If the current pathway of

lowest resistance is through the grain boundaries, the net effect of a change in pg

on R vs. T is to shift the curve to higher resistance values, without changing the

slope (p values) of the curve. Extrapolation of the curve to 0 OK will result in

higher values of 2(pg) when grain boundary resistance becomes larger, but not so

large as to significantly alter the current pathways. However, when 2(Po,) is

comparable to or greater than the intragranular resistivity, the percolation

distance of the transport increases as the current selects pathways which avoid

grain boundaries or other microstructural defects such as cracks, thereby

increasing p.

The same defects which increase normal state resistivity are responsible for

depressions in To and J.2829 Intragranular Josephson junctions resulting from

oxygen disorder or twinning locally depress the order parameter, and thus the To

within the grain. The most damaging defects are usually intergranular, and are

observed at high angle (0> 20 ) grain boundaries where dislocated regions often
have significant oxygen disorder.25 At high angle grain boundaries,










superconducting electrons are preferentially transported through the grain

boundaries at microbridges where the oxygen disorder is minimal (figure 2-9).

Interfacial impurity phases increase the separation between superconducting

regions, further degrading the interface and making it more difficult for

superconducting electrons to tunnel through the grain boundary. Because the

coherence length of superconducting electrons is longer in the <100> and <010>

than in the <001> direction, To and J, in (001) oriented films are not as sensitive

to grain boundaries as are films with other orientations. However, high angle

grain boundaries are highly dislocated regions where impurity phases tend to

accumulate, so the order parameter is depressed in (001) oriented films with

random in-plane orientation. By contrast, low angle grain boundaries are

generally free of interfacial phases and do not depress superconductivity. In

addition, 900 grain boundaries, which usually result from twinning or adjacent

(010), (100) and (001) oriented grains, generally do not contain significant

amounts of impurity phases, and hence do not reduce To or J, values.30


CU -T


6I, CRO-
RIDGE

I CUPRATE-312SUPEP-CONDUCT

Figure 2-9. Intergranular defect with one microbridge. Cuprate regions near the
grain boundary are normal conducting, and the superconducting order
parameter is diminished in these regions. Reference 26.













Nucleation and Epitaxial Growth


The superconducting properties of YBa2Cu3O07, are very dependent on

microstructure. A general section which describes nucleation and growth

processes is presented in order to establish a framework in which the electrical

properties of the YBa2CCu3O7. films can be correlated with the growth processes.

Nucleation and growth of thin films is a large and complex subject, with many

variables which can potentially affect the microstructure and orientation of the

film. While it is difficult to obtain direct experimental evidence for many of the

parameters which determine growth mode, it is possible to speculate about the

forces which were operative during film growth by examining microstructural

properties such as surface morphology, grain orientation, and the distribution and

orientation of interfacial phases. Because superconducting and barrier layer films

have been deposited on several different substrates under a variety of growth

conditions, many of the growth mechanisms which determine film microstructure

can be deduced. In most deposition processes involving high temperature

superconductors in which the film is grown from vaporized constituents striking

the substrate, the experimental parameters which are varied include substrate

temperature, oxygen partial pressure, and the energies of the atoms and ions as

they strike the substrate. A general discussion of the fundamental processes

involved in film growth and how they are affected by growth conditions is









36
presented in order to establish a framework by which film microstructure evolves.

In the specific case of YBa2Cu3O7-. and barrier layer growth, correlations between

growth conditions and the underlying growth mechanisms will lead to an

understanding of how film quality can be optimized.

The crystallinity and orientation of a film can be manipulated by changing the

growth rate. In order to grow a film, more adatoms must stick to the surface than

are evaporated, hence there must be a flux causing a supersaturation of atoms or

ions reaching the surface and forming stable nuclei. The flux of atoms impinging

on the surface (0) is given by:31

atoms P(T)
0 ( atoms_ P(T) (2-36)
cm2xsec (2xmkT)o05

where P(T) is the vapor pressure at the substrate surface, m is the mass of the

impinging species at the substrate surface, T is the substrate temperature, and k is

Boltzmann's constant. To achieve epitaxy, it is essential that the atom be able to

move freely across the substrate until it reaches a potential minimum. The jump

frequency, o, of an adatom on a substrate is given by:

-E
= v exp(-) (2-37)
KT

where v = number of jump attempts/sec (typically 1013/sec), and ED is the

activation energy for surface diffusion. The mean stay time for an adatom is:32











1 Ed (2-38)
Ve = exp(p (2-38)

where Ed is the energy for desorption of an adatom from the substrate. The

mean distance an adatom diffuses before being desorbed is given by:


= 2aoexpl[ (E, (2-39)
2KT,

where a0 is the single jump distance. This result shows that the surface diffusion

length increases with decreasing substrate temperature because the mean stay

time is increased. However, epitaxy is not necessarily improved by lowering the

substrate temperature because the jump frequency is lowered, hence the rate at

which equilibrium is reached is lowered. When the number of adatoms diffusing

across the surface at a given time approaches the density of surface sites, energy is

lost by adatom-adatom collisions, and adatom-substrate interactions are

diminished, thereby decreasing epitaxy. For the remainder of this study, we

assume that the flux of impinging species was low enough, and the substrate

temperature sufficiently high that the number of adatoms diffusing along the

surface was negligible compared to the number of surface sites, hence the

adatoms were able to reach their lowest energy configuration and equilibrium

conditions prevailed.

Until the nucleus reaches a critical size, it is more likely that the atoms will

dissociate and join a different nucleus or desorb, thus sub-critical nuclei do not

participate in the film growth process. In a real deposition system, the flux of








38
atoms striking the substrate consists of clusters of atoms as well as single atoms or

ions. Although it is difficult to determine what is the smallest stable nucleus size,

it is logical to expect that larger clusters striking the surface are more likely to

remain and form stable nuclei than are smaller clusters or atoms, since there are

more overlayer-substrate bonds formed. In addition, increasing the substrate

temperature is expected to increase the critical nucleus size, so the distribution of

clusters or atoms which eventually form stable clusters is further skewed towards

the higher cluster sizes.33

The relationship between critical cluster size and substrate temperature is

important because the orientation of the entire film can be heavily influenced by

the size of the original clusters from which stable nuclei are grown. Theoretical

and experimental data on face centered cubic metallic films deposited on NaCl

substrates indicates that for a critical cluster size of 3, the atoms will adopt a

triangular arrangement, and the film will grow with a (111) orientation.34

Similarily, when the smallest critical nucleus size is 4 atoms, the nucleus will

arrange itself in a square or rectangular mesh, and (100) oriented growth is

favored. In general, the first monolayer will grow with the surface mesh initiated

by the stable cluster, and subsequent layers will maintain this orientation and

grow with the closest packed planes parallel to the substrate. Based on this

analysis, face centered cubic films in which the critical cluster sizes are 3 or 4

atoms are expected to grow with (111) and (100) orientations, respectively. It has

been observed that several metallic films which grow with a (111) orientation at










low temperatures will adopt the (100) orientation when deposited at higher

temperatures.s Likewise, yittria-stabilized zirconia (Y-ZrO2) films deposited on

(1102) A1203 grow with mixed (111) and (100) orientations at temperatures below

780 *C, and are predominately (100) orientated when deposited at higher

temperatures.36

The relative surface free energies of the substrate, film surface, and interfacial

layer are some of the most important parameters which dictate the film

morphology, and whether film-substrate epitaxy is possible. In a thin film, the

surface energy can be a substantial portion of the total energy of the system. For

a planar interface between two phases, a and f,37


a dA = dU'l _TdS' E i ,dn (2-40)

where a is the energy required to create a surface of area A, dUx"" and dS'"

are the energy and entropy associated with creation of the new surface, and pi and

dn, are the chemical potentials and excess surface concentrations of species i at

the surface. The general expression for the surface energy is:


dUXes = dUMt- _dUa -dUJp (2-41)

where


dU" = TdS" -PadV* +, ifdr (2-42)


and











dU = TdS -PPdVP + iAn, (2-43)

Substituting the expressions for dU" and dUf into equation 2-42, we obtain

odA = dU'- TdSPD- pn, + P" dV + PdV (2-44)

Assuming the volume of the system is constant, the Helmholtz free energy is

obtained:


dF = -S"~dT+ E 1 II d,+ dA (2-45)

and


a aA .TSV (2-46)
aA

For a crystal which is freely grown from its supersaturated vapor (i.e. no

substrate effects on nucleation, growth energetic, or kinetics), the equilibrium

shape is given by:3

al 02 03 a
= --- = = constant (2-47)
1t 2 "3 i,

where ai is the surface energy of the ith face and Ai is the distance from the center

of the crystal to the face. This result implies that a crystal will grow so as to

minimize its surface energy, and in the case of thin film growth on a substrate,

orientations with the lowest surface energies are the most energetically favorable

to grow.

The morphologies which a growing film adopts generally fall into one of the

following three categories.39 Film growth can be classified as two dimensional, in










which case the film completely covers the substrate before another layer is

nucleated, and the film grows at a uniform rate normal to the substrate. Films

can also grow by a three-dimensional growth process, in which case the film atoms

agglomerate into islands and the islands coalesce to form the film. A third option

is for the film to grow by a mixture of two and three-dimensional growth. The

first growth mode, termed Frank-van der Merwe, is characterized by two

dimensional growth of the film. This occurs when the surface energy of the film

is less than that of the substrate and interface:


oia < a + os ra1~e -fi Lrface (2-48)

In the Frank-van der Merwe growth mode, the surface free energy of the system

is lowered by replacing the substrate surface with a different surface which has a

lower surface energy. For most electronic applications, in which it is often

desirable to grow a thin film with the same structure and orientation as the

substrate, two dimensional growth is preferred because the adatoms will tend to

reside in potential minima along the substrate, thus preserving the pattern of the

substrate surface. The relative adatom-adatom or adatom-substrate bonding

energies also influence the film growth mode, with strong adatom-substrate

bonding favoring two dimensional growth. When the surface energy of the film is

high,











afom > ,brae + abtrate filintrce (2-49)

the surface energy is minimized by the formation of three dimensional islands,

and the islands coalesce to form the film. This is known as the Volmer-Weber

growth mode, and is characterized by strong adatom-adatom bonding. The

orientations of these films are less influenced by the substrate than in the case of

two dimensional growth, and generally have more grain boundaries.

The third mode of film growth, Stranski-Krastanov, occurs when


ofia sbre 0 + Obte ilm inteface (2-50)

The Stranski-Krastanov mode contains features of both of the previous growth

modes, and is often characterized by two dimensional growth of the first

monolayer, followed by three dimensional growth of the rest of the film.

Assuming the film growth rate, surface energy, and adatom-substrate bond

strengths are such that two dimensional film growth predominates, the film will try

to minimize Osubstrate-film interface by growing epitaxially on the substrate, thus

eliminating the energy required to create an additional surface at the interface.

However, if the film has a different chemical composition than the substrate, the

lattice parameter of the film will be different and there will be a strain induced in

the film as it attempts to adopt the lattice spacings of the substrate. Once the

elastic limit of the film is exceeded, the strain is relieved via formation of

dislocations at the interface and the film is no longer completely epitaxial with the

substrate.










There are a variety of methods by which the film will attempt to minimize

differences in lattice parameters with the substrate. Starting with the definition of

lattice misfit ( = f)40


lattice misfit (=f)= af b -asub (2-51)
substrate
where anim and asubstrate are the lattice parameters of the film and substrate. It has

been observed that the orientation which a film adopts with respect to the

substrate is primarily a function of lattice misfit, as well as the adatom-film and

adatom-adatom bonding strengths. If the lattice parameters of the film and

substrate are very close (generally < 1%) and both have the same crystal

structure, the film will align itself so the film directions in the plane of the

interface match those of the substrate. For larger mismatches, the film will adopt

an in-plane orientation such that the maximum number of lattice points will be

commensurate with the substrate.41 Theoretical and experimental data indicate

that the unit cell over which the film minimizes misfit and becomes commensurate

with the substrate may exceed 500 A, hence the correlation between film and
substrate directions is not always obvious or simple to deduce. One commonly

observed method of minimizing strains induced by dissimilar surface meshes is to

strain the lattice of the film so as to match row spacings in one of the directions.42

The Nishiyama-Wassermand and Kurdjumov-Sachs matching are shown in figure

2-10. The difference between the two orientations is the Kurdjumov-Sachs










orientation achieves better row matching by rotating the film surface mesh with

respect to the substrate mesh.


Figure 2-10.


Matching of atomic rows with epitaxial configurations of dissimilar
rhombic meshes. The substrate unit cell PQRS is drawn in solid
lines, and the overlayer unit cells in dashed lines. (a) Nishiyama-
Wassermand orientation, with overlayer PABC matching rows
parallel to PR, and overlayer PERF matching rows parallel to SQ.
(b) Kurdjumov-Sachs orientation; after rotating the overlayer mesh
relative to the substrate, overlayer PABC matches rows parallel to
PQ and PS. Reference 42.


As the lattice mismatch continues to increase, elastic strain within the film

exceeds the shear stress limit, and strain energy is released by the formation of

misfit dislocations.43 Misfit dislocations usually consist of edge dislocations with

the glide direction parallel to the interface. The elastic strain which can be

accomodated before misfit dislocations form is a function of the film-substrate

bonding, with strong bonding leading to larger elastic strain accommodation and a









45
reduced tendency to form misfit dislocations. The strain energy per unit length of

an edge dislocation line is given by:


E= Gb In (2-52)
4n(1 -v) r0

where G is the shear modulus of the film, b is the Burger's vector of the

dislocation, v is Poissons's ratio, R and ro denote the outer and inner radii of the

strain field over which the dislocation acts. In bulk materials, R can be

approximated as the distance between parallel misfit dislocations, b/f. In thin
films, the area over which a strain field acts is limited by the film thickness, and if

the film thickness is less than b/f thick, the strain energy of a dislocation can be

substantially reduced.

In films used for electrical applications, one of the most damaging aspects of

misfit dislocations is that they create grain boundaries which degrade the

electrical performance of the films. For pure edge dislocations, the angle between

two adjacent grains, 0, is given by:


sin 0 = b (2-53)
D

where D is the distance between dislocations. Misfit dislocations also reduce the

film-substrate epitaxy. During the initial stages of film growth, islands are highly

mobile and wil try to attain epitaxy with the substrate. If the substrate-film

bonding is weakened by misfit dislocations, there will be a tendency for the

islands to rotate slightly about the ideal in-plane epitaxial positions.52 The extent

of island rotation is given by:








46

0 =fO (2-54)

where 0 is the island rotation, f is the lattice misfit, and 0 is the rotation of misfit
dislocations. The maximum island rotation is obtained when 0 = 1, and hence q
= f. The result of island rotation is that films deposited on substrates in which

the lattice mismatch is severe will contain a much larger number of grain

boundaries than will films deposited on a more closely lattice matched substrate.

Several of the mechanisms responsible for film growth, and the energetic

parameters which dictate the orientation and microstructure of films, have been

presented in this section. Because several of the parameters are interdependent

and are difficult to observe experimentally, it is impossible to provide a

quantitative analysis of the effect of each variable. However, a general

understanding of the parameters necessary to promote two dimensional, epitaxial

growth provides considerable insight into what defects are likely to emerge from

various film processing techniques, and suggest methods by which film quality can

be improved.


Thermally induced stresses


Thermally induced cracking is a problem which plagues thin films deposited

on substrates. There is usually a difference in thermal expansion coefficients

between the film and substrate, so if the temperature of the film is changed after

the deposition, thermally induced stresses will be generated in the film and

substrate. The least desirable way for the film to relieve these stresses is by










cracking, because cracks severely degrade the electrical properties of the films.

Thermally induced stresses have a significant impact on microstructural features

such as microcracking and grain size, and the microstructural processes by which

thermally induced stresses are relieved will be introduced.

Cracks are formed when the strain energy due to stresses in the film is greater

than the energy required to create new surfaces.45 The stress required to

propagate a crack is given by:


oa- 2E (y, ,)lo (2-55)
*ra

where Ef is the Young's modulus for the film, y, is the surface energy of the new

surface created by the crack, yp is the work of plastic deformation per unit area,

and a is the initial crack radius. A key point is that the fracture stress is

considerably lower at points where pre-existing flaws are present, such as at the

film-substrate interface.4

The model frequently used to calculate the stresses induced in thin films by

differences in the thermal expansion coefficients between the substrate and film(s)

begins with the assumption that the strain is entirely elastic.47 The dimensions of

the substrate are presumed to be unaffected by the presence of the film, and the

film is required to alter its shape in order to match the surface dimensions of the

substrate. Shear forces are generated at the film/substrate interface because of

the differing thermal expansion coefficients, and these shear forces cause the

substrate and film to bow slightly (figure 2-11). Assuming there is no plastic










deformation in the film or substrate, nor any slippage at the film/substrate

interface, stress in the film (af) can be obtained by measuring the radius of

curvature of the film/substrate combination:4


S ESt (2-56)
/ 6(1-vs)R t

where t, and tf are the thicknesses of the substrate and film, v, is the Poisson's

ratio of the substrate, and R is the radius of curvature of the film/substrate

combination. There are two basic requirements imposed by continuum mechanics

which this model satisfies.49 First, the sum of the forces per unit width acting on

the film and substrate must be zero. The force per unit width is defined as aiti,

where i denotes either the substrate or one of the film layers. As an example of

how the stresses are distributed in the film and substrate, consider the case in

which the film stresses are tensile (figure 2-11). These film stresses will cause the

substrate to bow slightly, so the top of the substrate (adjacent to the film) will be

in compression. The magnitude of the forces will vary within the substrate, with

the maximum compressive stress observed at the top, and the maximum tensile

stress at the bottom of the substrate. Near the middle, there will be a plane

through which the stress (and therefore strain) is zero, and this is called the zero

strain plane. In the films, the stress is assumed to be constant throughout the

thickness of the each film layer. In this example, the tensile aiti forces contributed

by the film and the portion of the substrate below the zero strain plane must

balance the compressive ait, contribution of the upper part of the substrate,








49






T
R STRESS
-- Tensile Compressive ----







Mi M1


Figure 2-11. Cross-sectional side view of a two-layered film on a substate,
showing how elastic film stresses are accomodated by the substrate.
Reference 47.


between the zero strain plane and the film-substrate interface. The second
requirement is that the sum of moments about the axis which runs through
the zero strain plane must equal zero. Continuum mechanics requires that the
sum of moments induced by various mechanical forces acting on a common axis
must cancel each other in order for the axis to remain stationary. In our case, the
axis running through the zero-strain plane is common to all the forces acting on
the films and substrate.










Several features regarding the stress distribution in a multilayered film are

implicit in the continuum mechanics model. Since the ot, force for each layer of

film is significantly less than oat, for the substrate, each layer of film is required to

adopt the surface dimensions of the relaxed substrate if there is no plastic

deformation at the film-substrate interface, or between any of the film layers. A

consequence of the small ot force of the film relative to that of the substrate is
that the stresses induced in each layer result almost entirely from interactions

between the film and substrate. According to this model, the stress in any given

layer of a multilayered film is independent of the sequence in which the films

were deposited, and the thermally induced stresses in each of the layers is given

by:46


a=- (af -a,)AT (2-57)

where Ef and vt are the Young's modulus and Poisson's ratio for the film, af and

a, are the thermal expansion coefficients for the film and substrate, and AT is the

change in temperature to which the film-substrate combination is subjected. A

key assumption of this model is that the chemical composition of each layer is

uniform. Since the composition is uniform, the thermal expansion coefficient

within each layer is constant, so the dimensions of the top of each layer will be

the same as the bottom. This assumption, coupled with the assumption of no

plastic deformation at any of the interfaces, leads to the conclusion that it is not

possible to reduce the thermally induced stresses in the outermost film by










depositing intermediate layers in which the thermal expansion coefficients

gradually bridge the gap between the substrate and outermost film.

Film stresses can significantly affect microstructural features such as grain size

and porosity in thin films. Using an energy minimization argument, Chaudhari5

showed that tensile stresses can significantly reduce the average grain size of a

film. Because the atomic packing density is lower in a region containing grain

boundaries than in a crystalline region, tensile stresses are reduced in films with

small grain sizes. A quantitative energy minima for the film is given by the

expression:

AEfmn = AEtra AE~ bndY
E 1 1 aa EE~ 1 1 (2-58)
[( -7)(-)+Eo]- r ( )
1-v do d 2 1-v do d

where do is the initial grain size, d is the final grain size, Eo is the initial strain in

the film, and a is a normalized distance parameter used to compare the atomic
density of the grain boundary region with that of a grain. If the boundary region

has the same atomic density as the grain, then a is zero. If the boundary has a
monolayer of atoms missing, then a is 1. P is a geometrical factor used to
characterize the shape of the grains. For grains with a square cross-section, p =
2. y is the grain boundary energy. For films subjected to tensile stresses, there is
some combination of strain and grain boundary energies for which the overall

energy of the film is a minimum. For films subjected to compressive stresses,

there is no overall energy minima, and the total film energy decreases as the grain

size increases.








52
In brittle materials, film cracking is caused by the motion and accumulation of

dislocations, which creates regions of high elastic energy and ultimately fracture.

The film microstructure can significantly alter the magnitude of stress required to

move dislocations and initiate plastic deformation.46,4751 Numerous experiments

have shown that the stress at which plastic deformation is initiated is greater in

thin films than in bulk materials. This is because dislocations are pinned at the

film-substrate interface, thus increasing the stress required to move dislocations

through the film. The minimum stress required to move a dislocation in a film is

given by:4


S=A[ b fv ln(2.6h)] (2-59)
2n(1-vf)h (vf+v,) b

where A is a geometrical constant which accounts for the angle between the

applied stress and Burgers vector, b is the Burgers vector of the dislocation, vf and

v, are the elastic shear moduli of the film and substrate, and h is the film

thickness. For brittle films, the stresses required to move dislocations are

associated with the onset of microcracking. Equation 2-59 shows that the stress

required to move dislocations is inversely proportional to the film thickness, and is

the reason why there is often a critical film thickness which, if exceeded, causes

cracks to form.

Depositing a second layer on top of the initial film significantly increases the

stress in the initial layer. This seems to contradict the elastic continuum model,

which asserts that the stress in a given layer is only a function of the differences in










thermal expansion coefficients between the each layer and the substrate, and is

independent of the other layers. However, the top layer creates an additional

interface which pins dislocations. Hence plastic deformation in the initial layer is

suppressed, and the yield strength is increased.

The elastic continuum model is a popular model for predicting film stresses in

the elastic limit. However, the assumptions of perfect bonding at each of the

interfaces, and chemical homogenity within each layer are not realistic in most

cases. Microstructural features such as chemical bonding at interfaces, grain size,

and film thickness will affect the stress at which the mode of deformation changes

from elastic to plastic, and must be taken into account when interpreting the

stresses required to form cracks.


YBa, jQCQx Film Growth


Shortly after the discovery of superconductivity in YBa2Cu3O.x, the difficulty

of growing stoichiometric YBa2Cu3O.. thin films by conventional methods such as

sputtering52',3 and electron beam evaporation4 became apparent. Depositing a

film with the correct Y:Ba:Cu ratios is a difficult process using these techniques.

Pulsed laser deposition became a popular technique for growing superconductor

films because stoichiometric films could be readily grown from a target with the

same composition.55 6 Pulsed lasers have generated stoichiometric films over a

wide range of pulse energies, laser wavelengths, and pulse durations. Nd-YAG

lasers operating at 1064, 532, and 355 nm and pulse energies ranging from 0.6 to








54
3.0 J/cm2 have generated YBa2Cu3O7. films with the correct stoichiometry.57'859

YBa2Cu3O7. films were also grown using a pulsed CO2 laser (10.6 pm), but a

Y-enriched target was required to produce a stoichiometric film.6 By increasing

the pulse energy density of a CO2 laser, stoichiometric YBa2Cu3O7.x films were

grown from a stoichiometric target, but the film morphology and superconducting

properties were seriously degraded by large globules in the film.61 The best films

have been grown using pulsed excimer lasers which operate at ultraviolet

wavelengths. The main problems associated with laser deposition is the presence

of fragments in the films which are ejected from the target, and nonuniform film

thicknesses.62 Shorter wavelength excimer lasers are popular because the

particulate size is smaller in these films than in films grown with longer

wavelength lasers.57

The ease with which stoichiometric films can be grown, and the presence of

particulates in the films both result from the laser-target interactions. There are

two basic ablation mechanisms: thermal and electronic.63 Thermal processes

result from rapid heating of the target surface and subsequent evaporation and

sublimation from the surface. Evaporation occurs when the laser power density

(Qo) exceeds the minimum power density necessary for evaporation, Q:64


Qo -" PoU( )0. (2-60)

where po is the target density, D is the thermal diffusivity, U is the sublimation

energy, and r is the duration of the laser pulse.








55
Assuming the optical absorption depth (a1) of the target is small compared to

thermal diffusion length, the relationship a(2Dt)o0 > > 1 is valid. The

temperature rise of the target surface layer (where the thickness, t, of the surface

layer heated by the laser is defined as t = (2DT)o), can be approximated by

comparing the energy absorbed during the laser pulse with the thickness of target

surface heated by the pulse:


AT= (1-R) (2-61)
C, p (2Dt)o
where R is the target reflectivity, I is the power density (W/cm2), t is the duration

of the laser pulse, and C, is the specific heat.

Since the thickness over which laser energy is absorbed in the target is

inversely proportional to the absorption coefficient, target materials with high

absorption coefficients will attain higher surface temperatures because the energy

is confined to a smaller volume. The absorption coefficient is also a function of

the laser wavelength (A), and the absorption coefficient for YBa2Cu3O7-. increases

as A decreases.5 Congruent evaporation from a multicomponent target occurs

because the components cannot segregate over a region greater than the thermal

diffusive region (2DT)0s during the time over which the target is irradiated.

Because the temperature within the thermal diffusive region is high enough to

evaporate all the components, the entire region is evaporated and the film

stoichiometry is the same as that of the target.








56
Electronic mechanisms are also operative during laser deposition.66 Photons

with energies greater than the first ionization potential energies (7.726, 5.512, and

6.38 eV for Cu, Ba, and Y, respectively) will excite the target atoms, thereby

breaking bonds and causing ejection of ions. Experiments in which YBa2Cu3O7-.

film morphology and electrical properties of films grown at different wavelengths

showed conclusively that smoother films with fewer and smaller particulates, as

well as lower normal state resistivities and higher J. values, were obtained in films

grown at shorter wavelengths (figure 2-12).57

The smaller particulate sizes which are observed when YBazCu3O7.x films are

deposited using shorter wavelength lasers is largely attributed to the higher

absorption coefficients in YBa2Cu3O7.x targets with decreasing A. The absorption

coefficients are 1.2 x 105, 1.5 x 10s, and 1.7 x 10s cm- at 1064, 532, and 355 nm,

respectively.57 A higher absorption coefficient results in a thinner layer at the

surface into which the laser energy is coupled, thus creating a hotter

plume with finer fragments. However, the improved microstructures which result

from short wavelength radiation cannot be completely attributed to the slightly

larger YBa2Cu3O7.- target absorption coefficients. Strong absorption by

photofragments in short wavelength radiation, and subsequent fragmentation into

smaller particles, probably contributes to the smooth morphology of films grown

at shorter wavelengths.57

A great deal of effort has been made to understand the mechanisms by which

material is transferred from a YBa2Cu3OT., target to the substrate via laser








57
ablation, and to optimize the film growing process. The film thickness across the

substrate varies as cos9", (n 4),67 which indicates the ejected particle distribution

from the target is highly peaked. In addition, the cation ratios in the film are no

longer stoichiometric at lateral distances greater than 20 degrees from the target


50 100 150 200
TEMPERATURE (K)


250 300


Figure 2-12.


Resistivity as a function of temperature for three YBa2Cu307.x films
deposited on (100) SrTiO3 substrates by a Nd:YAG laser and it's
second and third harmonics. Reference 57.


0.5








0.0
0








58
normal.66 Apparently the highly peaked, forward-directed component is primarily

responsible for stoichiometry in the growing film. One explanation for this

behavior is that the laser generated plasma results in rapid evaporation from the

target surface.6 The gas is initially at high pressures because the rate of

evaporation

is greater than the rate at which atoms and ions can leave the target surface. The

plasma expands into the vacuum, creating a supersonic molecular beam. Time-of-

flight measurements indicate that mean kinetic energies for Cu(1), Y(1), and

Ba(1) are 41.3, 43.4, and 47.9 eV, respectively for particles generated by a

193 nm laser.69 Time resolved optical spectroscopy measurements showed that at

oxygen partial pressures less than 1.3x104 Torr, the velocities of neutral and

ionized atomic species, as well as the diatomic species (such as YO, BaO, CuO)

were approximately 106 cm/sec. Increasing the oxygen pressure to 10'2 Tonr

reduced the velocities of the atomic and diatomic species to approximately

5x105 cm/sec.7

Although the ability to deposit a stoichiometric film is an extremely important

parameter for growing superconductor films with low surface resistivities and high

J, there have been other advances in film growth techniques which have

significantly improved the film quality. Optimization of the growth temperature

and oxygen pressure during deposition, as well as the oxygen pressure and cooling

rate after the deposition, have generated YBa2Cu3O7. films with nearly ideal

superconducting properties. Initially, amorphous films with the correct










stoichiometry were grown in vacuum (typically < 10-s Torr) onto unheated

substrates. The films were then post annealed in flowing 02 at 850 900 "C,

thereby forming the tetragonal YBa2Cu307- phase.5

During cooling from the high temperature tetragonal phase, YBazqCuO7.

undergoes large structural changes. Understanding what these changes are, and

how they are affected by temperature and oxygen pressure are critically important

for optimizing growth of YBa2Cu307.x films. The temperature at which

YBa2Cu307T. transforms from the non-superconducting tetragonal phase to the

superconducting orthorhombic phase is dependent on oxygen pressure.7 In 100%

oxygen, the transition occurs near 700 C. The oxygen content and change in

oxygent content as a function of temeprature are shown in figure 2-13 for bulk

YBa2Cu307-. in 100 % oxygen. If the oxygen pressure is lowered to 20% oxygen,

the tetragonal 6 orthorhombic transition temperature is lowered to 670 C, and
in 2% oxygen the transition is depressed to 620 C. Ordering of the oxygen

atoms into one-dimensional Cu-O chains along the <010> direction is the

primary mechanism responsible for the tetragonal orthorhombic transition.
The dramatic increase in the <010>, and decrease in <100> lattice parameters

as the orthorhombic phase is formed are shown in figure 2-14.

The temperature at which the transition occurs affects the kinetics of the

phase change. The transition occurs via a nucleation and growth process, in which

the ordering of oxygen along <010> and lengthening of the <010> lattice

parameter begins at a grain boundary or free surface.7 If the transition










temperature is suppressed by decreasing the oxygen pressure, growth of the

orthorhombic phase will be slow, and rapid cooling of the sample will result in


I
(a) Ox


I
Orthorhom


bic Tetragonal
--->*



Tp


dOx
(b) dox
dT

1 ^ -

j
I'


-- c I I


6.9

6.8

6.7

6.6

6.5

6.4

6.3

6.2

6.1

6.0


400 600
TEMPERATURE (oC)


800


Figure 2-13. Total oxygen content, and change in oxygen content as a function of
temperature for YBa2Cu3O7,.. Reference 72.


200


1000
























3.94
(a) 100% oxygen
3.92
[010] [10]t




S[[100]00]
3.90


S3.88

LU
0 3.86 -

[100]
3.84 -


3.82 1
0 200 400 600 800 1000
TEMPERATURE (OC)



Figure 2-14. The [100] and [010] lattice parameters of YBa2CuO37.x versus
temperature for a bulk sample heated in 100% oxygen. Reference
71.










incomplete growth of the orthorhombic phase. Rapidly cooled YBa2Cu3O7.

samples will be comprised of orthorhombic nuclei surrounded by an oxygen

depleted, tetragonal phase matrix. YBa2Cu307. films prepared by heating an

amorphous film to 850 900 C in oxygen, then slowly cooling through the

tetragonal-to-orthorhombic transition, have To values of approximately 85 "K on

nonreactive substrates such as SrTiO3, and To 75 "K on the more reactive A1203

substrates.55'56 Post-annealed films on (100) SrTiO3 substrates are polycrystalline,

but are preferentially textured with the (001) plane parallel to the substrate.

Major improvements in J, values were accomplished by depositing the films

in-situ at elevated temperatures and controlled oxygen pressures. Epitaxial

YBa2Cu307. films on (100) SrTiO3 and (100) Y-ZrO2 substrates were grown at

substrate temperatures ranging from 500 650 C and 200 mTorr 02, and

subsequently annealing the films at 450-500 C in 760 Torr 02 for 60 minutes.73'74

In-situ films grown on (100) SrTiO3 and (100) Y-ZrO, had Jc values of 5x106

amps/cm2 and 1x106 amps/cm2, respectively, at 77 K; both had To values of 90

K. In addition, films grown in-situ have much lower room-temperature

resistivities (160 /ohmxcm) than post-annealed films (- 1 milliohmxcm). This
behavior was attributed to the predominance of (001) orientation in the in-situ

film, whereas the post-annealed films contained a mixture of (001), (100) and

(010) oriented grains. The microstructures of in-situ films also formed an

epitaxial orientation with the substrate, whereas a "basket-weave" structure was

observed in post-annealed films.75 The basket-weave structure arises from the








63
growth kinetics of the (100) and (010) oriented grains. Both of these orientations

have the <001> axis, which is the slow growth direction, parallel to the substrate,

but the <100> directions are at 900 to each other.

For YBa2Cu30., films grown in-situ, the substrate temperature and lattice

mismatch at the film-substrate interface significantly affect the orientation of the

films.76 Films grown at 640 C on SrTiO3 and LaAIO3 substrates were primarily

(001) oriented, whereas films on MgO, Y-ZrO2, and A1203 substrates, which were

also deposited at 640 C, grew with an (001) orientation. Increasing the growth

temperature to 720 C resulted in (001) oriented films on SrTiO3 and LaA103

substrates. The variations in orientation were attributed to competition between

minimizing the surface energy of the film, which favors (001) orientation, and

minimizing structural coherence at the film-substrate interface during the early

stages of growth. Both SrTiO3 and LaA103 have the perovskite structure, as does

YBa2Cu307, and the lattice matching between these substrates and YBa2Cu30..

is reasonably close. At low growth temperatures, the reduced surface mobility

and possibility of film-substrate coherence is not as energetically favorable when

there is a large lattice mismatch between the film and substrate, which is the case

for YBa2Cu30., films deposited on MgO, Y-ZrO2, and A1203 substrates. Hence

the surface free energies are minimized by incoherent growth of the (001)

oriented grains.

Above 670 C, the tetragonal, oxygen depleted phase of YBa2Cu307. is

stable, and at a typical growth temperatures of 750 C the non-superconducting










YBa2Cu307. phase is formed. At 750 C, the perovskite lattice is

thermodynamically stable at oxygen pressures greater than 150 mTorr.7 Below

this pressure, YBa2Cu30.. decomposes into its component oxides, hence in-situ

growth is dependent on oxygen pressure, and is usually performed at a PO2 of

approximately 200 mTorr (figure 2-15). The tetragonal-to-orthorhombic transition

temperature is a function of oxygen pressure, with a maximum temperature of 700

C. This transition is usually induced by backfilling the vacuum chamber to 10 -

760 Torr 02 after the deposition, and slowly cooling to 450 C. The film is kept

at 450 "C for approximately 30 minutes to insure oxygenation of the Cu(1) atoms,

and ordering of the 0(1) atoms in the <010> direction. In the tetragonal phase,

0(1) and 0(5) sites are randomly occupied by O atoms, whereas in the

orthorhombic phase the 0(1) sites are completely full and the 0(5) sites are

empty (figure 2-16).7'79

YBa2Cu307. films deposited in-situ at 650 750 C have higher Jc values than

post-annealed films. In-situ films have a higher ratio of (001)/(100) oriented

grains, and have better in-plane epitaxy in the <100> and <010> directions. It

has been established that the penetration depth is increased, and Je values are

lowered by weakly coupled grains separated by high-angle grain boundaries or

non-superconducting interfacial phases. Hence the improved electrical properties

of films grown in-situ results from the reduction of weak links at the grain

boundaries.



















TEMPERATURE (oC)
900 800 700 600 500 400
103
YBa2C30Oy Ortho-2
Ortho-1
Tetragonal

S102 -y=6.0 y=6.5
1 \ Sputtering
CL y=6.9
W 0


ILI Y2BaCuO5 + Ablation
I BaCuO2
L BaCuO2 + Sputtering
-1 Cu20
Thermal
C 100

Z Electron
(B- eam/Thermal
S10'1
o C



10-2 I I I
0.8 0.9 1.0 1.1 1.2 1.3 1.4 1.5
1000/T[K-1]


Figure 2-15. Oxygen partial pressure versus temperature plot showing the critical
stability line for YBa2Cu3O7-. at y = 6.0. Reference 77.























YBa2Cu307


0(2)


Figure 2-16. Crystal structure of YBa2Cu3O7., illustrating the CuO chains and
CuO2 planes.













Barrier Layer Technology


Because of the increased performance which could be realized by replacing

metallic interconnects and microstrip lines with superconductors, significant effort

has been expended towards finding deposition techniques which enable

YBa2Cu3O0.x films with high To and J. values to be deposited on silicon and A1203

substrates. Sapphire is an attractive substrate material for microwave applications

because it is relatively inexpensive, mechanically strong, and has a much lower

dielectric loss tangent (tan 6) than other substrate materials (table 2-3).17



Table 2-3. Dielectric properties of substrate and barrier layer materials.
Reference 17.


Material Dielectric Loss
constant tangent
Sapphire (A1203) 9.4 1 x 10-6
Silicon (Si) 12 1 x 10-3
Y-ZrO2 27 6 x10
LaAIlO 25 5.8 x 10-4
SrTiO3 305
MgO 9.65 4 x 104








68
Devices fabricated on Si substrates would benefit from the lower attenuation and

signal distortion which superconducting interconnects can provide. Because

silicon and sapphire are chemically reactive with YBa2Cu307., and attempts to

grow YBa2Cu3O7. films directly on these substrates result in interfacial phases

which damage both the substrate and film," substantial efforts to grow

intermediate barrier layers on the substrate prior to YBa2Cu3O7. film deposition

have been made. In order to understand why some barrier layers are successful

and thus obtain the expertise necessary to design better barrier layer structures, it

is helpful to characterize YBa2Cu3O7. films grown on inert single crystal

substrates. The first class of substrates are materials which are chemically inert to

YBa2Cu3O7 lattice match reasonably well, and have the perovskite structure.

SrTi03 is cubic with a0 = 3.905 at 25 "C,8 and LaA103 is rhombohedral with ao

= 7.586 and and included angle of 9001'. However, LaAO13 undergoes a cubic &

rhombohedral transformation at 430 C, and at 650 C LaA103 is cubic with ao =

3.818 A,82 so in-situ films are grown on a cubic perovskite LaA103 substrate.

YBa2Cu3O7. films grown at 650 *C on (100) SrTiO3 substrates show To at 89 K

and Jc = 7.5x106 amps/cm2 at 77 K. Cross-sectional transmission electron

microscopy indicates the interface is atomically flat and abrupt, and the film does

not contain any secondary phases.83'84'85 The films grew with the <001> direction

normal to the substrate, the films are heavily faulted, with staggered,

discontinuous (001) layers. Close to the substrate, Y or Ba combine with Cu and

O to form perovskite subshells which are epitaxial with the substrate.








69
Perpendicular to the substrate, each perovskite subshell contains both Y and Ba,

but Y and Ba segregate into different domains along the interface. Away from

the interface, the (001) layers become continuous, but very wavy. Increasing the

substrate temperature to 720 780 "C reduced the number of defects. Films

grown at the higher temperatures were completely (001) oriented, with less wavy

(001) layers, and Jc values of 2.2x 106 amps/cm2 at 77 K.

YBa2Cu3O7-. films with To = 90 K and J, = 1x106 amps/cm2 have been

deposited on (1102) LaAlO, at 750 C.86 These films were highly (001) oriented

during the first 4000 A of their growth. At greater thicknesses however, the
growth mode abruptly changed, and the dominant growth mode was with the

(100) and (010) planes parallel to the substrate. Computer simulations based on

nucleation densities and growth rates of (001) and (100)-oriented grains help

explain the experimentally observed film microstructures.87 Assuming the

nucleation density of (001)-oriented grains is 109/cm2, and for (100) and (010)

oriented grains is 2x108/cm2, and the growth rate in the <100> and <010>

directions is 10 times larger than in the <001> direction, a microstructure

emerges in which the film surface is initially dominated by (001)-oriented grains.

As the film thickens, the (100) and (010)-oriented grains, which have a high

growth rate normal to the substrate, coalesce and the film is covered by these

orientations.

YBa2Cu307.x films have also been deposited on several substrates which have

distorted perovskite structures. The difference between these substrates and the








70
cubic perovskite materials (such as SrTiO3) is that the angle between the <100>

and <001> directions is not 90 0. YBa2Cu3O7- films with To values of 92, 89,

and 88 "K have been deposited on single crystal LaGaO3," PrGaO3,89 and

YbFeO39 substrates, respectively. In each case, the YBa2Cu307-. films were

strongly (001) oriented. Although each of these substrates contain elements which

are known to suppress superconductivity (La, Pr, and Fe), interdiffusion was not

observed.

Despite a large lattice mismatch, highly (001) oriented YBa2Cu307.x films with

To = 92 K have been deposited on LiNbO, substrates.91 The unit mesh of Y-cut

LiNbO3, onto which the YBa2Cu30-.x films were grown, is 5.148 x 6.932 A. The J.

values for these films were 2x10s amps/cm2 at 77 K, and the reduction in Je

(relative to YBa2CuO7.x films deposited on SrTiO3) was attributed to the high

concentration of Li in the film. The diffusion of Li was so rapid that the

concentration of Li within the YBa2CU3O7-. film film was peaked at the surface.

The second group of substrates are materials which are chemically inert to

YBa2Cu3O7., but do not lattice match nor do they have the perovskite lattice. Y-

ZrO2 amd MgO have been widely studied because high quality YBa2Cu3O7. films

have been grown on these materials. Despite the large YBa2Cu3O7.x/Y-ZrO2

lattice mismatch of 5.95 % in the YBa2Cu3OT7. direction,

YBa2Cu3O7 films with To = 88 K and J, = 1x106 amps/cm2 at 77 K were

deposited on (100) Y-ZrO2.92'93 Tietz et. al.94 proposed a model for










YBa2Cu307O. growth on Y-ZrO2 substrates which asserts that YBazCuO37. grows

in a manner which permits matching of the oxygen sublattices. This model also

proposes that the large lattice misfits are accomodated by 90 degree boundaries

and stacking faults. Norton et. al.92 reported that YBa2Cu3O7. films with To = 90

K and Jc = 11,000 amps/cm2 at 77 K were grown on a polished, randomly

oriented Y-ZrO2 substrate at 680 C. The films were strongly (001) oriented, and

the electrical properties of the films were more sensitive to the substrate

temperature during growth than were films grown on SrTiO3 and LaAlO3.

Increasing the growth temperature to 730 C resulted in YBa2Cu3O7. films with

Jc values of 1000 amps/cm2. The drop in Jc was attributed to increased chemical

interaction at the Y-ZrO2/YBa2Cu307. interface, with subsequent formation of

BaZrO3. Presumably, BaZrO3 diffused through the grain boundaries and reduced

intergranular conduction.

The lower Jc values for YBa2Cu307. films on polycrystalline Y-ZrO2 relative

to single crystal Y-ZrO2 was attributed to the presence of high angle grain

boundaries. Garrison et al.95 demonstrated that by altering the deposition

conditions, YBa2Cu3307 films could be deposited such that matching of the

YBa2Cu307-x <100>/Y-ZrO2 <100> or <110> directions could be induced. When one

or the other of these orientations was dominant, YBa2Cu3O7. films with Je = 106

amps/cm2 at 77 K were observed. However, if both orientations were present in

the same film, the Jc values were only 102 104 amps/cm2 at 77 OK. This










behavior was attributed to the high angle grain boundaries which resulted from

the mixed in-plane orientations of the YBa2Cu307.x films.

Films grown on (100) MgO substrates also showed To values of 89 K and J,

= lx106 A/cm2 at 77 OK.9697 Similar to Y-ZrO2, there is a large YBa2Cu3O-.x/

MgO lattice mismatch (5.9 and 4.2% in the YBa2Cu3OTx <1o> and <01o0/MgO,0o>

directions, respectively). YBa2Cu3O7-x films grown onto MgO at 670 C were

predominantly (001) oriented. Unlike Y-ZrO2, there was no evidence of

interfacial reactions at the YBa2Cu3O/7MgO interface. Similar to the case for

Y-ZrO2, it was speculated that YBa2Cu3O7., grows on MgO in a manner which

permits matching of the oxygen sublattices. The model also proposed that the

large lattice misfits are accomodated by 90" boundaries and stacking faults.

The third group of substrate materials is characterized by materials which

chemically react with YBa2Cu307x, thus making growth of high quality

YBa2Cu3O7.x directly onto these substrates very difficult. Unfortunately, the two

substrate materials most widely used for electronic applications--(100) silicon for

integrated circuits and A1203 (sapphire) for micro and millimeter-wave electronics,

belong to this group. A 1.5 pm thick YBa2Cu3O7-x film grown in-situ at 700 C

onto a Si substrate showed To = 70 K, and an interfacial reaction layer 0.5 pm

thick was observed.98 Auger electron spectroscopy detected Si at the film surface,

indicating Si diffusion through grain boundaries and microcracks. The poor film

quality was also attributed to microcracks generated by the large difference in

thermal expansion coefficients for Si (3.8x10-6/C) and










YBa2Cu307. (13.6x 10-6/C). Chourasia et al.99 showed that in a YBa2Cu3O7.

film deposited directly onto a Si substrate at room temperature, Si diffused into

YBa2Cu3O7, near the interface, and formed a Si suboxide which depleted oxygen

from the CuO planes. After annealing the film at 860 C for 3 hours, the Si

diffusion was much more extensive, and SiO2 was detected at the surface. The

authors concluded that oxidation of Si at the expense of CuO was the primary

reason that diffusion of Si into YBa2Cu307, degraded the superconducting

properties of YBa2Cu30,7.

Better films have been grown directly on sapphire. YBa2Cu37.,, films grown

on (1102) A1203 at 700 C showed To = 87 "K and JC values of 2x106 amps/cm2

at 4.2 OK.1 Attempts to deposit films at higher temperatures resulted in rapid

deterioration of the film because of interfacial reactions, probably BaO and CuO

reacting with Al203.

Because of the enormous technological advances which would result from

successful deposition of high quality YBa2Cu30.x films on Si and sapphire, a great

deal of research has been dedicated towards overcoming the interfacial reaction

problem. The most common approach has been to grow an intermediate barrier

layer prior to YBa2Cu307.-. The principal requirements of the barrier layer are

that it must be chemically inert to both the substrate and the YBa2Cu3O7. film,

and should lattice match YBa2CU3O7. and the substrate. The materials which

have been most successful as barrier layers have been the materials which are also

the best substrate materials, such as SrTiO3, Y-ZrO2, and MgO. Although








74
excellent quality YBa2CuO37., films have been deposited on single crystal LaAO03

substrates, LaA103 has not been reported to be a good barrier layer material for

growth of YBa2CU307. on either A1203 or Si substrates. Attempts to grow

LaAO03 films on a variety of substrates at 760 C showed that epitaxial films

could be deposited on SrTiO3 and LaAO03 substrates, while attempts to grow

LaAIO3 films on Si, A1203, and MgO substrates resulted in amorphous films.10

This indicates that matching both the lattice parameters and crystal structures at

the film-substrate interface are important parameters which critically influence the

crystallinity and orientation of the YBa2Cu3O7.x films.

Experiments in which LaAlO3 and YBa2Cu3O07. were simultaneously deposited

onto an MgO substrate showed that To gradually dropped as the fraction of

LaAO03 in the film increased.t10 YBa2Cu3O., films which contained no LaA103

had a To = 87 "K, while the transition temperature dropped to 77.5 and 30 K

for YBa2Cu3O7. films containing 9 and 13 mole percent LaAlO3, respectively.

When the transition from Tsen to To was determined by measuring the magnetic

response of the film to a magnetic field (inductive response), the transition

remained fairly sharp (< 5 K) for all the films. However, when measured by the

electrical resistance technique, the transitions became much broader as the

fraction of LaA103 in the films increased. Apparently, the formation of LaA103

at the grain boundaries decreased coupling of the superconducting wave function

between the grains. The discrepancies in the widths of the temperature ranges

over which the transitions occurred was attributed to the higher sensitivity of the








75
inductive measuring technique to intragranular conductivity, whereas the electrical

resistance technique is more sensitive to the presence of intergranular defects.

To date, Y-ZrO2 has been the most successful barrier layer for YBa2Cu307-x

films grown onto Si substrates. 1 pm thick YBa2Cu3O0.x films with To = 82 K

have been grown at substrate temperatures of 650 C, using 0.1 pm Y-ZrO2

barrier layers.83 The orientation of the Y-ZrO2 layer greatly influences the

orientation of YBa2Cu3O.7-, and the best superconducting films are grown onto

(100) oriented Y-ZrO, barrier layers.1" Fork et al. showed that the ratio of

(200)/(111) x-ray diffraction peaks from Y-ZrO2 is strongly influenced by the

depostion temperature, and to a lesser extent, the oxygen partial pressure during

growth. The conditions under which highly (100) oriented Y-ZrO, barrier layers

were grown on Si were at a substrate temperature of 780 "C and Po2 = 7 x 10

Torr. Growth of (100) Y-ZrO2 was also promoted by degreasing and cleaning the

silicon substrates in a flowing N2 hood, then transferring the substrates to the

deposition chamber via a N2 purged glove box, thus insuring a hydrogen

terminated Si surface. Thin YBa2Cu3OT7. films (305 A) deposited at 750 TC, and

grown on 500 A Y-ZrO2 layers have To = 86 88 oK, and Jc values of 2.2x106

amps/cm2 at 77 K. These values are comparable to those obtained from

YBa2CuO37x films deposited on SrTiO3 substrates. However, increasing the

YBa2Cu307-. film thickness to 1300 A reduced the Jc to 1.5 x 10 A/cm2 at 77 K.

Table 2-4 lists some of the more successful YBa2CU307.x/barrier layer structures

deposited on Si substrates, along with the film thicknesses. These data show that










reductions in To and J, as the YBa2Cu3O7-. film thickness increases result from

film cracking caused by tensile stresses in the YBa2Cu3O7-. films. The cracking

seriously degrades the superconducting film properties as the YBa2Cu3O7- film

thickness exceeds approximately 500 A.
In addition to the stress created by the thermal expansion coefficient

differences between Si and YBa2Cu307., stresses in the YBa2Cu307, films are

exacerbated because the thermal expansion coefficient of YBa2Cu307. is highly

anisotropic.14 Table 2-5 tabulates the thermal expansion coefficients of

YBa2Cu3O7.x in three directions, and in different temperature regimes. The

differences in thermal expansion coefficents in the <100> and <010> directions

are caused by ordering of oxygen atoms along the <010> direction, and

formation of the orthorhombic phase. The largest thermal expansion coefficients

are observed in the <001> direction. For YBa2Cu3O7T. films in which the <001>

direction is normal to the substrate, nucleation of a grain with the <010>

direction along a given direction parallel to the substrate will also result in

nucleation of another grain with the <010> direction orthogonal to the first

grain. This is the method by which the system minimizes the stresses caused by

differing thermal expansion coefficients in the <100> and <010> directions. If

an (001) oriented YBa2Cu3O7-. film is deposited on a substrate (such as SrTiO3 or

LaA103) in which the thermal expansion coefficient is close to the average

thermal expansion of YBa2Cu3O7-. in the <100> and <010> directions, film

cracking does not appear to be a problem. However, for substrates with lower








77
thermal expansion coefficients (Si and A1203), YBa2CuO37.x film cracking caused

by thermally induced stresses is a significant problem.

The cracking problems caused by the low thermal expansion coefficients of Si

have been circumvented by using a silicon-on-sapphire structure.105 YBa2Cu307-.

films grown onto a Y-ZrO2 barrier layer, which in turn was deposited on a Si film

grown on a sapphire substrate, had Jc values of 4.6x 106 A/cm2 at 77 "K for

YBa2Cu307.- film thicknesses up to 4000 A.

Y-ZrO2 has also been successfully used as a barrier layer for YBa2Cu3O7.x

films deposited on the (1102) plane of A1203. Highly (100)-oriented Y-ZrO2 films

were deposited on (1102) A1203 at substrate temperatures greater than 780 "C,

whereas growing the Y-ZrO2 layer at lower temperatures increased the ratio of

(111)/(100) Y-ZrO2 x-ray diffraction peak intensities.36 YBa2Cu3O7.. films

deposited on highly (100) oriented Y-ZrO2 barrier layers had To = 90 K, and J.

values of 1.2x 106 amps/cm2 at 77 K.

SrTiO3 has emerged as a very good barrier layer material for growth of

YBa2Cu307-. on (1102) A1203 substrates. 1.2 pm thick YBa2Cu307-. films grown

on a 4000 A SrTiO3 barrier layer, had To values of 86.5 K and Jc values of x 106

amps/cm2 at 77 K.106 X-ray diffraction indicated that the SrTiO3 layers

preferentially grew with a (110) orientation, although the (200) peak was

significant. The YBa2Cu3O7-, films were (001) oriented. Secondary ion mass

spectroscopy showed a drastic reduction of Al concentration in films grown on









Table 2-4. Superconducting transition temperatures and critical current densities for YBa2CU3O7dbarrier layer films
on silicon and LaAO03 substrates.


Substrate Barrier layer and YBazCu30,x Transition Critical current Reference
Thickness (A) thickness (A) temperature density
(K) (amps/cm2)
Si Y-ZrO2 (1000) 10,000 82 -- 83
Si BaTiO3/MgAl204 --- 70 -- 107
(3500/5000)
Si BaTiO3/MgAl204 1000 86-87 6x104 at 77 oK 108
(3500/750)
Si Y203/Y-ZrO2 600 82-84 1x106 at 77 OK 109
(100/900)
Si Y-ZrO2 130 86-88 2x106 at 77 OK 103
(500)
Si Y-ZrO2 1350 --- 1x10 at 77 OK 103
(500)
LaAO03 --- 1300 88-90 5x106 at 770K 73











Table 2-5. Thermal expansion of YBa2Cu307Ox. Reference 104.


Thermal Expansion (xl06/OC). Dila-
tometer
<100> <010> <001> Average average

Orthorhombic
25-400 oC 14.3 5.8 25.5 15.2 12.9
400-610 oC 37.5 0.0 39.5 25.7 25
25-610 oC 22.6 3.5 30.3 18.8 16.6
Tetragonal
25-800 oC 11.5 17.0 13.3 10.9


SrTiO3 barrier layers relative to superconducting films grown directly on A1203,

confirming that SrTiO3 is an excellent barrier to Al diffusion. Char et al.o10 found

that growth of the YBa2Cu3O7/SrTiO3 film on (1102) A1203 at 750 C produced

films with To = 86.5 "K and Jc values of 2x 106 amps/cm2 at 74 K. Higher

YBa2Cu3O7., deposition temperatures and hence better in-plane epitaxy, which

were made possible by the SrTiO3 barrier layers, were credited as the cause for

improved electrical properties, relative to YBa2Cu30. films deposited directly

onto (1102) Al203.













CHAPTER 3
EXPERIMENTAL TECHNIQUES



Film Growth by Laser Deposition


Barrier layer and YBa2Cu3O7. films were sequentially deposited at 730 750

C on silicon (Si), aluminum oxide (A1203), yittria-stabilized zirconia (Y-ZrO2),

lanthanum aluminate (LaAlO), or strontium titanate (SrTiO3) substrates using a

pulsed laser deposition system. A Questek model 2560 pulsed excimer laser,

using KrF gas and operating at 248 nanometers, 30 nanosecond pulses, and 5

pulses/second, was focused to 2.5 3.0 Joules/cm2 with a 50 centimeter focal

length lens onto a one-inch diameter YBa2Cu3O7. or barrier layer target. A

schematic diagram of the deposition system is shown in figure 3-1. The

stoichiometric YBa2Cu307 target was obtained from Ceracon, Inc. and was 96%

dense. The barrier layer targets were fabricated by mixing stoichiometric ratios of

the powders, calcining at 950 OC, then pressing the powders into disks and re-

firing at 950 OC for 12 hours. The barrier-layer targets were approximately 65%

dense. Up to three targets could be mounted on a stainless steel holder. By

rotating the target holder, sequential films were deposited without breaking











PULSED LASER DEPOSITION


Excimer Laser
248 nm, 30 ns pulses


Figure 3-1. Schematic diagram of the pulsed laser deposition system.









vacuum or reducing the substrate temperature. This holder did not allow

continuous rotation of the target during deposition, but the target was moved

slightly every 800 pulses to expose a new surface to the laser radiation. The

deposition temperatures were measured by a thermocouple spot-welded to the

heater block. The barrier layer films were approximately 1000 A thick, and unless
specified otherwise, were grown in 40 mTorr 02. The YBa2CCu307 films were

2000 3000 A thick, and were always deposited at Po2 = 200 mTorr. After the

YBa2Cu30O7. films were deposited, the chamber was filled with oxygen, and the

temperature was maintained at 730 750 C for 20 minutes in order to facilitate

the tetragonal-to-orthorhombic transition. The films were cooled to 450 C over

60 minutes, held at 450 C for 45 minutes at approximately 300 Torr of oxygen to

ensure complete oxygenation, then cooled to room temperature.

All of the substrates used in this study were single crystal. The SrTiO3 and Si

substrates were cut parallel to the (100) planes, LaAO03 was cut parallel to the

(1102) planes, and Al203 was cut parallel to the (1102), (1210), or (0001) planes.

The Y-ZrO2 substrates were cut 5 12 degrees from the (100) planes, as

determined by Laue back-diffraction patterns. By depositing the YBa2Cu307.

films on off-axis Y-ZrO2 substrates, we increased the tendency for high-angle

grain boundaries to form in the barrier layer and superconducting films. This

feature enabled us to determine whether the barrier layers would passivate the

YBa2Cu307.x films from the defects introduced by these substrates, and allow

growth of superconducting films with high Jc values.










A variety of techniques were used to characterize the superconducting and

barrier-layer microstructures. Film orientation and interdiffusion at the barrier

layer/substrate and YBa2Cu307.,/barrier layer interfaces critically affected the

electrical properties of the YBa2Cu307. films, and evaluation of the interfacial

reactions and diffusion phenomena which promoted various types of

microstructures were required in order to correlate film microstructure with

electrical performance. Several analytical techniques, including x-ray diffraction

(XRD), scanning Auger electron spectroscopy (AES), scanning electron

microscopy (SEM), and Raman spectroscopy were used to evaluate the film

microstructures. Because the beam/sample interactions and detection techniques

were different for each of the measurement techniques, various microstructural

features could be examined. By understanding the mechanisms by which the data

was generated, the sample volume which was probed by each technique, and the

factors which were likely to reduce the validity of the data (such as electron

charging in AES), a complementary set of data were obtained which uncovered

many of the microstructural features which influenced the superconducting

properties of the films. Similarily, electrical and magnetic measurements provided

essential information about the film microstructures, and the suitability of the

YBa2Cu3O7-./barrier layer/substrate combinations for various devices. A basic

understanding of how the techniques work, and the potential sources of error are

essential in order to assess the data. In this section, a description of the various










experimental techniques, and the regimes in which they were used for this study,

is presented.


X-ray Diffraction


X-ray diffraction was initially used to verify that YBa2Cu307-. and barrier layer

films were being grown. Interfacial phases were also detected using x-ray

diffraction (XRD). A Phillips model APD 3720 x-ray diffractometer operating at

40 kilvolts and 20 milliamps was used to generate Cu Ka radiation of A = 1.54060
and 1.54439 1. A graphite monochromater filtered out most of the Cu 1kf
radiation. Peaks within the 20 range of 5 65* were detected, and the x-ray
detector was rotated at 3" per minute. Interplanar spacings were calculated using

the Bragg equation:"1


nAL = 2dsine (3-1)

where n is an integer, A is the photon wavelength, d is the interplanar spacing,
and is the angle (relative to the sample surface) at which the x-rays enter and

leave the sample. By matching the experimentally observed interplanar spacings

with those predicted by the Joint Committee of Powder Diffraction Standards

index, the phases and orientations of the films were determined. Although the

graphite monochromater eliminated over 99% of the Kfl radiation, samples which
produced extremely large Ka diffraction peaks, such as the single crystal
substrates, also produced measurable Kf x-ray diffraction peaks.
The orientations of the YBa2Cu3O7.. films were highly dependent on the

orientations of the barrier layer films. To determine whether phase information










about the barrier layer or interfacial phases buried beneath the superconducting

film could be obtained, the x-ray attenuation depths for the various films were

calculated. Assuming the x-ray intensity decreases exponentially as it enters the

sample, the attenuation for each compound was calculated using the expression:


S= exp[-() pt] (3-2)
o0 P

where I/I0 is the fraction of the incident x-ray intensity which penetrates to a

depth = t, (u/p) is the mass absorption coefficient for each phase, and p is the
density of the phase. (u/p) was calculated from the weighted fractions of the
mass absorption coefficients for the individual elements:


(!),I = wl([) + W2( ) + W3( ) + ...+ (1) (3-3)
P P P2 P3 Pn

where Wn is the weight fraction of element n in the phase, and (u/p)n is the mass

absorption coefficient for element n. For the calculation, the penetration depth at

which I/Io was equal to 0.368 was taken to be the absorption depth (= t). Of the

films examined in this study, YBa2Cu30. had the smallest absorption depth (= 9

/im). Since the YBa2Cu3O7x films were typically 3000 A thick, and the barrier

layer films were 1000 A thick, we concluded that diffraction data was obtained
from all of the films in the multilayered structure. The calculation of the x-ray

absorption depth for YBa2Cu30. is presented in appendix A.













Scanning Electron Microscopy


A JEOL JSM 35C scanning electron microscope (SEM) operated at 15 20

kilovolts accelerating voltage and 100 microamps beam current was used to

visually determine the surface morphology and microstructures of the films.

Microstructural features are readily imaged in the SEM because secondary

electron detection is highly sensitive to the angle between the ejected electrons

and detector, so the number of secondary electrons detected varies as the

topography of the film changes.112 YBa2Cu30 films deposited on highly reactive

substrates or barrier layers tended to be cracked or have rough surfaces, whereas

superconducting films grown on inert substrates were smooth and featureless.

Correlations between SEM micrographs and electrical data were instrumental for

clarifying the types of microstructures which resulted in YBa2Cu307-. film

degradation.


Scanning Auger Electron Spectroscopy


Film uniformity and interdiffusion between the various layers and Si substrates

were analyzed using Auger electron spectroscopy (AES). A Phi model 660

scanning Auger microprobe, controlled by an Apollo domain series 3500

computer, was used to determine the film composition as a function of depth.

The operating parameters for the Auger were 5 kilovolts accelerating voltage, 25 -










35 nanoamps beam current, 132 volts emission voltage, and 40 60 microamps

emission current. Interdiffusion at the YBa2Cu3O.7x/barrier layer and barrier

layer/substrate interfaces was observed by ion sputtering a crater in the films, so

as to expose the barrier layer and substrate, then making a line scan across the

edge of the crater. With this method, errors induced by electrical charging and

Auger peak shifting in the insulating barrier layers were minimized.

Auger is often used to qualitatively measure the relative concentrations of

components within the top 10 A of the surface. The type of element is
determined by the energy of electrons emitted as a result of the Auger process.

The Auger process is started by an incident electron beam with sufficient energy

to remove an inner shell electron, which creates a core hole. The ion energy is

reduced by filling the core hole with an electron from a more shallow energy

level, and emitting another electron from a shallow energy level. The energy of

the emitted electron is given by:13


Kinetic Energy = EA E- Ec (3-4)

where EA is the energy of the core level electron, EB is the energy of the shallow

level electron which fills the core hole, and Ec is the energy of the shallow level

electron which is emitted. Although EA, Eg and Ec are all sensitive to the

chemical state of the atom, the time constant for Auger emission is short, so the

peaks are broad. Therefore the energies of electrons emitted via the Auger

process are less sensitive to the chemical state of the element than are electrons

emitted by other techniques, such as x-ray photoelectron spectroscopy. Hence








88
AES is widely used to determined which elements are present at the surface, with

limited determination of chemical state.

Quantification of the surface concentration is a difficult process because there

are many factors which influence the Auger yield. For an Auger transition from

species i at a site (x,y,z), where Ni is the background Auger count and dN, is the

number of Auger electrons resulting from the transition:114

dN, = (incident electron flux of energy Eprimary at xy,z)

x (ionization cross-section of EA for species i at E,)

x backscatteringg factor for Eprimary at the incident direction)

x (probability of decay of EA for species i to give the Auger

transition)

x (probability of no loss escape of electrons from region (x,y,z))

x (acceptance angle of analyzer)

x (instrumental detection efficiency).

To as much of an extent as possible, the Auger operating parameters were

kept constant in these experiments so that comparisons between the atomic

concentrations of different samples could be made. The incident electron flux was

dependent on the beam current, and was maintained at 30 40 nanoamps. The

ionization cross section is heavily influenced by the incident beam energy; low

beam energies are not adequate to produce core holes, and high beam energies

reduce the Auger yield from the shallow core levels. Generally, the optimal beam

energy is 3 5 times the binding energy of the deepest core level of interest. In










this set of experiments the beam accelerating voltage was always 5 kilovolts

because the sensitivity of the YLMM transition is highest at this accelerating

voltage.

Atoms can be ionized by backscattered and secondary electrons as well as

primary electrons, and the backscattering and secondary-electron yields are

sensitive to the chemical environment and electrical properties of the sample.119

Discrepancies in ionization cross section resulting from different backscattering

yields of the same element in different compounds is the primary phenomena

which limits quantitative Auger analysis. In this set of experiments, interdiffusion

between YBa2Cu3O7-, the barrier layers, and the Si substrates were of interest.

Since the chemical environment and backscattering yields of each of the structures

were similar, semi-quantitative comparisons between the chemical compositions of

these structures could be made. Two final parameters which significantly effect

the Auger yield are the angle between the sample surface and the incident

electron beam, and the angle between the surface and detector. As the angle

between the beam and surface normal increases, the incident beam path length in

the surface region increases by a factor of sec 0,14 and the Auger yield increases.

Experimentally, this factor was kept constant by always keeping the angle between

the incident beam and surface normal at 60 .













Electrical Resistance Measurements


Electrical resistance versus temperature data were taken in order to correlate

normal state resistances, the onset of superconductivity (Tont), and the

temperatures at which the DC resistance dropped to zero (To), with YBa2Cu3O7.x

film microstructures. Resistance measurements were obtained using a four-point

probe apparatus in which current was transported through the film by the outer

two terminals, and the voltage drop was measured across the inner two terminals.

The four probes were mechanically pressed against the sample, and either 10 or

100 microamps were transported through the sample. To was determined when

the voltage drop was less than 1 microvolt (R < 0.1 or 0.01 ohms. Resistance vs.

temperature data were also obtained on many of the samples at NASA Lewis

Research Center. The system at NASA was more sensitive than the one at the

University of Florida; gold contacts leads were wire bonded directly to the

superconducting films, which increased the sensitivity in the low resistance regime

near the transition temperature. The criteria for superconductivity in this system

was R < 0.001 ohms. Despite the differences, both measurement apparatuses

produced very similar normal state resistance and transition temperature data

when the same samples were tested on both systems.

For thin films, resistivity rather than film resistance is usually plotted because

resistivity is a material property, and the calculations used to obtain resistivity








91
values take into account film thickness, probe spacings, and sample geometry. In

this set of experiments, resistance was documented because the variations in

resistance caused by microstructural features overshadowed the relatively minor

changes in resistivity caused by varying YBa2Cu3O7. film thicknesses and sample

geometries. An explanation of how resistivities are calculated, and how they are

affected by sample geometries is presented in order to support the hypothesis that

electrical resistance was the more appropriate parameter to monitor. Film

resistivity is given by:


p =FRt (3-5)

where F is a correction factor, R is the measured resistance, and t is the film

thickness. Assuming the probes are equally spaced, there are two correction

factors which will improve the accuracy of the resistivity measurements."1 The

first correction factor is given by the thickness correction factor, F(t/a). As the

sample length, a, becomes significantly greater than the film thickness, F(t/a)

approaches 1. Since the films were less than 3000 A thick, and the length of the
substrates were usually greater than 1 cm, F(t/a) had a negligible influence on the

measurement. The second correction factor, F2, is a geometric correction factor

which takes into account the increased current densities in the sample caused by

narrow samples with relatively large distances between the probes. Films

deposited on narrow substrates have large a/d ratios, where d is the sample width.

If we assume the YBa2Cu3O7. films were deposited on rectangular substrates in

which the length was twice the width (a/d = 2), the appropriate correction factors










for the film resistivity as the ratio of sample width to probe spacing (d/s) are

presented in table 3-1.

There were several practical difficulties which made calculating the

resistivities difficult. First, there was a trade-off between growing the highest

quality films and accurately assessing the film thickness, because masking off an

area of the substrate to create a step for profilometer analysis created thermal

gradients in the sample, which damaged the superconducting film quality. Second,

the substrates on which the films were deposited had a variety of shapes, so the

precise geometrical correction factor was difficult to establish. Although the

geometric factors were diverse, it is unlikely that they altered the resistivities by

more than +/- 50%, since the sample width/probe spacing ratios were

constrained by the design of the R vs. T measurement apparatus to be between 2

and 3. On the other hand, measured resistances often varied by an order of

magnitude or more, depending on the film microstructure and type of substrate

used. The additional information which could have been obtained by determining

the film resistivities would have been minor, and the fundamental

microstructural/electrical property correlations which controlled the electrical

properties were more fully uncovered by correlating data obtained by electrical

resistance measurements with other types of microstructural data.












Table 3-1. Resistivity correction factors as the sample width to probe spacing
increases. The sample length to width is kept constant (a/d = 2).


Ratio of sample width Resistivity correction
to probe spacing (= d/s) factor (= F2), assuming
a/d = 2.
1.50 1.4788
1.75 1.7196
2.00 1.9454
2.50 2.3532
3.00 2.7000
4.00 3.2246
5.00 3.5746
7.50 4.0361
10.00 4.2357
15.00 4.3947
20.00 4.4553
40.00 4.5129
00 4.5324




Full Text
20
T,, N 4ttc2 %
K(g,a>)= :-J1-h
q2 + k2
o)2-o)2q
].
(2-19)
where V (q, between the k and k* states, aq is the phonon frequency at wave-vector q, and k is
the wave number of superconducting electrons at the Fermi surface. During an
electron transition between two states (k to k'), a phonon may be absorbed. If (o
< ft)q, and the transition lowers the potential energy of the system, the electron-
phonon interaction creates an energy gap in the energy vs. momentum spectrum,
and the most favorable way for two electrons with p > Pf (p = momentum at the
Fermi level) to lower their energy below 2Ef is to form a bound state, in which
electrons with equal and opposite momenta combine to form Cooper pairs. The
wavefunction for a Cooper pair is given by:4
- (M0)
p h
where W is the amplitude of the wavefunction (also known as the "order
parameter"), multiplied by the travelling wave expression in which P = electron
momentum, r = position, h is Planck's constant, and | *P |2 is the density of
superconducting electrons. Cooper pairs obey Bose-Einstein statistics, hence it is
energetically favorable for all the pairs can have the same momentum. Because
all the superconducting pairs have the same momentum, and thus the same
wavelength, superposition of the coherent waves results in another wave with the
same wavelength. Superconductors possess macroscopic quantization because all


INTENSITY (ARBITRARY UNITS)
127
Figure 4-22. X-ray diffraction pattern from a YBa2Cu307.x/Zr02 film deposited
on Y-Zr02.


233
grains. The lack of x-ray diffraction from these non-(OOl) peaks indicates that the
grain size was too small to yield diffraction peaks. The small size of the 500 cm'1
Raman peak from the YBa2Cu307.x/LaA103/Al6Si2013 film shows that the fraction
of non-(OOl) oriented grains is much smaller, but the wide breadth of this peak
indicates disorder along the oxygen chains or substitution of Cu(l) by impurity
atoms (such as A1 or Si). Both types of phenomena could have changed the
bonding environment of the 0(4) atoms, and cause the spread in the 500 cm'1
peak. According to this data, intragranular defects should have been a significant
cause of degradation in the YBa2Cu307_x/LaA103/Al6Si2013 film on Si. Except for
the * 2x difference in the normal state resistance values for the YBa2Cu307.x/
(YA103 vs LaA103)/Al6Si2013 films on Si, the resistance versus temperature
patterns were very similar. Using the Halbritter equation to uncover the
predominant type of defect responsible for the electrical behavior, the normal
state resistances were dominated by cracking, and the cracking was less
pronounced in the YBa2Cu307.x/LaA103/Al6Si2013 film. However, the Raman
spectra shows the intragranular microstructures were quite different for the two
films.
Comparisons between YBa2Cu307.x/(YA103 or Y203)/Al6Si20i3 films
deposited on (1102) A1203 provided insights into the role which the mechanical
properties of the barrier layers play in the overall degradation of the YBa2Cu307..x
films. Since the thermal expansion coefficient of Al6Si2013 is smaller than that of
(1102) or any of the other films used in this study, prediction of how the Al6Si2013


22
1.0
0.8
<
£ 0.4
0.2
0
T/Tc
Figure 2-7. The superconducting energy gaps of lead, tin, and indium versus
temperature. Reference 4.
and shows the gap is a maximum at T = O K. To illustrate the fraction of the
ideal superconducting bandgap which commercial high temperature
superconducting devices will be expected to operate at, we assume T0 to be 90
K, and the device operating temperature to be 77 K. These values correspond
to the temperature at which zero resistance is achieved in YBa2Cu307.x, and the
boiling point of liquid nitrogen. Using equation 2-23, we see that A(T)/A(0)
drops to 0.38 at T/Tc = 0.86, and disappears at T = T0. Increasing the
superconducting energy gap at 77 K is one of the prime motivations for studying
higher Tc materials such as the bismuth and thallium-based superconductors.
A parameter which critically affects film and device quality is the coherence
length (£), which is the decay length for the wave function created by the
formation of Cooper pairs,7
5 =
hvF
71A
(2-24)


244
consistent with Gibb's phase rule, there must be a total of 4 three-phase regions
connected to each four-phase equilibria. Once the four-phase equilibria are
determined, the five-phase equilibria can be determined using the same process.
In this study, the Y-Al-Si-Cu-O flow chart was constructed from the
existing binary oxide phase diagrams, and the Y203 A1203 Si02 liquidus phase
diagram. In our flow chart diagrams we treated the metallic oxides as
components, and did not take into account the effect of oxygen nonstoichiometiy
on reaction temperature; hence the most complex flow chart was the quaternary
diagram. For clarity, only ternary and quaternary reactions which involve a liquid
phase are listed in the flow charts.
This flow chart was constructed in order to uncover the chemical reactions
which caused the microstructures observed in figure 4-54 to emerge after
deposition of the superconductor on the YA103 barrier layer. Since CuO has the
lowest melting point of the metallic oxides which comprise YBa2Cu307.x (the
melting points of CuO, BaO, and Y203 are 1362, 1923, and 2410 C, respectively),
we hypothesized that CuO lowers the liquidus temperatures in the system, and
chemical interactions between CuO in the superconductor and the Y Al Si O
system were responsible for the microstructures in the YBa2Cu307_x/YA103 films
on Si.
The ternary flow charts are shown in figures Al A4. The binary phase
diagrams used to construct the flow charts are shown in figures A6 All. The
Y203 CuO phase diagram was not found, so the types of reactions which are


29
proximity effect entails diffusion of superconducting electron pairs and normal
electrons across the interface to create a weak superconducting layer in the
normal metal. There is a coherence length for superconducting pairs in the
normal metal which is given by:23
hVw.
normal metal *
(2-31)
where vF is the Fermi velocity of the normal metal, and the electron mean free
path is assumed to be longer than the coherence length ("clean limit").
Conversely, if the electron mean free path is shorter than the coherence length
("dirty limit"), the coherence length is given by:
* uAt
b /
where 1 is the electron mean-free path. The critical current across an SNS
junction is given by:
(2-32)
, (2-33)
where B is a constant, and a is the normal metal thickness.
The short coherence length in YBa2Cu307.x is the primary reason why
microstructural defects have a significant effect on the superconducting properties.
Dimos et al. conducted a series of experiments designed to determine the
dependence of Jc and T0 on temperature and magnetic field, and reveal the types
of weak links responsible for reduced superconducting properties.24 For epitaxial


238
which occured because the thermally induced stresses exceeded the fracture
strength of YBa2Cu307_x. The stresses in the (103) oriented YBajQ^O^ film
were greater than in the (001) oriented film because the thermal expansion
coefficient is highest in the <001> direction, and the contribution of the <001>
direction to the in-plane thermal expansion of the film is greater in the (103)
oriented films than in the (001) YBajQ^O^ films.
YA103 and LaA103 films deposited on (1102) A1203 were both amorphous,
but the electrical properties of the YBa2Cu307x films deposited on the barrier
layers were quite different. The normal state resistance of the YBa2Cu307_
x/YA103 film deposited on (1102) A1203 was metallic, with a transition to the
superconducting state at 82 K. However, the YBa2Cu307_x/LaA103 film
deposited on (1102) A1203 was less metallic, and the transition to the
superconducting state was depressed to 57 K. The difference in electrical
properties was attributed to La segregation to the grain boundaries and decreased
intergranular coupling of the superconducting hole pairs.
Because of the large difference in thermal expansion coefficients between
YBa2Cu307.x and Si (lSxlO'VC vs 3.8xl0'6/oC, respectively) thermally induced
cracking is prevalent in the YBa2Cu307.x films. Elastic continuum mechanics
predicts that the thermally induced stress in a film will result almost entirely from
the difference in thermal expansion coefficients between the film and substrate,
and will be virtually independent of the presence of any other layers.


APPENDIX B
FLOW CHART OF THE Y AI Si Cu O SYSTEM
A flow chart of the Y Al Si Cu O system was constructed in order to
determine the chemical reactions responsible for the microstructure observed in
the YBa2Cu307.x/YA103 films on Si and A1203 substrates.
Construction of the flow chart requires that all phases appear or disappear at
equilibrium in a manner consistent with the Gibb's phase rule:
P+F=C+1 (A1)
where P = the number of phases in equilibrium, F = degrees of freedom, C =
the number of components, and the pressure is arbitrarily held at 1 atmosphere.
In this construction, only invarient points where F = 0 are considered. Three-
phase fields must begin or terminate at binary three-phase equilibria, ternary four-
phase equilibria, or quaternary five-phase equilibria (i.e. conditions where F = 0,
so the temperature at which the phases are in equilibrium, and the compositions
of each phase, are fixed). Since the temperatures of the three-phase equilibria
are known for the binary systems, the four-phase equilibria and temperature range
in which they must occur can be determined using a self-consistent analysis. In
this analysis, all three-phase regions emanating from or terminating at binary
phase diagrams are required to connect ternary four-phase equilibria. To be
243


20KU XI1000 0192 ^0U UFMSE
Figure 4-4. Scanning electron micrograph of YI^CujO^ film on (1102)
LaA103.


INTENSITY (ARBITRARY UNITS)
H*
8


RESISTANCE (OHMS)
126
100 150 200
TEMPERATURE (K)
250 300
Figure 4-21. Resistance versus temperature for a YBa2Cu307.x/Zr02 film
on Y-Zr02.


240
YBa2Cu307.x/Zr02 film on the Y-Zr02 substrate was 9X104 amps/cm2 at 4.5 K.
These data confirm that excess Y203 in the grain boundaries significantly
improves intergranular coupling of the superconducting hole pairs, and Jc values
similar to those obtained from epitaxial YBa2Cu307.x films grown onto single
crystal substrates.
Selected barrier layers were used to dramatically improve the superconducting
properties of YBajQ^O^ films deposited on various substrates. The type of
barrier layer which worked best was different for each of the substrates, and was a
function of the lattice matching, and chemical reactivity at both the barrier
layer/substrate and YBa2Cu307_x/barrier layer interfaces. In this study, we
determined that SrTi03, YA103, and Y203 were the best barrier layer materials
for YBa2Cu307_x films deposited on (1102) A1203, Si, and Y-Zr02 substrates,
respectively.


42
^film ^ substrate + substrate -film Interface
(2-49)
the surface energy is minimized by the formation of three dimensional islands,
and the islands coalesce to form the film. This is known as the Volmer-Weber
growth mode, and is characterized by strong adatom-adatom bonding. The
orientations of these films are less influenced by the substrate than in the case of
two dimensional growth, and generally have more grain boundaries.
The third mode of film growth, Stranski-Krastanov, occurs when
Gfilm & substrate + ^substrate -film interface
(2-50)
The Stranski-Krastanov mode contains features of both of the previous growth
modes, and is often characterized by two dimensional growth of the first
monolayer, followed by three dimensional growth of the rest of the film.
Assuming the film growth rate, surface energy, and adatom-substrate bond
strengths are such that two dimensional film growth predominates, the film will try
to minimize orsubstratc.nira interface by growing epitaxially on the substrate, thus
eliminating the energy required to create an additional surface at the interface.
However, if the film has a different chemical composition than the substrate, the
lattice parameter of the film will be different and there will be a strain induced in
the film as it attempts to adopt the lattice spacings of the substrate. Once the
elastic limit of the film is exceeded, the strain is relieved via formation of
dislocations at the interface and the film is no longer completely epitaxial with the
substrate.


48
deformation in the film or substrate, nor any slippage at the film/substrate
interface, stress in the film (af) can be obtained by measuring the radius of
curvature of the film/substrate combination:48
ar
EX
6(1 ~vs)Rtf
(2-56)
where ts and tf are the thicknesses of the substrate and film, vs is the Poisson's
ratio of the substrate, and R is the radius of curvature of the film/substrate
combination. There are two basic requirements imposed by continuum mechanics
which this model satisfies.49 First, the sum of the forces per unit width acting on
the film and substrate must be zero. The force per unit width is defined as (ft,
where i denotes either the substrate or one of the film layers. As an example of
how the stresses are distributed in the film and substrate, consider the case in
which the film stresses are tensile (figure 2-11). These film stresses will cause the
substrate to bow slightly, so the top of the substrate (adjacent to the film) will be
in compression. The magnitude of the forces will vary within the substrate, with
the maximum compressive stress observed at the top, and the maximum tensile
stress at the bottom of the substrate. Near the middle, there will be a plane
through which the stress (and therefore strain) is zero, and this is called the zero
strain plane. In the films, the stress is assumed to be constant throughout the
thickness of the each film layer. In this example, the tensile eft forces contributed
by the film and the portion of the substrate below the zero strain plane must
balance the compressive (ft contribution of the upper part of the substrate,


I certify that I have read this study and that in my opinion it conforms to
acceptable standards of scholarly presentation and is fully adequate, in scope and
quality, as a dissertation for the degree of Doctor of Philosophy.
Paul H. Holloway, Chair u
Professor of Materials Science
and Engineering
I certify that I have read this study and that in my opinion it conforms to
acceptable standards of scholarly presentation and is fully adequate, in scope and
quality, as a dissertation for the degree of Doctor of Philosophy.
4teza Abbaschian
Professor of Materials Science
and Engineering
I certify that I have read this study and that in my opinion it conforms to
acceptable standards of scholarly presentation and is fully adequate, in scope and
quality, as a dissertation for the degree of Doctor of Philosophy.
Timothy J. (Anderson
Professor of Chemical Engineering
I certify that I have read this study and that in my opinion it conforms to
acceptable standards of scholarly presentation and is fully adequate, in scope and
quality, as a dissertation for the degree of Doctor of Philosophy.
Richard Associate Professor of Materials Science
and Engineering


77
thermal expansion coefficients (Si and A1203), YBa2Cu307.x film cracking caused
by thermally induced stresses is a significant problem.
The cracking problems caused by the low thermal expansion coefficients of Si
have been circumvented by using a silicon-on-sapphire structure.105 YBa2Cu307.x
films grown onto a Y-Zr02 barrier layer, which in turn was deposited on a Si film
grown on a sapphire substrate, had Jc values of 4.6 xlO6 A/cm2 at 77 K for
YBa2Cu307.x film thicknesses up to 4000 A.
Y-ZrOz has also been successfully used as a barrier layer for YBajQijO^
films deposited on the (1102) plane of A1203. Highly (lOO)-oriented Y-Zr02 films
were deposited on (1102) A1203 at substrate temperatures greater than 780 C,
whereas growing the Y-ZrOz layer at lower temperatures increased the ratio of
(111)/(100) Y-Zr02 x-ray diffraction peak intensities.36 YBa2Cu307.x films
deposited on highly (100) oriented Y-Zr02 barrier layers had T0 = 90 K, and Jc
values of 1.2X106 amps/cm2 at 77 K.
SrTi03 has emerged as a very good barrier layer material for growth of
YBa2Cu307_x on (1102) A1203 substrates. 1.2 pm thick YBa2Cu307_x films grown
on a 4000 SrTi03 barrier layer, had T0 values of 86.5 K and Jc values of lxlO6
amps/cm2 at 77 K.106 X-ray diffraction indicated that the SrTi03 layers
preferentially grew with a (110) orientation, although the (200) peak was
significant. The YBa2Cu307.x films were (001) oriented. Secondary ion mass
spectroscopy showed a drastic reduction of A1 concentration in films grown on


235
the intragranular regions for both the YBa2Cu307.x/(YA103 or Y203)/Al6Si2013
films on (1102) A1203 were nearly ideal, with the asymmetry of the 335 cm*1 peak
and the small, narrow peak at 500 cm'1 indicating highly oxygenated, (001)
oriented YBa2Cu307_x grains in both films.
We believe chemical interdiffusion and incorporation of Si into the
YBa2Cu307.x film was not the mechanism responsible for degradation of
YBa2Cu307_x films when an Al6Si2013 layer was added to the barrier layer
structure. The degradation was caused by increased cracking in the YBa2Cu307.x
films. This data shows that the Al6Si2013 layer increased the magnitude of the
tensile stresses in the YBa2Cu307.x films, which indicates that the "continuum
mechanics model"48 for stress transfer between the substrate and various film
layers is not completely valid. However, this model does the best job of
explaining the experimentally observed behavior. The data also suggests that
cracking is not necessarily initiated in the YBa2Cu307.x layer, but the poor
electrical properties of the YBa2Cu307.x/Y203/Al6Si2013 film may result from the
poor fracture strength of the Y203 layer, which would allow cracks to nucleate
and propagate into the YBa2Cu307.x layer. The tensile stresses which accompany
crystallization may have also increased the stress in the Y203 layer, and induced
cracking. (222) Y203 diffraction peaks were observed, whereas the Y-A1-0 layer
did not crystallize. The thermally induced stresses in the amorphous Y-Al-O layer
were less than in the Y203 layer, so cracking was suppressed.


ACKNOWLEDGEMENTS
Several people have helped me complete this thesis. I would like to thank my
advisor, Professor Paul Holloway for being an excellent role model and showing me
how a scientist solves problems. I am especially grateful for his guidance during the
early stages of the project, and for letting me choose my own directions as the project
matured. I was fortunate to have very talented people on my committee, and
Professors Abbaschian, Anderson, Connell, and DeHoff each provided insights into
materials properties which were important to this project and will be valuable
throughout my career.
I am grateful for the help I recieved from co-workers within our group, especially
Kelly Truman and Ludie Hampton. Much of the data was collected at the Major
Analytical Instrumentational Center at the University of Florida, and Eric Lambers,
Wayne Aeree, and Richard Crockett did a superb job of keeping the instruments in
top condition, which made the data collection and interpretation possible.
I would like to thank Drs. Kul Bhasin, Felix Miranda, Mark Stan, and Crystal
Cubbage of NASA Lewis Research Center for their help. A large portion of this
thesis would not have been possible without their help.
n


223
of the interleaving oxygen layers, which occured between laser pulses, when the
cation flux was zero.
SrTi03 films were deposited on (100) MgO substrates at 40 mTorr and 200
mTorr oxygen. In both cases, the SrTi03 films were (100) oriented. Table 5-1
shows that the lattice mismatch is a minimum in the <100> SrTi03 // MgO
directions, hence we expect the <100> SrTi03 direction to be parallel to the
substrate. However, this system differs from the SrTiO3/(1102) A1203 case
because the interfacial energy can be further minimized by matching the <010>
SrTi03 // MgO directions, so growth of (100) oriented SrTi03 is favored.
In the SrTi03 //(1102) A1203 system, only the <110> SrTi03 direction provides a
reasonably close lattice match, so minor changes in the deposition parameters can
induce either the (100) or (110) orientations to form.
A YBa2Cu307.x film subsequently deposited on the (110) SrTi03 barrier layer
at 200 mTorr was (103) oriented, which minimized lattice mismatch at the
YBa2Cu307_x/barrier layer interface. The resistance versus temperature behavior
of this film was slightly metallic, which was partially attributed to the lack of an
easy direction parallel to the substrate in which the charge carriers could travel.
In terms of the Halbritter equation, both the grain boundary and non-intrinsic
intragranular resistance were significant compared to the intrinsic intragranular
resistances. YBa2Cu307_x films deposited onto (110) SrTi03 substrates show
similar normal state resistances, but exhibit a sharp transition to the
superconducting state at 86 K. On the other hand, YBa2Cu307.x films deposited


54
3.0 J/cm2 have generated YBa2Cu307.x films with the correct stoichiometry.57,58,59
YBa2Cu307.x films were also grown using a pulsed C02 laser (10.6 /im), but a
Y-enriched target was required to produce a stoichiometric film.60 By increasing
the pulse energy density of a C02 laser, stoichiometric YBa2Cu307.x films were
grown from a stoichiometric target, but the film morphology and superconducting
properties were seriously degraded by large globules in the film.61 The best films
have been grown using pulsed excimer lasers which operate at ultraviolet
wavelengths. The main problems associated with laser deposition is the presence
of fragments in the films which are ejected from the target, and nonuniform film
thicknesses.62 Shorter wavelength excimer lasers are popular because the
particulate size is smaller in these films than in films grown with longer
wavelength lasers.57
The ease with which stoichiometric films can be grown, and the presence of
particulates in the films both result from the laser-target interactions. There are
two basic ablation mechanisms: thermal and electronic.63 Thermal processes
result from rapid heating of the target surface and subsequent evaporation and
sublimation from the surface. Evaporation occurs when the laser power density
(Q0) exceeds the minimum power density necessary for evaporation, Q^64
r> 0.5
where pQ is the target density, D is the thermal diffusivity, U is the sublimation
energy, and r is the duration of the laser pulse.


2
For each of these applications, Jc values greater than ~ 10s amps/cm2 are
required.3 The sensitivity of Jc to interfacial phases and high-angle grain
boundaries in YBa2Cu307.x is one of the most significant problems, and it has
inhibited commercialization of YBa2Cu307_x films. To date, YBa2Cu307.x films
with Jc values above 10s amps/cm2 at 77 K have only been deposited on single
crystal substrates such as SrTi03, LaA103, and Y-Zr02 which nearly lattice match
YBa2Cu307_x. The ability to deposit films with high Jc values on randomly
oriented or polycrystalline substrates would be a tremendous boost towards
commercialization, since this would allow a variety of substrates with different
shapes and sizes to be used.
In this thesis, the literature is reviewed in chapter 2 in order to provide a
background for this work. The experimental techniques used to deposit and
characterize the films are described in chapter 3. The data original to this thesis
is presented in chapter 4, and is discussed in detail in chapter 5.
Chapter 2 shows how superconductivity evolves. The temperature at which
the resistance disappears (T0) and Jc are closely tied to the microstructure, and an
overview of the microstructural phenomena which most critically affects the
superconducting properties is introduced. The laser deposition process is
described, and the mechanisms which enable multicomponent films to grow with
the same stoichiometry as the target are presented. Chapter 2 concludes with a
survey of the work directed towards depositing YBa2Cu307_x films on various
substrate materials.


242
I/Io = exp -(fi/p)pt, where I/Iq is the ratio of x-ray intensity at a given depth (= t)
relative to the incident x-ray intensity, and p is the density of the phase under
consideration (= 6.38 g/cm3 for YBa2Cu307_x). Assuming the x-ray penetration
depth is equal to the depth at which I/Iq drops to 0.368, we obtain:
0.368 = exp -(174.0)(6.38)t.
Solving for t yields an x-ray penetration depth of 9 pm.


86
Scanning Electron Microscopy
A JEOL JSM 35C scanning electron microscope (SEM) operated at 15 20
kilovolts accelerating voltage and 100 microamps beam current was used to
visually determine the surface morphology and microstructures of the films.
Microstructural features are readily imaged in the SEM because secondary
electron detection is highly sensitive to the angle between the ejected electrons
and detector, so the number of secondary electrons detected varies as the
topography of the film changes.112 YI^Q^O^ films deposited on highly reactive
substrates or barrier layers tended to be cracked or have rough surfaces, whereas
superconducting films grown on inert substrates were smooth and featureless.
Correlations between SEM micrographs and electrical data were instrumental for
clarifying the types of microstructures which resulted in YBa2Cu307.x film
degradation.
Scanning Auger Electron Spectroscopy
Film uniformity and interdiffusion between the various layers and Si substrates
were analyzed using Auger electron spectroscopy (AES). A Phi model 660
scanning Auger microprobe, controlled by an Apollo domain series 3500
computer, was used to determine the film composition as a function of depth.
The operating parameters for the Auger were 5 kilovolts accelerating voltage, 25 -


RESISTANCE (OHMS)
116
Figure 4-13. Resistance versus temperature data for a YBa2Cu307.x/Y-Zr02 film
deposited on Y-Zr02.


92
for the film resistivity as the ratio of sample width to probe spacing (d/s) are
presented in table 3-1.
There were several practical difficulties which made calculating the
resistivities difficult. First, there was a trade-off between growing the highest
quality films and accurately assessing the film thickness, because masking off an
area of the substrate to create a step for profilometer analysis created thermal
gradients in the sample, which damaged the superconducting film quality. Second,
the substrates on which the films were deposited had a variety of shapes, so the
precise geometrical correction factor was difficult to establish. Although the
geometric factors were diverse, it is unlikely that they altered the resistivities by
more than + /- 50%, since the sample width/probe spacing ratios were
constrained by the design of the R vs. T measurement apparatus to be between 2
and 3. On the other hand, measured resistances often varied by an order of
magnitude or more, depending on the film microstructure and type of substrate
used. The additional information which could have been obtained by determining
the film resistivities would have been minor, and the fundamental
microstructural/electrical property correlations which controlled the electrical
properties were more fully uncovered by correlating data obtained by electrical
resistance measurements with other types of microstructural data.


TABLE OF CONTENTS
ACKNOWLEDGEMENTS ii
ABSTRACT vi
CHAPTER
1. INTRODUCTION 1
2. LITERATURE REVIEW 4
Devices 4
Microstracture/Superconductor Relations 18
Weak link Behavior 27
Nucleation and Epitaxial Growth 35
Thermally Induced Stresses 46
YBa2Cu307.x Film Growth 53
Barrier Layer Technology 67
3. EXPERIMENTAL TECHNIQUES 81
Film Growth by Laser Deposition 80
X-ray Diffraction 84
Scanning Electron Microscopy 86
Scanning Auger Electron Spectroscopy 86
Electrical Resistance Measurements 90
Critical Current Density 94
Raman Spectroscopy 96
Millimeter-Wave Transmission Measurements 99
4. RESULTS 101
YBa2Cu307.x on (1102) LaA103 101
IV


53
thermal expansion coefficients between the each layer and the substrate, and is
independent of the other layers. However, the top layer creates an additional
interface which pins dislocations. Hence plastic deformation in the initial layer is
suppressed, and the yield strength is increased.
The elastic continuum model is a popular model for predicting film stresses in
the elastic limit. However, the assumptions of perfect bonding at each of the
interfaces, and chemical homogenity within each layer are not realistic in most
cases. Microstructural features such as chemical bonding at interfaces, grain size,
and film thickness will affect the stress at which the mode of deformation changes
from elastic to plastic, and must be taken into account when interpreting the
stresses required to form cracks.
YBaoCugO-, Film Growth
Shortly after the discovery of superconductivity in YBa2Cu307.x, the difficulty
of growing stoichiometric YBa2Cu307.x thin films by conventional methods such as
sputtering52,53 and electron beam evaporation54 became apparent. Depositing a
film with the correct Y:Ba:Cu ratios is a difficult process using these techniques.
Pulsed laser deposition became a popular technique for growing superconductor
films because stoichiometric films could be readily grown from a target with the
same composition.55,56 Pulsed lasers have generated stoichiometric films over a
wide range of pulse energies, laser wavelengths, and pulse durations. Nd-YAG
lasers operating at 1064, 532, and 355 nm and pulse energies ranging from 0.6 to


8
magnetic flux travels. This dramatically increases the area over which the flux is
measured. For two junctions placed in parallel,
1,(4 = /c(0)cos^l (2-5)
0
The high sensitivity of SQUID devices make them attractive for a variety of
biomedical, geological, and military applications.
l*oA HoA
APPLIED MAGNETIC FIELD STRENGTH, Ha
Figure 2-2. Periodic variations in the critical current density with increasing
magnetic field strength. Reference 4.


L + Y2Si05 Y4S3012 + Y2Cu2Os
Y2S05 Y4Si3012 Y2Cu2Os
L Y4Si3Or Y2Cu205
~i l
L + Y4Si3012 Y2Si207 + Y2Cu205
Y4S3012 Y2Si207 Y2Cu205
L Y2Si207 Y2Cu205
I i ,
L + Y2Si207 ** Si02 + Y2Cu2Os
Y2S207 Si02 Y2Cu2Os
L Si02 Y2Cu205
Ni
Ul
Figure A-3. Continued.


RESISTANCE (OHMS)
140
100 150 200 250
TEMPERATURE (K)
300
Figure 4-33. Resistance versus temperature for a YBa2Cu307.x film on (100)
SrTi03.


RESISTANCE (OHMS)
164
Figure 4-53. Resistance versus temperature data for YBa2Cu307.x/YA103 on: (O)
(100) Si; () (1102) A1203.


260
A1203 Y203
Figure A-6. Phase diagram of the Y203 A1203 system. Reference 136.


232
very similar. This indicates that the Al6Si2013 layer was equally successful at
preventing the diffusion of Y or La to the substrate. Comparison of these
micrographs with the morphologies of YBa2Cu307.x/(YA103 or LaA103) films
deposited on Si also showed that the microstructures of YBa2Cu307_x films grown
on Si using YA103 or LaA103 barrier layers (without the Al6Si2013 layer) resulted
largely from a rare earth silicide type of reaction. X-ray diffraction shows that
both the YBa2Cu307.x/(YA103 or LaA103)/Al6Si2013 films on Si were (001)
oriented, and the smaller degree of cracking relative to the YBa2Cu307_x/LaA103
film on Si was attributed to the reduced strain associated with a highly (001)
oriented film. YBa2Cu307_x/LaA103 films deposited on Si contained both (001)
and (103) oriented grains, and the cracking was much more extensive.
The normal state electrical resistances for the YBa2Cu307_x/(YA103 or
LaA103)/Al6Si2013 films on Si were both slightly metallic, with T0 values of 61.5
and 59 K, respectively. Analysis of this behavior in terms of the Halbritter
equation suggests that the predominant contribution to the normal state electrical
resistances were intergranular defects (i.e. cracks), and the tail in the resistance
versus temperature profile could be attributed to weakened coupling of
superconducting charge carriers across the grains. However, examination of the
Raman spectra show that the intragranular microstructures of these two films are
markedly different. Comparison of Raman data from the YBa2Cu307_x/(YA103 or
LaA103)/Al6Si2013 films on Si shows that the film containing YA103 has a
pronounced peak at 500 cm"1, indicative of a large fraction of non-(001) oriented


33
p oo E (psi)t* pi) <2-35)
where ^(p#,) is the sum of the resitivities associated with all grain boundaries, and
(a'T + p0Ll) describes the intragranular resistivity. The percolation parameter, p,
accounts for increases in the percolation distance of the transport current resulting
from microstructural defects including grain boundaries and cracks. When p^ is
small compared to p(a1 + p0L), intergranular defects have little effect on p, and
separation of inter and intragranular defects is difficult. If the current pathway of
lowest resistance is through the grain boundaries, the net effect of a change in p^
on R vs. T is to shift the curve to higher resistance values, without changing the
slope (p values) of the curve. Extrapolation of the curve to 0 K will result in
higher values of ^(p#,) when grain boundary resistance becomes larger, but not so
large as to significantly alter the current pathways. However, when ^ip#,) is
comparable to or greater than the intragranular resistivity, the percolation
distance of the transport increases as the current selects pathways which avoid
grain boundaries or other microstructural defects such as cracks, thereby
increasing p.
The same defects which increase normal state resistivity are responsible for
depressions in T0 and J,..2829 Intragranular Josephson junctions resulting from
oxygen disorder or twinning locally depress the order parameter, and thus the T0
within the grain. The most damaging defects are usually intergranular, and are
observed at high angle (> 20 ) grain boundaries where dislocated regions often
have significant oxygen disorder.25 At high angle grain boundaries,


43
There are a variety of methods by which the film will attempt to minimize
differences in lattice parameters with the substrate. Starting with the definition of
lattice misfit ( = f),40
lattice misfit (=/) = -(tbn.?ub^te (2-51)
substrate
where afllm and asubstrate are the lattice parameters of the film and substrate. It has
been observed that the orientation which a film adopts with respect to the
substrate is primarily a fimcton of lattice misfit, as well as the adatom-film and
adatom-adatom bonding strengths. If the lattice parameters of the film and
substrate are very close (generally < 1%) and both have the same crystal
structure, the film will align itself so the film directions in the plane of the
interface match those of the substrate. For larger mismatches, the film will adopt
an in-plane orientation such that the maximum number of lattice points will be
commensurate with the substrate.41 Theoretical and experimental data indicate
that the unit cell over which the film minimizes misfit and becomes commensurate
with the substrate may exceed 500 , hence the correlation between film and
substrate directions is not always obvious or simple to deduce. One commonly
observed method of minimizing strains induced by dissimilar surface meshes is to
strain the lattice of the film so as to match row spacings in one of the directions.42
The Nishiyama-Wassermand and Kurdjumov-Sachs matchings are shown in figure
2-10. The difference between the two orientations is the Kurdjumov-Sachs


25
The penetration depth varies with superconducting electron density, and hence
temperature. The Gorter-Casimir expression for A is:11
A(7) =
T 4 -
(2-7)
where A is the penetration depth at 0 K, and varies with film quality. For a
highly-(001)-oriented, 3000 thick YBa2Cu307_x film grown on LaA103, A0 = 1800
was measured.17 Films with larger fractions of non-(001) oriented grains, or
which contain high-angle grain boundaries, had larger penetration depths.
The way in which a superconductor responds to a magnetic field is dependent
on the Ginzburg-Landau parameter k, defined as4
k=0.966 (2-27)
*0
The magnetic field may be applied, or may be induced by the transport current (a
self-field). Superconductors with k < ]2 are classified as Type 1, while those
with k > ] 2 are Type 2. In both types, there is a magnetic contribution which
increases the free energy density, AG, while the electron ordering associated with
the formation of Cooper pairs lowers AG. If the AG associated with electron
ordering occurs over a shorter distance from the surface than the magnetic
contribution ( as in Type 2 superconductors), there will be a minimum in the free
energy density at the surface, and it becomes energetically favorable to form an
interface between the normal and superconducting regions (figure 2-8). Thus
there are isolated circular regions in the sample which are normal, while the rest


139. A. Bondar and F.Y. Galakhov, in Phase Diagrams for Ceramists. 1969
Supplement, p. 165.
140. A.M.M. Gadalla and J. White, in Phase Diagrams for Ceramists. 1969
Supplement, p.13.
141. A.S. Berezhnoi, L.I. Karyakin, and I.F. Dudavskij, in Phase Diagrams for
Ceramists, edited by E.M. Levin, C.R. Robbins, and H.F. McMurdie
(American Ceramic Society, Columbus, 1964), p. 86.


YBa2Cu307.x on La AI03 (1102)
Figure 4-2. Magnetization (M) versus magnetic field intensity for a YBa2Cu307_x film on (1102) LaA103. AM is the
width of the magnetization curve at a given magnetic field intensity, and Jc is derived from Jc = 15 AM/r,
where is the radius of the film. The Jc of this film was 5xl07 amps/cm2 at 4.5 K. _*


Chapter 3 describes the film growth technique, and the methods used to
characterize the films. Each of the characterization tools relies on different
physical phenomena to probe the film microstructures, so different aspects of the
microstructures were uncovered. A brief description of the operating principles of
each of the probes is given in this chapter.
The experimental results are presented in chapter 4. The sections are
arranged so as to compare the effectiveness of each barrier layer to induce growth
of optimal quality YBajCujQ^ films on different substrates. The text explains
what information the data is providing, and points out the most notable features
in each figure.
Chapter 5 compares and contrasts the data in order to explain the observed
phenomena. The primary objective of this thesis is to uncover the microstructural
features which were primarily responsible for the normal state and
superconducting properties. By using a variety of experimental techniques, we
arrived at a more clear understanding of film microstructures than if only one or
two techniques were used; thus we were able to correlate the film microstractures
with the electrical properties.


CHAPTER 3
EXPERIMENTAL TECHNIQUES
Film Growth by Laser Deposition
Barrier layer and YBajCi^O^ films were sequentially deposited at 730 750
C on silicon (Si), aluminum oxide (A1203), yittria-stabilized zirconia (Y-Zr02),
lanthanum alumnate (LaA103), or strontium titanate (SrTi03) substrates using a
pulsed laser deposition system. A Questek model 2560 pulsed excimer laser,
using KrF gas and operating at 248 nanometers, 30 nanosecond pulses, and 5
pulses/second, was focused to 2.5 3.0 Joules/cm2 with a 50 centimeter focal
length lens onto a one-inch diameter YBa2Cu307.x or barrier layer target. A
schematic diagram of the deposition system is shown in figure 3-1. The
stoichiometric YT^CujOy.* target was obtained from Ceracon, Inc. and was 96%
dense. The barrier layer targets were fabricated by mixing stoichiometric ratios of
the powders, calcining at 950 C, then pressing the powders into disks and re
firing at 950 C for 12 hours. The barrier-layer targets were approximately 65%
dense. Up to three targets could be mounted on a stainless steel holder. By
rotating the target holder, sequential films were deposited without breaking
80


108
Raman spectra from a LaA103 substrate, as well as 500 and 1000 thick
YBa2Cu307.x films deposited at 750 C. Increasing the deposition temperature
significantly reduced the height of the 500 cm'1 peak. For thicknesses greater than
1000 , Raman transitions from the substrate were masked by the film.
YBa3Qi3Q:, It/Y-ZrO, Films on Si. Y-ZrCX,. and LaAlQ3 Substrates
The most notable microstructural differences between YBa2Cu307_x films
deposited on Si versus Y-ZrOz substrates was the presence of cracks in
superconducting films on Si (figure 4-7). In all of the YBa2Cu307.x/barrier layer
films deposited on Si substrates, cracking was observed to some extent. SEM
micrographs of Y-Zr02 films deposited at 730 C on Si were featureless (figure
4-8), indicating that the cracking occurred during the YBa2Cu307.x film growth
and cooling process. We will show that strain induced by the difference in the
thermal expansion coefficients (table 4-1) was the principal cause of YBa2Cu307.x
film cracking on Si substrates, and reduction of the thermally induced stresses
significantly reduced YBa2Cu307.x cracking.
A primary objective of this study was to identify the mechanisms by which the
thermally induced stresses were relieved. Figure 4-9 shows the resistance vs.
temperature data for YBa2Cu307.x grown directly on an off-axis single crystal Y-
Zr02 substrate, and on Si with a Y-Zr02 buffer layer. After multiplication of the
normal state resistance values of the YBa2Cu307_x/Y-Zr02 film on Si by 0.061, the
resistance versus temperature plot is similar to that of YBa2Cu307.x films grown


Figure 4-26. Scanning electron micrograph of a YBa2Cu307.x/Zr02 film on (100)
SrTi03.


23
where vF is the velocity of electrons at the Fermi level and A is the size of the
superconducting energy gap. An equivalent definition is that the coherence length
is the minimum distance over which the density of superconducting electrons can
vary, and hence the minimum distance between superconducting and
nonsuperconducting regions. The coherence length (£) is much shorter in
YBa2Cu307.x than it is in conventional superconductors such as Nb3Sn, and is
highly dependent on crystallographic direction (table 2-1). The short coherence
length in YBajQ^O^ (table 2-2) is primarily due to the small Fermi velocity of
the superconducting electrons, and an important consequence of the short
coherence length is that microstructural defects, such as grain boundaries,
impurity atoms, dislocations, or chemically unstable surfaces which create
imperfect or disordered regions of similar size to the coherence length, can
significantly alter the superconducting wave function, especially in the <001>
direction. By contrast, microstructural defects are considerably less degrading to
the order parameters of conventional superconductors.
Table 2-1. Normal state parameters of conventional metals (such as Nb3Sn) and
YBa2Cu307_x. Reference 7.
Parameters
Conventional
Metals
YBa2Cu307 x
II (001)
1 (001)
m*
1-1.5 mc
5 mc
25 mc
Er (eV)
5-10
0.3
0.3
kF (cm)1
108
5xl07
5xl07
vF (cm/sec)
1-2x108
107
2xl06


CRITICAL CURRENT DENSITY (10 A/CM )
205
MAGNETIC FIELD (T)
Figure 5-3. Jc versus magnetic field intensity for YB2l2Oi301.JY203 films on Y-
Zr02 and (100) SrTi03 substrates. Data taken at 4.5 K. The solid
and dashed lines are fit of the experimental data to the flux creep
model.


L + Y4A]209 YA103 + Y2Cu205
Y4A1209 YA103 Y2Cu205
L Y4A1209 Y2Cu205
1 I
L + YA103 Y3A15012 + Y2Cu2Os
yaio3 Y3A15012 Y2Cu205
L Y3A15012 Y2Cu205
1 1
L + Y3A15012 ** Y2Cu205 + A1203
Y3A15012 Y2Cu205 A1203
L Y2Cu205 A1203
Figure A-4. Continued.


57
ablation, and to optimize the film growing process. The film thickness across the
substrate varies as cosn0, (n ~ 4),67 which indicates the ejected particle distribution
from the target is highly peaked. In addition, the cation ratios in the film are no
longer stoichiometric at lateral distances greater than 20 degrees from the target
Figure 2-12. Resistivity as a function of temperature for three YBa2Cu307.x films
deposited on (100) SrTi03 substrates by a Nd:YAG laser and it's
second and third harmonics. Reference 57.


94
-Critical .CuirentPensity
For commercial applications such as interconnect lines, the critical current
density (Jc) is the most important figure of merit of a superconducting film. In
this set of experiments, Jc values were determined using the Bean model,116,117,118
in which circulating currents are induced in the superconducting films by an
applied magnetic field (Ha). In this technique, a DC magnetic field is applied
perpendicular to the film surface. Circular eddy currents are generated in the
film, and these currents propagate near the edges of the film in order to shield
the interior of the film from Ha. Because of the shielding currents, the magnitude
of H is less near the center of the film than at the edges, and the relationship
between H and Jc is given by:
(3-6)
10
To measure Jc, the intensity of Ha is increased beyond the point at which the
film can shield its interior from the magnetic field via circulating currents, so the
magnetic field completely penetrates the sample, and the film becomes a normal
conductor. Ha is then reduced, and the circulating currents reverse direction in
response to the change in Ha. When Ha = 0, the remanant magnetization (Br)
created by the circulating currents is given by:


61
Figure 2-14. The [100] and [010] lattice parameters of YBa2Cu307.x versus
temperature for a bulk sample heated in 100% oxygen. Reference
71.


72
behavior was attributed to the high angle grain boundaries which resulted from
the mixed in-plane orientations of the YBa2Cu307_x films.
Films grown on (100) MgO substrates also showed T0 values of 89 K and Jc
= lxlO6 A/cm2 at 77 K.96,97 Similar to Y-Zr02, there is a large YBa2Cu307.x/
MgO lattice mismatch (5.9 and 4.2% in the YBa2Cu307.x <100> and <0io>/MgO<100>
directions, respectively). YBa2Cu307.x films grown onto MgO at 670 C were
predominantly (001) oriented. Unlike Y-Zr02, there was no evidence of
interfacial reactions at the YBa2Cu307.x/Mg0 interface. Similar to the case for
Y-Zr02, it was speculated that YBa2Cu307.x grows on MgO in a manner which
permits matching of the oxygen sublattices. The model also proposed that the
large lattice misfits are accomodated by 90 boundaries and stacking faults.
The third group of substrate materials is characterized by materials which
chemically react with YBa2Cu307_x, thus making growth of high quality
YBa2Cu307_x directly onto these substrates very difficult. Unfortunately, the two
substrate materials most widely used for electronic applications-(lOO) silicon for
integrated circuits and A1203 (sapphire) for micro and millimeter-wave electronics,
belong to this group. A 1.5 /m thick YBa2Cu307.x film grown in-situ at 700 C
onto a Si substrate showed T0 = 70 K, and an interfacial reaction layer ~ 0.5 fim
thick was observed.98 Auger electron spectroscopy detected Si at the film surface,
indicating Si diffusion through grain boundaries and microcracks. The poor film
quality was also attributed to microcracks generated by the large difference in
thermal expansion coefficients for Si (3.8xlO^/C) and


INTENSITY (ARBITRARY UNITS)
114
Figure 4-11. X-ray diffraction patterns from a YBajO^O^ film on Y-ZrO,, and a
YBa2Cu307.x/Y-Zr02 film on (100) Si.
(200)


INTENSITY (arbitrary units)
141
<0
8
YBa2Cu307_x/SrT103
on AI2O3 (1102)
CO
o
CO
8
m
8
52 I
CM
8

CO
o



SITIOg

AI2P3
I

SrTlC>3 on Al203 (1102).
Deposited at 200 mTorr 02>
-Q
a
CM
O
1
o
o
1
1
I
Jl
11
10
20 30 40 50
DIFFRACTION ANGLE (20)
60
Figure 4-34. X-ray diffraction patterns for SrTi03 and YBa2Cu307.x/SrTi03 films
deposited on (1102) A1203. SrTi03 films deposited at an oxygen
pressure of 200 mTorr.


90
Electrical Resistance Measurements
Electrical resistance versus temperature data were taken in order to correlate
normal state resistances, the onset of superconductivity (Tonsct), and the
temperatures at which the DC resistance dropped to zero (T0), with YBa2Cu307.x
film microstructures. Resistance measurements were obtained using a four-point
probe apparatus in which current was transported through the film by the outer
two terminals, and the voltage drop was measured across the inner two terminals.
The four probes were mechanically pressed against the sample, and either 10 or
100 microamps were transported through the sample. T0 was determined when
the voltage drop was less than 1 microvolt (R < 0.1 or 0.01 ohms. Resistance vs.
temperature data were also obtained on many of the samples at NASA Lewis
Research Center. The system at NASA was more sensitive than the one at the
University of Florida; gold contacts leads were wire bonded directly to the
superconducting films, which increased the sensitivity in the low resistance regime
near the transition temperature. The criteria for superconductivity in this system
was R < 0.001 ohms. Despite the differences, both measurement apparatuses
produced very similar normal state resistance and transition temperature data
when the same samples were tested on both systems.
For thin films, resistivity rather than film resistance is usually plotted because
resistivity is a material property, and the calculations used to obtain resistivity


INTENSITY (ARBITRARY UNITS)
161
Figure 4-50. X-ray diffraction patterns from LaA103 and YBa2Cu307.x/LaA103
films on (1102) A1203.


L ** CuA102 + CuO
L ** CuO + SiO-,
(1060 C)
L ** Si02 + CuA102 + CuO
SiO-, CuAlO-, CuO
it it
Figure A-2. Continued.


RESISTANCE (OHMS)
193
60
50
40
30
20
10
0
Figure 4-76.
Resistance versus temperature data for: (O) YB2^Ca307.J
YAlOj/A^SiA, on (1102) A1203; () YBa^O^/YA/AljSijO*
on (1102) A1203.


101
YB^Ci^Ot x on (1102) LaAlQ3
Before examining the properties of the various YBa2Cu307.x/barrier layer
structures, the electrical properties of a nearly ideal YBa2Cu307.x film are
reviewed in order to establish a benchmark to which the YBajQ^O^/barrier
layer films deposited on Si and A1203 substrates can be compared. Figure 4-1
shows resistance versus temperature data for two YBa2Cu307_x films deposited on
(1102) LaA103 substrates. The first film was cooled to 550 C before filling the
chamber with oxygen. The resistance versus temperature profile of this film was
non-linear, with a T0 of 79 K. For the second film, the chamber was filled with
oxygen while the temperature was still 730 C, and this film had nearly ideal
electrical properties. The normal state resistance was metallic, and extrapolated
to 0 ohms at 0 K. There was a sharp transition to the superconducting state at
88 K, and the critical current density was 5xl07 amps/cm2 at 4.2 K (figure 4-2).
X-ray diffraction data shows the film was predominately (001) oriented, with
minor (100) and (010) peaks also observed (figure 4-3). SEM micrographs (figure
4-4) show the film is smooth, and the most prominent features are debris from the
target.
The Raman spectra from YBa2Cu307_x films deposited on (1102) LaA103
substrates was sensitive to the deposition temperature and the film thickness.
Raman spectra from the film deposited at 730 C shows a pronounced peak at
500 cm*1 and an asymmetric peak at 335 cm*1 (figure 4-5). Figure 4-6 shows


INTENSITY (ARBITRARY UNITS)
151
Figure 4-42. X-ray diffraction patterns from SrTi03 and YBa2Cu307.x/SrTi03
films on (0001) A1203.


217
Effects of Lattice Matching
Minimization of the lattice mismatch and matching of the surface meshes
strongly influenced the orientation of SrTi03 films deposited on A1203 substrates.
SrTi03 films deposited at 200 mTorr onto (1102) A1203 were predominately (100)
oriented, while films grown on (0001) A1203 under identical conditions were
strongly (111) oriented. Because the two dimensional mesh of the (1102) A1203
surface has an almost tetragonal mesh (the angle between the <100> and
<012> directions is 86), it is not suprising that SrTi03 adopts a cubic overlayer
mesh. Similarity, the oxygen atoms of the (0001) A1203 surface are hexagonalty
close-packed, so the SrTi03 overlayer adopts the (111) orientation, in which the
oxygen atoms are also hexagonalty close-packed. Note that the (111) SrTi03
orientation is predominant despite a lattice mismatch of 15.8% with the A1203
(0001) substrate. Matching the type of surface mesh appears to be as important
as the lattice mismatch in the growth of highly oriented SrTi03 films. Oriented
SrTi03 films were not observed on (1210) A1203, and this was attributed to the
lack of a close-packed surface mesh on the A1203 substrate. The smallest surface
mesh for the (1210) surface is a tetragonal mesh with lattice parameters of 12.99
x 8.24 (see figure 5-6).


Figure 4-47. X-ray diffraction pattern from a YBa^O^/LaAlOa film on (100)
INTENSITY (ARBITRARY UNITS)
ro
o
O
T1
n
33
>
O
H
O u
O
r~
m
13
3 *
o
Ol

(002)
(003)
(005)
-<
CD
O
F
r

>
cP
O
3
W
8
(006)
-j


55
Assuming the optical absorption depth (a1) of the target is small compared to
thermal diffusion length, the relationship a(2Dt)0^ > > 1 is valid. The
temperature rise of the target surface layer (where the thickness, t, of the surface
layer heated by the laser is defined as t = (2DT)0,5), can be approximated by
comparing the energy absorbed during the laser pulse with the thickness of target
surface heated by the pulse:
Ar=(i-)-
it
cvPo(2Dty
0.5
(2-61)
where R is the target reflectivity, I is the power density (W/cm2), t is the duration
of the laser pulse, and Cv is the specific heat.
Since the thickness over which laser energy is absorbed in the target is
inversely proportional to the absorption coefficient, target materials with high
absorption coefficients will attain higher surface temperatures because the energy
is confined to a smaller volume. The absorption coefficient is also a function of
the laser wavelength (A), and the absorption coefficient for YBa2Cu307.x increases
as A decreases.65 Congruent evaporation from a multicomponent target occurs
because the components cannot segregate over a region greater than the thermal
diffusive region (2DT)0-5 during the time over which the target is irradiated.
Because the temperature within the thermal diffusive region is high enough to
evaporate all the components, the entire region is evaporated and the film
stoichiometry is the same as that of the target.


24
A final parameter which is important for understanding the magnetic and
electrical transport properties of superconductors is the magnetic penetration
depth, A, which is a measure of how deeply a magnetic field penetrates into a
superconductor. The depth to which a magnetic field can penetrate into a
superconductor is limited because the superconductor will generate circular eddy
currents, which create an internal magnetic field that nullifies the applied field. A
is defined as4
f~B(x)dx=XB0 (2-25)
where B0 is the applied magnetic field, and B(x) decays exponentially as it enters
the superconductor (figure 2-4). By virtue of the Maxwell equation,21
Vx£ = + 4nJ, (2-26)
dt
A is the maximum depth at which a transport current (J) can flow, hence it is
analogous to the skin depth in normal metals.
Table 2-2. Superconducting parameters of conventional metals (such as Nb3Sn)
andYBa2Cu307_x. Reference 7.
Parameter
Conventional
Metal
YBa2Cu307 x
II (001)
X (001)
T0 (K)
< 23
95
95
2A/kBT0
< 4.4
5-8
2-3.5
A/Ep
lo-4
2x1o1
lxlO'1
to (A)
lOMO4
15
7


179

%

%
20KU XI 1000 0070
1.0U UFMSE
Figure 4-66. Scanning electron micrograph of a YI^Q^O^/YjOj film on (100)
SrTi03.


62
incomplete growth of the orthorhombic phase. Rapidly cooled YBa2Cu307.x
samples will be comprised of orthorhombic nuclei surrounded by an oxygen
depleted, tetragonal phase matrix. YBa2Cu307.x films prepared by heating an
amorphous film to 850 900 C in oxygen, then slowly cooling through the
tetragonal-to-orthorhombic transition, have T0 values of approximately 85 K on
nonreactive substrates such as SrTi03, and T0 * 75 K on the more reactive A1203
substrates.55,56 Post-annealed films on (100) SrTi03 substrates are polycrystalline,
but are preferentially textured with the (001) plane parallel to the substrate.
Major improvements in Jc values were accomplished by depositing the films
in-situ at elevated temperatures and controlled oxygen pressures. Epitaxial
YBa2Cu307.x films on (100) SrTi03 and (100) Y-Zr02 substrates were grown at
substrate temperatures ranging from 500 650 C and 200 mTorr 02, and
subsequently annealing the films at 450-500 C in 760 Torr 02 for 60 minutes.73,74
In-situ films grown on (100) SrTi03 and (100) Y-Zr02 had Jc values of 5X106
amps/cm2 and lxlO6 amps/cm2, respectively, at 77 K; both had T0 values of 90
K. In addition, films grown in-situ have much lower room-temperature
resistivities (160 ^ohmXcm) than post-annealed films (~ 1 milliohmXcm). This
behavior was attributed to the predominance of (001) orientation in the in-situ
film, whereas the post-annealed films contained a mixture of (001), (100) and
(010) oriented grains. The microstructures of in-situ films also formed an
epitaxial orientation with the substrate, whereas a "basket-weave" structure was
observed in post-annealed films.75 The basket-weave structure arises from the


125
. 7 j|jJ ft
"'ft
%
#
20KU XI 1000 0058 ~ UFMSE
Figure 4-20. Scanning electron micrograph of YBa2Cu307 J Y-Zr02 film on
(1102) LaA103.


188


*
>

fe



1%

%
1
it %
m
w
20KU XI1000 0004
*
1.0U
0
A
"UFMSE
Figure 4-73. Scanning electron micrograph of a YBa2Cu307_x/LaA103/Al6Si2013
film on (100) Si.


34
superconducting electrons are preferentially transported through the grain
boundaries at microbridges where the oxygen disorder is minimal (figure 2-9).
Interfacial impurity phases increase the separation between superconducting
regions, further degrading the interface and making it more difficult for
superconducting electrons to tunnel through the grain boundary. Because the
coherence length of superconducting electrons is longer in the <100> and <010>
than in the <001 > direction, T0 and Jc in (001) oriented films are not as sensitive
to grain boundaries as are films with other orientations. However, high angle
grain boundaries are highly dislocated regions where impurity phases tend to
accumulate, so the order parameter is depressed in (001) oriented films with
random in-plane orientation. By contrast, low angle grain boundaries are
generally free of interfacial phases and do not depress superconductivity. In
addition, 90 grain boundaries, which usually result from twinning or adjacent
(010), (100) and (001) oriented grains, generally do not contain significant
amounts of impurity phases, and hence do not reduce T0 or Jc values.30
Figure 2-9. Intergranular defect with one microbridge. Cuprate regions near the
grain boundary are normal conducting, and the superconducting order
parameter is diminished in these regions. Reference 26.


INTENSITY (ARBITRARY UNITS)
172
Figure 4-60. X-ray diffraction pattern of a YBa2Cu307.x/YA103 film on (1102)
ai2o3.


27
Weak Link Behavior
Because of the short coherence length, microstructural defects significantly
affect the electrical properties of YBa2Cu307_x. Defects are regions in which the
amplitude of the superconducting wave function is depressed, and the density of
superconducting electrons is lowered. Since these defects reduce the
superconducting order parameter, reductions in T0 and Jc result if the transport
current is forced to travel through the defects.6
Defects are classified as either superconductor-insulator-superconductor (SIS)
or superconductor-normal-superconductor (SNS), depending on the electrical and
magnetic properties of the junction.4 For the case in which two superconductors
are separated by a thin insulating layer, there is a finite probability that
superconducting electron pairs will tunnel through the insulator and a
superconducting current will be maintained. Because of the insulating layer, the
wave function is no longer continuous between superconductors, but there is a
phase difference, AO, between the wavefunction at each of the interfaces which
determines the maximum supercurrent, Is, which can pass through the insulator:5
/5=icsinA (2-1)
If the transport current exceeds ic, AO is no longer constant with time, and a
voltage appears across the junction. The currents which result from the time
varying AO and subsequent voltage across the insulator are given by:4


148
which was due to the superconducting charge carriers, was greater than the real
part of conductivity at temperatures below 80 K (figure 4-40).
Lowering the oxygen pressure during growth of the SrTi03 barrier layer
resulted in a damaged YBajCujO^ film. The (103) oriented YBa2Cu3O7.x/(110)
SrTi03 film on (1102) A1203 was only slightly metallic, and did not reach T0 until
77 K. The critical current density for this film was 5x10s amps/cm2 at 4.2 K.
The electrical properties of YBa2Cu307.x/SrTi03 films deposited on the
(1210) and (0001) planes of A1203 were severely degraded. Figure 4-41 shows the
YBa2Cu307.x/SrTi03 film on (0001) A1203 had high normal state resistance values
which were almost independent of temperature, and there was a broad
temperature range (88 51 K) over which the resistance dropped from 36 to 0
ohms, with a long tail near 0 ohms. Similarily, the normal state resistances of
YBa2Cu307.x/SrTi03 on (1210) A1203 were much higher than the resistance values
of an ideal YBa2Cu307.x film, and the temperature dependence was
semiconducting, with a broad transition to T0 = 52 K. X-ray diffraction data
shows that SrTi03 grew with a strong (111) orientation on (0001) A1203, and the
YBa2Cu307.x film deposited on this structure was (001) oriented (figure 4-42).
SrTi03 grew with a weak (110) orientation on (1210) A1203, and the YBa2Cu307.x
film was weakly (103) oriented (figure 4-43).
Raman spectra from YBa2Cu307.x/SrTi03 films show the 335 cm'1 peak is
higher than the 500 cm'1 peak for the film on (1102) A1203, but the 500 cm'1 peak
is larger than the 335 cm'1 peak for the film on (1210) A1203 (figure 4-44).


230
resulted from Al diffusion to the YBajCojO^, grain boundaries and reduced the
intergranular coupling between grains, which would account for the broadened
transition between Tonset and T0. Despite the adverse effects of A1 incorporation
into the YBa2Cu307_x film, the combination of Y and A1 oxides in the barrier layer
caused only a slight degradation in T0 of the YBajQ^O^ film.
Much of the phenomena responsible for the microstructure observed in the
YBa2Cu307.x/YA103 films can be inferred from phase equilibria. The flow chart
diagram for the Y-Al-Si-Cu-O system is shown in appendix B. This diagram
shows that incorporation of A1 into the Y-Cu-Si-O system lowers the liquidus
temperature. The flow chart predicts that the lowest temperature for a liquid
phase occurs at the l^Si02-Al203-CuA102-Cu0 five-phase equilibria, and SEM
micrographs of YBa2Cu307.x/YA103 films on Si indicate a liquid phase may have
formed at the grain boundaries.
The sharpest transition for YBajQijO^ films deposited on Si substrates were
observed when YA103 barrier layers were used to prevent interdiffusion. The
sharp transition was attributed primarily to the greatly reduced cracking (relative
to the YBa2Cu307.x/Y-Zr02 film on Si). The most probable reason for reduced
cracking in these films was that the thermally-induced tensile stresses were
relieved by viscoelastic relaxation of the YBa2Cu307.x/YA103 film. We believe
A10x diffused to the grain boundaries of the YBa2Cu307.x film during the = 1 hour
anneal at 730 C (required to form the orthorhombic YBa2Cu307.x phase),
forming liquid and CuA102 phases. It is probable that incorporation of BaO


88
AES is widely used to determined which elements are present at the surface, with
limited determination of chemical state.
Quantification of the surface concentration is a difficult process because there
are many factors which influence the Auger yield. For an Auger transition from
species i at a site (x,y,z), where N¡ is the background Auger count and dN¡ is the
number of Auger electrons resulting from the transition:114
dN¡ = (incident electron flux of energy Eprimary at x,y,z)
x (ionization cross-section of EA for species i at Ep)
x (backscattering factor for Eprimaiy at the incident direction)
x (probability of decay of EA for species i to give the Auger
transition)
x (probability of no loss escape of electrons from region (x,y,z))
x (acceptance angle of analyzer)
x (instrumental detection efficiency).
To as much of an extent as possible, the Auger operating parameters were
kept constant in these experiments so that comparisons between the atomic
concentrations of different samples could be made. The incident electron flux was
dependent on the beam current, and was maintained at 30 40 nanoamps. The
ionization cross section is heavily influenced by the incident beam energy; low
beam energies are not adequate to produce core holes, and high beam energies
reduce the Auger yield from the shallow core levels. Generally, the optimal beam
energy is 3 5 times the binding energy of the deepest core level of interest. In


64
YBa2Cu307.x phase is formed. At 750 C, the perovskite lattice is
thermodynamically stable at oxygen pressures greater than 150 mTorr.77 Below
this pressure, YBa2Cu307.x decomposes into its component oxides, hence in-situ
growth is dependent on oxygen pressure, and is usually performed at a P02 of
approximately 200 mTorr (figure 2-15). The tetragonal-to-orthorhombic transition
temperature is a function of oxygen pressure, with a maximum temperature of 700
C. This transition is usually induced by backfilling the vacuum chamber to 10 -
760 Torr 02 after the deposition, and slowly cooling to 450 C. The film is kept
at 450 C for approximately 30 minutes to insure oxygenation of the Cu(l) atoms,
and ordering of the 0(1) atoms in the <010> direction. In the tetragonal phase,
0(1) and 0(5) sites are randomly occupied by O atoms, whereas in the
orthorhombic phase the 0(1) sites are completely full and the 0(5) sites are
empty (figure 2-16).78,79
YBa2Cu307_x films deposited in-situ at 650 750 C have higher Jc values than
post-annealed films. In-situ films have a higher ratio of (001)/(100) oriented
grains, and have better in-plane epitaxy in the <100> and <010> directions. It
has been established that the penetration depth is increased, and Jc values are
lowered by weakly coupled grains separated by high-angle grain boundaries or
non-superconducting interfacial phases. Hence the improved electrical properties
of films grown in-situ results from the reduction of weak links at the grain
boundaries.


FLOW CHART FOR Y Cu Si 0
Y,Q, CuO
Y^O-t SiO-7
TERNARY
CuO Si02
L + Y303 + Y2CU2O5
L * Y2Si05 + Y4Si3012
(1900 C)
L ~ Y203 + Y2Si05
(1800 C)
L ** Y4Si3012 + Y2SiOs
(1775 C)
L ** Y2S2O7 + Si02
(1660 C)
L + Y203 ** Y2Cu205 + Y2Si05
Y203 Y2Cu205 Y2Si05
L Y2Cu205 Y2Si05
Figure A-3. Flow chart of the Y Cu Si O system.


239
YBa2Cu307_x films (2500 ) were deposited onto Si substrates using either Y-Zr02
or YA103 barrier layers. The YBa2Cu307.x/Y-Zr02 film became superconducting
at 76 K, and was heavily cracked. However, the superconducting transition
temperature of the YBa2Cu307.x/YA103 film was 78 K, and the cracking was not
as extensive. Scanning Auger analysis showed diffusion of A1 to the interface to
form an Al-Si-O layer, followed by Y-Al-Si-O and Y-Al-O layers. The reduced
cracking in the YBa2Cu307.x/YA103 film deposited on Si was attributed to
relaxation of the thermally induced stresses by viscoelastic strain relaxation at the
YA103/Si interface.
The Jc values of YBa2Cu307.x films deposited on single crystal Y009Zr091O195
substrates cut 5 12 degrees from the (100) planes were significantly enhanced by
depositing a Y-Zr02 barrier layer film prior to the YBa2Cu307.x film. The normal
state properties of YBa2Cu307.x films deposited with or without Y-Zr02 barrier
layers were similar, with superconducting transition temperatures ranging from 86
to 89 K in both cases. However, the Jc for the YBajQ^O^ film deposited
directly onto the Y-Zr02 substrate was 6.8 xlO3 amps/cm2 at 4.5 K, while the Jc
of the YBa2Cu307_x/Y-Zr02 film was lxlO7 amps/cm2 at 4.5 K. The
composition of the Y-Zr02 target was Y013Zr0 87O193, so the increased critical
current density was attributed to pinning of the magnetic flux by excess Y203 in
the grain boundaries. To test this hypothesis, Y203 or Zr02 barrier layer films
were deposited prior to the YBajQrjO^ films. The Jc of the YBa2Cu307.x/Y203
film was lxlO7 amps/cm2 at 4.5 K, while the critical current density of the


231
lowered the liquidus temperature below the temperature suggested by the flow
charts. The glass transition temperature for Y-Al-Si-O is 930 C,135 and SEM
micrographs of a YA103 film deposited on Si at 730 C were featureless.
However, after the deposition and in-situ annealing of the YBajQ^O^ films, the
growth of circular, overlapping grains is apparent. The circular shape of the
grains is evidence that the viscosity of the film was sufficiently low during the
YBa2Cu307_x growth process to adopt a configuration which minimized the surface
tension. The overlap of the YBa2Cu307_x grains resulted from YBajCojO-^ film
contraction, which was the mechanism by which the tensile stresses were relieved.
We conclude that incorporation of A10x into the YBajCujO,.* film was
necessary for viscoelastic relaxation of the thermal stresses. However, the A10x
also contributed to degradation of the superconducting properties of the film
because it degraded intergranular coupling of the superconducting holes at the
grain boundaries.
Al6Si3Q13 Barrier Lavers
Deposition of an Al6Si2013 film onto (100) Si substrates prior to growth of
YA103 or LaA103 barrier layers and the YBa2Cu307_x film dramatically altered
the microstructures of the YBa2Cu307_x films, compared to the cases in which an
Al6Si2013 was not deposited. Figures 4-72 and 4-73 show that the surface
morphologies for the YBa2Cu307.x/(YA103 or LaA103)/Al6Si2013 films on Si are


15
P = fi (2-15)
Kg
where Ag is the wavelength of the electromagnetic signal as it propagates in the
microstrip line. The speed at which electrical signals propagate through the metal
is given by the phase velocity, vp, where
(2-16)
H jiQux and e = e0e where fi0 and e0 are the magnetic permeability and
dielectic permittivity of free space, and (i00)'05 is equal to the speed of light.
Given that nt and er are the relative permeability and dielectric constant of the
microstrip material, the phase velocity of an electromagnetic signal propogating in
the metal is given by:
(ery
0.5
(assuming pr = 1)
(2-17)
Changes in the electrical conductivity, and hence dielectric constant of the metal
as a function of frequency, will also result in phase velocities which vary with
frequency. This phenomena, termed dispersion, causes electrical pulses composed
of various Fourier components to become spread out as they propagate along a
transmission line, and is the primary reason why normal metal conductors are
inadequate for long transmission lines. Dispersion in metallic interconnect lines
makes them inadequate for delay times greater than 1 microsecond, or for
wideband lines carrying short pulses (figure 2-5). For example, assuming typical


93
Table 3-1. Resistivity correction factors as the sample width to probe spacing
increases. The sample length to width is kept constant (a/d = 2).
Ratio of sample width
to probe spacing (= d/s)
Resistivity correction
factor (= F2), assuming
a/d = 2.
1.50
1.4788
1.75
1.7196
2.00
1.9454
2.50
2.3532
3.00
2.7000
4.00
3.2246
5.00
3.5746
7.50
4.0361
10.00
4.2357
15.00
4.3947
20.00
4.4553
40.00
4.5129
00
4.5324


Figure A-7. Phase diagram of the Y203 Si02 system. Reference 137.


7
Josephson supercurrent with induced normal state AC currents will produce sum
and difference harmonic currents within the junction. When the ac Josephson
frequency and radiation frequency are harmonics of each other, there is a mixing
harmonic, and a step in the current vs. voltage spectrum appears at zero
frequency and at the mixing frequency. By varying the voltage and thereby the
supercurrent frequency across the junction, a large range of frequencies can be
sampled via the conversion factor:
1 microvolt = 484 MHz. (2-3)
Surface Quantum Interference Devices (SQUID's) are used to detect weak
magnetic fields, and operate on the principle that the number of flux lines which
pass through a Josephson junction must be quantized.4,9 If the magnetic field
being detected is not a quantized number of fluxons, the Jc of the junction will
change so that the magnetic flux from the magnetic field plus the flux generated
by the current through the junction is quantized (figure 2-2). The magnetic
intensity of a fluxon, is given by:
= = 2.07 x 1015 Webers (2-4)
2e
Thus the I V profile of a Josephson junction is modulated by the magnetic field.
In practice, the magnetic sensitivity of a dc SQUID is significantly improved by
placing two Josephson junctions in parallel (figure 2-3), which makes the loop
defined by the film plus the two junctions the area through which the modulating


Figure 4-58. Auger line scan across crater edge of a YBa2Cu307.i[/YA103 film on (100) Si.
o


65
TEMPERATURE (C)
Figure 2-15. Oxygen partial pressure versus temperature plot showing the critical
stability line for YBajQ^O^ at y = 6.0. Reference 77.


137
for (111) is greater than that of (200) Y-Zr02 (100 vs. 25). The highest
(200)/(lll) ratio was observed for the Y-Zr02 film deposited on (0001) A1203.
Although highly (200) oriented Y-Zr02 barrier layers have been reported to
improve the in-plane epitaxy of YBa2Cu307.x films grown on (102) A1203, the
electrical performance of YBa2Cu307.x/Y-Zr02 deposited on (0001) A1203 was
poor. Raman spectroscopy (figure 4-31) of the YBa2Cu307_x/Y-Zr02 films
deposited on the (1102) and (0001) planes of A1203 show that both films had
minor concentrations of (100) and (010) oriented grains, and there was no
significant difference between the two spectra. Scanning electron micrographs of
a YBa2Cu307.x/Y-Zr02 film on (0001) A1203 revealed small cracks (figure 4-32).
These cracks were not observed in films deposited on (1102) or (1210) oriented
ai2o3.
YBa3Cu3Q7 ./SrTiC^ Films on A13Q3 Substrates
YBa2Cu307.x films deposited directly onto both (100) SrTi03 substrates had
optimal electrical properties (figure 4-33). The normal state resistance was
metallic, and there was a sharp drop to zero resistance at 89 K.
SrTi03 was also very effective as a barrier layer for the growth of YBa2Cu307_x
on (1102) A1203. The orientation adopted by the SrTi03 barrier layer was highly
sensitive to the oxygen pressure during growth. At PG2 = 200 mTorr, SrTi03 grew
with the (100) orientation (figure 4-34), whereas dropping the PG2 to 40 mTorr


My family provided a strong emotional base which allowed my to complete the
program. I would like to thank my parents, Gerhard and Lillian Mueller, for
stressing the importance of finishing a project. I would also like to thank my
brothers, Don, Lloyd, and Keith, my sister Jan, and their families for their
encouragement and for helping to keep things in perspective.
Finally, I would like to thank the friends who made my stay in Gainesville so
enjoyable. I will miss the Sunday afternoon soccer games.
in


139
Figure 4-32. Scanning electron micrograph of a YBa2Cu307.x/Y-Zr02 film on
(0001) ai2o3.


FLOWCHART FOR Y Al Si Cu O
TERNARY
L + Y203 Y4A12
TERNARY
QUATERNARY
TERNARY
09 + Y2Si05
L + Y203 ** Y2Cu205 + Y2SOs
-L + Y203 ** Y2Cu205 + Y4A1209
L + Y203 ** Y4A1209 + Y2Cu205 + Y2Cu2Os
L Y4A1209 Y2Si05 Y2Cu205
L + Y4A1209 o Y2SOs + YAI03
L + Y4A1209 ~ YA103 + Y2Cu2Os
L + Y4A1209 ~ Y2Si05 + Y2Cu205 + YA103
L Y2Si05 Y2Cu2Os YAIO,
L + YA103 o Y3A15012 + y2so5
L + YA103 *> Y3A15012 + Y2Cu2Os
L + YA103 ~ Y3A15012 + Y2SiOs + Y2Cu205
L Y3A15012 Y2Si05 Y2Cu205
L + Y2SiOs Y3A15012 + Y4Si3012
1
L + Y2SiOs o Y4S3012 + Y2Cu2Os
L + Y2SiOs ** Y3A15012 + Y4S3012 + Y2Cu205
L Y3A15012 Y4Si3012 Y2Cu2Os
N
Figure A-5. Flow chart of the Y-Al-Si-Cu-O system.


71
YBa2Cu307.x growth on Y-Zr02 substrates which asserts that YBa2Cu307.x grows
in a manner which permits matching of the oxygen sublattices. This model also
proposes that the large lattice misfits are accomodated by 90 degree boundaries
and stacking faults. Norton et. al.92 reported that YBa2Cu307_x films with T0 = 90
K and Jc = 11,000 amps/cm2 at 77 K were grown on a polished, randomly
oriented Y-ZrOz substrate at 680 C. The films were strongly (001) oriented, and
the electrical properties of the films were more sensitive to the substrate
temperature during growth than were films grown on SrTi03 and LaA103.
Increasing the growth temperature to 730 C resulted in YBa2Cu307.x films with
Jc values of 1000 amps/cm2. The drop in Jc was attributed to increased chemical
interaction at the Y-Zr02/YBa2Cu307_x interface, with subsequent formation of
BaZr03. Presumably, BaZr03 diffused through the grain boundaries and reduced
intergranular conduction.
The lower Jc values for YBajQijO^ films on polycrystalline Y-Zr02 relative
to single crystal Y-ZrOz was attributed to the presence of high angle grain
boundaries. Garrison et al.95 demonstrated that by altering the deposition
conditions, YBa2Cu307.x films could be deposited such that matching of the
YBa2Cu307.x <100>/Y-ZrO2 <100> or <110> directions could be induced. When one
or the other of these orientations was dominant, YBa2Cu307_x films with Jc = 106
amps/cm2 at 77 K were observed. However, if both orientations were present in
the same film, the Jc values were only 102 104 amps/cm2 at 77 K. This


36
presented in order to establish a framework by which film microstructure evolves.
In the specific case of YBa2Cu307.x and barrier layer growth, correlations between
growth conditions and the underlying growth mechanisms will lead to an
understanding of how film quality can be optimized.
The crystallinity and orientation of a film can be manipulated by changing the
growth rate. In order to grow a film, more adatoms must stick to the surface than
are evaporated, hence there must be a flux causing a supersaturation of atoms or
ions reaching the surface and forming stable nuclei. The flux of atoms impinging
on the surface ( atoms ^ P(T)
cm 2 x sec (2nmkT)0S
(2-36)
(
where P(T) is the vapor pressure at the substrate surface, m is the mass of the
impinging species at the substrate surface, T is the substrate temperature, and k is
Boltzmann's constant. To achieve epitaxy, it is essential that the atom be able to
move freely across the substrate until it reaches a potential minimum. The jump
frequency, a), of an adatom on a substrate is given by:
(2-37)
where v = number of jump attempts/sec (typically 1013/sec), and ED is the
activation energy for surface diffusion. The mean stay time for an adatom is:32


213
ohms at 0 K, indicating that grain boundary resistance was negligible compared
to the intrinsic intragranular resistance of YT^Q^O^. The critical current
density for this film was 2.5xl06 amps/cm2 at 4.2 K, which was lower than the
value observed for YBa2Cu307.x films deposited on (1102) LaA103 substrates, but
over 100 times greater than the Jc value of the YBajQ^O^/Y-Zr02 film on
(1102) A1203. The high Jc value for this film indicates there was a large degree of
in-plane alignment between the YI^CujO^ and SrTi03 films, hence the number
of high-angle YBa2Cu307.x grain boundaries were minimal. Raman spectra shows
the 335 cm'1 peak was larger than the 500 cm'1 peak, which indicates the film was
highly (001) oriented, and the asymmetry of the 335 cm'1 peak shows the film was
highly oxygenated. Raman spectra shows the intragranular microstructure of the
film was nearly optimal. The reason the YBa2Cu307_x film grew with a high
degree of (001) texture was attributed to the predominately (100) growth of the
SrTi03 barrier layer. Since SrTi03 has the same perovskite crystal structure and
similar lattice parameters as YBa2Cu307.x (a<100> for SrTi03 = 3.90 , a<100> <010>
= 3.82 and 3.89 for YBa2Cu307.x), YBa2Cu307_x grows epitaxially on single
crystal SrTi03. The Jc for this film was not as high as the Jc observed in
YBa2Cu307.x films deposited on LaA103 substrates, possibly because of the small
amount of (110) oriented SrTi03 grains which could have led to high-angle grain
boundaries in the YBa2Cu307.x film and reduced Jc across these grains. SEM
micrographs show the surface to be composed of clearly defined grains, whereas


INTENSITY (ARBITRARY UNITS)
175
Figure 4-62. X-ray diffraction pattern from a YBa2Cu307.x/Y203 film on Y-Zr02.


44
orientation achieves better row matching by rotating the film surface mesh with
respect to the substrate mesh.
R
Figure 2-10. Matching of atomic rows with epitaxial configurations of dissimilar
rhombic meshes. The substrate unit cell PQRS is drawn in solid
lines, and the overlayer unit cells in dashed lines, (a) Nishiyama-
Wassermand orientation, with overlayer PABC matching rows
parallel to PR, and overlayer PERF matching rows parallel to SQ.
(b) Kurdjumov-Sachs orientation; after rotating the overlayer mesh
relative to the substrate, overlayer PABC matches rows parallel to
PQ and PS. Reference 42.
As the lattice mismatch continues to increase, elastic strain within the film
exceeds the shear stress limit, and strain energy is released by the formation of
misfit dislocations.43 Misfit dislocations usually consist of edge dislocations with
the glide direction parallel to the interface. The elastic strain which can be
accomodated before misfit dislocations form is a function of the film-substrate
bonding, with strong bonding leading to larger elastic strain accomodation and a


264
2000
1700
1400
1100
i 1 1 1 1 1 1 1 r
Li q u i d
Liq.
+
ai2o3
\ CuO ai2o3
\ +
CuO-Al203 -f CuO
i i i i
ALO, 20 40 60 80 CuO
Mol.% i/2cu20
Figure A-10. Phase diagram of the A1203 CuO system. Reference 140.


190
barrier layer structures which contained an Al6Si2013 layer. Conversely, the
normal state behavior of the YBa2Cu307.x/Y203/Al6Si2013 film on (1102) A1203
was semiconducting, and T0 was less than 39 K. X-ray diffraction data (figure 4-
77) show that the YBa2Cu307.x layer was primarily (001) oriented for both
structures, with minor (100) contributions. The predominant barrier layer peaks
in the film with a Y203 barrier layer were from the (222) Y203 phase; there were
no YA103 peaks in the YBa2Cu307.x/YA103/Al6Si2013 film on (1102) A1203. The
Raman spectra for the two films are nearly identical, with major peaks at 335
cm"1, and smaller peaks at 500 cm"1 (figure 4-78).


45
reduced tendency to form misfit dislocations. The strain energy per unit length of
an edge dislocation line is given by:
EgL-ln.*
(2-52)
4ji(1-v) r0
where G is the shear modulus of the film, b is the Burger's vector of the
dislocation, v is Poissons's ratio, R and r0 denote the outer and inner radii of the
strain field over which the dislocation acts. In bulk materials, R can be
approximated as the distance between parallel misfit dislocations, * b/f. In thin
films, the area over which a strain field acts is limited by the film thickness, and if
the film thickness is less than b/f thick, the strain energy of a dislocation can be
substantially reduced.
In films used for electrical applications, one of the most damaging aspects of
misfit dislocations is that they create grain boundaries which degrade the
electrical performance of the films. For pure edge dislocations, the angle between
two adjacent grains, 0, is given by:
sin 0 =
D
(2-53)
where D is the distance between dislocations. Misfit dislocations also reduce the
film-substrate epitaxy. During the initial stages of film growth, islands are highly
mobile and wil try to attain epitaxy with the substrate. If the substrate-film
bonding is weakened by misfit dislocations, there will be a tendency for the
islands to rotate slightly about the ideal in-plane epitaxial positions.52 The extent
of island rotation is given by:


91
values take into account film thickness, probe spacings, and sample geometry. In
this set of experiments, resistance was documented because the variations in
resistance caused by microstructural features overshadowed the relatively minor
changes in resistivity caused by varying YBa2Cu307_x film thicknesses and sample
geometries. An explanation of how resistivities are calculated, and how they are
affected by sample geometries is presented in order to support the hypothesis that
electrical resistance was the more appropriate parameter to monitor. Film
resistivity is given by:
p=FRt (3-5)
where F is a correction factor, R is the measured resistance, and t is the film
thickness. Assuming the probes are equally spaced, there are two correction
factors which will improve the accuracy of the resistivity measurements.115 The
first correction factor is given by the thickness correction factor, F(t/a). As the
sample length, a, becomes significantly greater than the film thickness, F(t/a)
approaches 1. Since the films were less than 3000 thick, and the length of the
substrates were usually greater than 1 cm, F(t/a) had a negligible influence on the
measurement. The second correction factor, F2, is a geometric correction factor
which takes into account the increased current densities in the sample caused by
narrow samples with relatively large distances between the probes. Films
deposited on narrow substrates have large a/d ratios, where d is the sample width.
If we assume the YBa2Cu307_x films were deposited on rectangular substrates in
which the length was twice the width (a/d = 2), the appropriate correction factors


CHAPTER 4
RESULTS
The superconducting properties of YBajQ^O^ films must be nearly optimal
in order to be used in commercial devices. When deposited on substrates which
are chemically inert and have a similar lattice structures, lattice spacings, and
thermal expansion coefficients to YBa2Cu307.x, films with transition temperatures
near 90 K and Jc values greater than 4xl07 A/cm2 at 4.2 K can be grown.
However, when YBajCujO-^ is deposited on Si or A1203 substrates, there are a
number of microstructural defects which degrade the superconductivity, such as
high angle grain boundaries, grain boundary phases, and poor lattice matching
between YBa2Cu307_x and the substrate or barrier layers. Microcracking is a
particularly troublesome defect which results from the tensile stresses caused by
different thermal expansion coefficients between YBa2Cu307_x and the substrate.
The types of defects which are mainly responsible for YBa2Cu307.x film
degradation are often related to the substrate properties, and in this chapter we
will compare and contrast the effectiveness of various barrier-layer materials on
different substrates.
100


CRITICAL CURRENT DENSITY (10 A/CM )
Figure 5-4. Jc versus magnetic field intensity for YBa2Cu307.x/Zr02 film on
Y-Zr02 substrate. Data taken at 4.5 K.


121
in figure 4-17, and Jc versus H is given in figure 4-18. X-ray diffraction shows that
the Y-ZrOz film was predominately (100) oriented, but small (220) peaks were
also apparent (figure 4-19). The only diffraction peaks observed from the
YBa2Cu307.x film were from the (001) orientation. SEM indicates pinholes were
prevalent in the YBa2Cu307_x film. Through the pinholes, a fine-grained Y-Zr02
barrier layer is visible (figure 4-20).
In order to correlate yittria content in the barrier layers with Jc, Zr02 films
without yittria were deposited prior to the YBa2Cu307.x films. Growth of a Zr02
barrier layer on a Y-Zr02 substrate prior to the YBajQ^O^ deposition resulted
in a superconducting film with highly metallic normal state electrical properties,
and a zero resistance temperature of 88.9 K (figure 4-21). The x-ray pattern
from this film was identical to the pattern from the YBa2Cu307_x/Y-Zr02 film
deposited on Y-Zr02 (figure 4-22). However, the Jc was only 9.4 X104 amps/cm2
at 4.5 K. SEM shows the film is comprised of 0.5 -1.0 fim grains (figure 4-23).
The normal state resistance values and x-ray pattern of a YBa2Cu307.x/Zr02
film deposited on (100) SrTi03 were very similar to the YBa2Cu307_x/Zr02 film
deposited on Y-ZrOz. The resistance versus temperature curve was metallic, and
the T0 was 88.1 K (figure 4-24). Only x-ray peaks from the (001) YBa2Cu307.x
phase were detected; no ZrOa peaks were observed (figure 4-25). However, the Jc
of this film (3.8 X106 A/cm2) at 4.5 K was over a factor of ten higher than a


262
Al203-Si02
Figure A-8. Phase diagram of the A1203 Si02 system. Reference 138.


INTENSITY (ARBITRARY UNITS)
119
Figure 4-15. X-ray diffraction patterns from Y-ZrO, and YBa,Cu,07 /Y-ZrO,
films on Y-Zr02.


10
Switching speed (2-6)
2n A (7)
where A(T) is the temperature dependent superconducting energy gap. The
intrinsic switching speed for niobium-based Josephson junctions is 0.22
picoseconds. As an example of the increased performance which can be achieved
by replacing semiconducting switching devices with superconducting Josephson
junctions,10 a four-bit data processor made with GaAs transistors had a clock
speed of 72 MHz, and a power dissipation of 2.2 watts. A processor which
performed the same functions using Josephson junctions had a clock speed of 770
MHz and dissipated 5 milliwatts. More recently, superconducting electronics were
used to fabricate a four-bit shift register using 3 pm linewidths, which operated at
9.6 GHz and dissipated 40 microwatts. By comparison, devices made from GaAs
or Si which used 0.5 pm linewidths, dissipated approximately 100 milliwatts. The
lower power dissipation associated with Josephson switching devices is an
important advantage, since heat generation and removal limit the density and
bandwidth of circuits based on semiconducting electronics.
Superconducting devices employing Josephson junctions have great potential
and, in theory, could completely change the materials and operating principles on
which high speed electonics are currently based. However, reproducible and
reliable Josephson junctions are difficult to fabricate. Efforts to create high
speed, commercially acceptable switching devices based on low temperature
superconducting electronics have not been successful. Because the high
temperature superconductors are much more sensitive to microstructural defects


Abstract of Thesis Presented to the Graduate School
of the University of Florida in Partial Fulfillment of the
Requirements for the Degree of Doctor of Philosophy
MICROSTRUCTURE/ELECTRICAL PROPERTY CORRELATIONS FOR
YBa2Cu307.x/BARRIER LAYER FILMS DEPOSITED ON A1203, SILICON, AND
YITTRIA-STABILIZED ZIRCONIA SUBSTRATES
By
CARL HENRY MUELLER
December, 1992
Chairperson: Professor Paul Holloway
Major Department: Materials Science and Engineering
YBa2Cu307.x and barrier layer films were deposited on single-crystal silicon (Si),
A1203, yittria-stabilized zirconia (Y-Zr02), SrTi03, and LaA103 substrates. A pulsed
laser deposition process was used to deposit the films at a substrate temperature of
730 750 C, and the films were cooled in an oxygen ambient. The films were
characterized using resistance versus temperature, critical current density (Jc), x-ray
diffraction (XRD), scanning electron microscopy (SEM), Auger electron spectroscopy
(AES), and Raman spectroscopy.
Growth of barrier layers on Si and A1203 substrates prior to the superconductor
suppressed chemical interdiffusion between the superconductor and substrate. For
(1102) A1203, the best barrier layer was a SrTi03 film deposited at 200 mTorr of
oxygen. The YBa2Cu307.x film had a zero resistance temperature of 83 K, and the
vi


160
on (1210) A1203 was semiconducting, and there was no transition to the
superconducting state. X-ray diffraction peaks from LaA103 barrier layers
deposited on both A1203 substrates were amorphous (figures 4-50 and 4-51), and
the only diffraction peaks observed were the Ka and Kfi peaks from the substrate.
An explanation is presented in the discussion section chapter as to why LaA103
films grown on (1102) A1203 appeared to be amorphous, whereas films such as
SrTi03, which has a crystal structure similar to LaA103, is highly oriented.
Despite the lack of x-ray diffraction from the LaA103 layers, the orientation of
the YBa2Cu307.x films was dependent on the orientation of the A1203 substrate.
In all cases, the intensities of x-ray diffraction signals from YI^CujO^ films
deposited on LaA103 barrier layers were much smaller than those obtained from
YBa2Cu307.x films deposited on single crystal LaA103 substrates. Raman spectra
of the YBa2Cu307.x/LaA103 films deposited on (102) A1203 and Si substrates
(figure 4-52) show large peaks at 500 cm'1, and shoulder peaks at 440 cm'1, which
indicate a large fraction of non-(001) oriented grains.
YBa^CugO-^/YAlOj Films on Si and A13Q3
YA103 barrier layers effectively limited interdiffusion between YBa2Cu307.x
films and Si or (1102) A1203 substrate. Figure 4-53 compares the resistance vs.
temperature performance of YBa2Cu307.x/YA103 films on Si and (1102) A1203
substrates. The YBa2Cu307.x/YA103 film on (1102) A1203 showed metallic
normal state behavior, with a sharp drop to T0 = 82 K. To compare the


temperature is suppressed by decreasing the oxygen pressure, growth of the
orthorhombic phase will be slow, and rapid cooling of the sample will result in
60
7.0
6.9 h
6.8
6.7
x
O
w 6.6
O
CO 6*5
>-
- 6.4
x
O
6.3
6.2
6.1
6.0
o,
Orthorhombic
<
I
Tetragonal
>-
\
\ T.
\
\
(b)
dOx
dT
\
\
\
\
\ -

l
J
200 400 600 800
TEMPERATURE (C)
1000
Figure 2-13. Total oxygen content, and change in oxygen content as a function of
temperature for YBajQ^O^. Reference 72.


I certify that I have read this study and that in my opinion it conforms to
acceptable standards of scholarly presentation and is fully adequate, in scope and
quality, as a dissertation for the degree of Doctor of Philosophy.
Robert T. DeHoffy
Professor of Materials Science
and Engineering
This dissertation was submitted to the Graduate Faculty of the College of
Engineering and to the Graduate School and was accepted as partial fulfillment of
the requirements for the degree of Doctor of Philosophy.
December, 1992
Winfred M. Phillips
Dean, College of Engineering
Madelyn M. Lockhart
Dean, Graduate School


PULSED LASER DEPOSITION
Vacuum Chamber
Excimer Laser
248 nm, 30 ns pulses
Focusing Lens
50 cm Focal length
Figure 3-1. Schematic diagram of the pulsed laser deposition system.


266
REFERENCES
1. A.W. Kleinsasser and W.J. Gallagher, in Superconducting Devices, edited by
S.T. Ruggiero and D.A. Rudman (Academic Press, Boston, 1990) p. 325.
2. C.M. Chorey, K.S. Kong, K.B. Bhasin, J.D. Warner, and T. Itoh, IEEE
Trans, on Mic. Theory and Tech., 1480 (1991).
3. L.M. Sheppard, Bull. Am. Cer. Soc., 21, 1242 (1992).
4. AC. Rose-Innes and E.H. Rhoderick, Introduction to Superconductivity. 2nd
ed.(Pergamon Press, New York, 1978).
5. B.D. Josephson, Phys. Lett.,1, 251 (1962).
6. T. Van Duzer and C.W. Turner, Principles of Superconductive Devices and
Circuits (Elsevier, New York, 1981).
7. V.Z. Kresin and S.A. Wolf, Fundamentals of Superconductivity. (Plenum
Press, New York, 1990).
8. H. Hayakawa, in Superconducting Devices, edited by S.T. Ruggiero and
DA. Rudman (Academic Press, Boston, 1990), p. 101.
9. Semiconductor International, p. 16, November, 1991.
10. J. Clarke in Superconducting Devices, edited by S.T. Ruggiero and DA.
Rudman (Academic Press, Boston, 1990), p. 51.
11. D.R. Tilley and J. Tilley, Superfluidity and Superconductivity (John Wiley
and Sons, New York, 1974).
12. FA. Miranda, W.L. Gordon, K.B. Bhasin, V.O. Heinen, and J.D. Warner, J.
Appl. Phys. 2Q, 5450 (1991).
13. F. Wooten, Optical Properties of Solids (Academic Press, New York, 1972).
14. T.C. Edwards, Foundations for Microstrip Circuit Design (John Wiley and
Sons, New York, 1981).
15. O.K. Kwon, B.W. Langley, R.F.W. Pease, and M.R. Beasley, IEEE Elec.
Dev. Lett. £, 582 (1987).
16. H.Y. Lee, and T. Itoh, IEEE Trans. Mic. Theory and Tech. 22, 1904 (1989).


MICROSTRUCTTJ RE/ELECTRICAL PROPERTY CORRELATIONS FOR
YBa2Cu307.x/BARRIER LAYER FILMS DEPOSITED ON A1203, SILICON, AND
YITTRIA-STABILIZED ZIRCONIA SUBSTRATES
By
CARL HENRY MUELLER
A DISSERTATION PRESENTED TO THE GRADUATE SCHOOL
OF THE UNIVERSITY OF FLORIDA IN PARTIAL FULFILLMENT
OF THE REQUIREMENTS FOR THE DEGREE OF
DOCTOR OF PHILOSOPHY
UNIVERSITY OF FLORIDA
1992


52
In brittle materials, film cracking is caused by the motion and accumulation of
dislocations, which creates regions of high elastic energy and ultimately fracture.
The film microstructure can significantly alter the magnitude of stress required to
move dislocations and initiate plastic deformation.46,47,51 Numerous experiments
have shown that the stress at which plastic deformation is initiated is greater in
thin films than in bulk materials. This is because dislocations are pinned at the
film-substrate interface, thus increasing the stress required to move dislocations
through the film. The minimum stress required to move a dislocation in a film is
given by:48
a, = A [ ] [ ln(^-)]
f l2n(l-Vf)hi l(Vf+v4) b
(2-59)
where A is a geometrical constant which accounts for the angle between the
applied stress and Burgers vector, b is the Burgers vector of the dislocation, vf and
vs are the elastic shear moduli of the film and substrate, and h is the film
thickness. For brittle films, the stresses required to move dislocations are
associated with the onset of microcracking. Equation 2-59 shows that the stress
required to move dislocations is inversely proportional to the film thickness, and is
the reason why there is often a critical film thickness which, if exceeded, causes
cracks to form.
Depositing a second layer on top of the initial film significantly increases the
stress in the initial layer. This seems to contradict the elastic continuum model,
which asserts that the stress in a given layer is only a function of the differences in


154
YBa3Cu307j,/LaAlQ3 Films on Si and A13Q3 Substrates
The optimal normal state and superconducting properties of YBajQ^O^*
films deposited on (1102) LaA103 substrates have been described earlier. LaA103
was considerably less effective as a barrier layer material than as a substrate
material, and YBa2Cu307.x/LaA103 films deposited Si were severely degraded.
The normal state resistance for YBa2Cu307_x/LaA103 on Si was high and
semiconducting, and there was a broad transition to the superconducting state,
from Tonset = 90 K to T0 less than 46 K (figure 4-45). AES data do not show
La-Si interdiffusion or formation of an interfacial phase at the substrate/barrier
layer interface (figure 4-46), and there are no well defined impurity phases at the
YBa2Cu307.x/barrier layer interface. Diffraction peaks from the rare-earth silicide
or any other impurity phases were not observed in the x-ray pattern (figure 4-47).
The YBa2Cu307_x film was weakly (001) oriented, and a smaller (103) peak was
also observed. The film surface was very rough (figure 4-48), and an
interconnected crack network was observed throughout the film.
LaA103 was also a poor barrier layer material for YBajQ^O^ films deposited
on the (102) or (1210) faces of A1203. The normal state resistance values of the
YBa2Cu307.x/LaA103 film on (1102) A1203 were moderately metallic, with a T0 of
58 K (figure 4-49). The resistance values of the YBa2Cu307.x/LaA103 film


INTENSITY (arbitrary units)
Figure 4-44. Raman spectra of YBa2Cu307_x/SrTi03 films on (1210) and (102) A1203 substrates.
LA
LO


Figure 4-65. X-ray diffraction pattern from a YBa2Cu307.I/Y303 film on (100)
INTENSITY (ARBITRARY UNITS)


67
Barrier Layer Technology
Because of the increased performance which could be realized by replacing
metallic interconnects and microstrip lines with superconductors, significant effort
has been expended towards finding deposition techniques which enable
YBa2Cu307.x films with high T0 and Jc values to be deposited on silicon and A1203
substrates. Sapphire is an attractive substrate material for microwave applications
because it is relatively inexpensive, mechanically strong, and has a much lower
dielectric loss tangent (tan <5) than other substrate materials (table 2-3).17
Table 2-3. Dielectric properties of substrate and barrier layer materials.
Reference 17
Material
Dielectric
constant
Loss
tangent
Sapphire (A1203)
9.4
1x10*
Silicon (Si)
12
1 x 103
Y-Zr02
27
6x 104
LaA103
25
5.8 x 104
SrTi03
305
MgO
9.65
4x 104


263
Figure A-9. Liquidus projection of the Y203 A1203 Si02 system. Reference
139.


COMPLEX CONDUCTIVITY (S/m)
149
Figure 4-40. Real (ctj) and imaginary (a2) parts of the electrical conductivity for a
YBa2Cu307.x/SrTi03 film on (1102) A1203. Data taken at 36
gigahertz.


CHAPTER 2
LITERATURE REVIEW
Devices
Much of the interest in high temperature superconductivity is due to the large
number of applications which could benefit from replacing normal metals and
conventional solid state electronic devices with superconducting materials.
Basically, applications which could utilize high temperature superconducting thin
films can be divided into two categories: active devices which require one or
more weak link junctions, and passive devices which utilize the intrinsic properties
of superconductors for enhanced performance.3
In superconductors, zero resistance and macroscopic quantization occur
because the wave functions of the electron or hole pairs are coherently coupled to
each other. This coherence enables calculations of the phase difference between
two different points of the superconductor to be made.4 Weak-link junctions are
thin insulating or normal metal regions which separate two superconducting
volumes, and the lower density of superconducting charge carriers causes a
gradient in the phase of the superconducting electron (or hole) wave function
across the weak link.
4


L + YA103 Y3A15012 + Y2Si05
YA103 Y3A15012 Y2SiOs
L Y3A15012 Y2Si05
L + Y2Si05 Y3A15012 + Y4Si3012
Y2Si05 Y3A15012 Y4Si3012
L Y3A15012 Y4Si3012
L ** Y3A15012 + A1203
(1760 C)
L + Y3A15012 ** Y4Si3012 + A1203
Y3A15012 Y4Si3012 AI203
L Y4Si3012 A1203
Figure A-l. Continued


95
fi=£ (3-7)
r 3
when the plane of the film is circular. Hp is the strength of the applied magnetic
field which completely penetrates the film. The polarity of Ha is then reversed,
and the magnitude is increased beyond -Hp. By cycling the sample through the
applied magnetic fields, a B vs Ha hysteresis loop is generated. The Bean model
asserts that Jc is constant throughout the sample, so
j = E (3-8)
e dr
where r is the radius of the film. Combining equations 3-7 and 3-8, then
integrating:
J = (3-9)
c 2 r
where AM is the remnant magnetization per unit volume (emu/cm3), and is the
magnetization difference for increasing and decreasing fields. The radius is given
in centimeters. After correcting for the units (1 emu/cm3 = 10 amps/cm), the
final expression is obtained:
j = iM (3-10)
c r
where Jc is given in amps/cm2. In practice, rectangular samples were cut, and the
films were oriented with the (001) planes perpendicular to Ha. r(cm) was the
radius of the largest circle which could be inscribed in the rectangular sample.


Si02 A1203 CuA102 CuO
is)
Lf
VO
Figure A-5. Continued.


YBajQ^CVx films deposited on this structure were highly (001) oriented. The Jc
was lxlO4 amps/cm2 at 4.2 K.
YBa3Cu3Q7l[/Y3Q3 Films on Si. Y-ZrQ3. and SrTiQ3 Substrates
Y203 barrier layers dramatically improved the Jc values of YBajQ^O^ films
deposited on Y-Zr02 substrates. The normal state resistance of the YBa2Cu307_x/
Y203 film on Y-Zr02 was metallic, with a T0 of 89.6 K (figure 4-61). The Jc
versus H data was similar to that of the YBa2Cu307.x/Y-Zr02 film on Y-Zr02,
with zero field Jc values of 7x10s A/cm2 at 77 K, and lxlO7 A/cm2 at 4.5 K.
X-ray diffraction of the YBa2Cu307_x/Y203 film on randomly oriented Y-ZrOz
show small peaks from the (001) YBa2Cu307_x phase (figure 4-62). The peak
intensities from this film were much smaller than the peak intensities from the
YBa2Cu307.x/Y-Zr02 film deposited on Y-Zr02. SEM shows this film is fine
grained, with numerous pinholes (figure 4-63).
The normal state resistance and Jc data for a YBa2Cu307.x/Y203 film
deposited on (100) SrTi03 (figures 4-64) were very similar to those observed on
Y-Zr02 substrates. The Jc values of the YBa2Cu307.x/Y203 film on (100) SrTi03
were 7x10s A/cm2 at 77 K, and lxlO7 A/cm2 at 4.5 K. X-ray diffraction
patterns from this film show the Y203 barrier layer was highly (100) oriented, and
only YBa2Cu307.x peaks from the (100) orientation were detected (figure 4-65).
SEM micrographs show the YBa2Cu307_x/Y203 film is smooth (figure 4-66).


189
diffraction patterns for both of these films show strongly (001) oriented
YBa2Cu307.x peaks (figure 4-74). The T0 for the YBa2Cu307_x/YA103/Al6Si2013
film on Si film is significantly lower than the YBa2Cu307_x/YA103 film on Si (61.5
versus 78 K), whereas T0 is higher for the YBa2Cu307.x/LaA103/Al6Si2013 film
on Si than for the YBa2Cu307_x/LaA103 film on Si (59 versus 46 K). In addition,
the normal state resistance values in the film containing the LaA103/Al6Si2013
barrier layer were a factor of 10 lower than in the film with only a LaA103 barrier
layer. Although the electrical and x-ray diffraction data indicate similar
microstructures for the YBa2Cu307_x/(LaA103 or YA103)/Al6Si2013 films
deposited on Si, the Raman spectra for these films are quite different (figure
4-75). The 500 cm"1 peak of the film with a LaA103 layer is smaller than the 335
cm"1 peak, which indicates a high degree of (001) orientation. By contrast, the
500 cm"1 peak in the film with a YA103 barrier layer is more pronounced.
Comparing these data to the Raman spectra obtained from the YBa2Cu307.x
/YA103 film on Si, we see that addition of the Al6Si2013 layer resulted in more
(100) and (010) oriented YBajC^O^ grains.
The effectiveness of YA103/Al6Si2013 relative to Y203/Al6Si2013 as a barrier
layer structure to (1102) A1203 presents a unique comparison, since both are
comprised of the same elements, but form different phases. The resistance vs.
temperature data (figure 4-76) show the normal state behavior of theYBa2Cu307.x
/YAlcyAlgSiAa film on (1102) A1203 was metallic, with a T0 = 75 K. This
was the highest transition temperature observed for YBa20u3O7.x films grown on


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AUTHOR:
TITLE:
Mueller, Carl
Microstructure/electrical property correlations for YBa2Cu307-
x/barrier layer films deposited on A1203 silicon and yittria-
stabilized zirconia substrates / (record number: 1904702)
PUBLICATION
DATE:
1992
I) ff- )0liA.*-/bsl
, as copyright holder for the
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INTENSITY (ARBITRARY UNITS)
YBa2Cu307.x/Y203/AI6Si2013 on Al203 (1102)
Figure 4-77. X-ray diffraction patterns for YBajO^O-^/YAlCVAl^O^ and
YBa2Cu307.x/Y203/Al6S2013 films on (1102) A1203.


187
2@tu X8600 0004 1.0U UFMSE
Figure 4-72. Scanning electron micrograph of a YE^Q^O^/YA103/Al6Si2013
film on (100) Si.


L <* Y-yCu7Or + CuO
(1060 C)
L ** Y2Cu205 + CuO 4- Si02
L ** CuO + Si02
(1050 C)
Y2Cu205 CuO Si02
N>
Ut
w
Figure A-3. Continued.


227
Although Tonset was 88 K, the transition was very broad and 0 ohm resistance
was not observed. X-ray diffraction showed only a weak LaA103 peak, and the
YBa2Cu307.x film grew with (103) and (001) orientations. All of the YBa2Cu307_x
peaks were weak, which indicates non-crystallized Y-Ba-Cu-O constituents. The
weak x-ray peaks may have also resulted from a crystalline YBa2Cu307_x film in
which the grain size was too small to yield x-ray diffraction peaks. Although the
LaA103 barrier layer films did not yield well defined x-ray peaks, the orientation
of the YBa2Cu307.x films was dependent on the orientation of the A1203
substrates. This dependence supports the hypothesis that the LaA103 barrier
layer and YBa2Cu307_x films were more crystalline than the x-ray diffraction
suggests, but the very small grain size broadened the x-ray diffraction peaks so
they were undetected. Like the YBa2Cu307_x/LaA103 film deposited on (1102)
A1203, the wide transition temperature range was attributed to intergranular
damage caused by La diffusion to the grain boundaries.
YBa2Cu307_x/LaA103 films deposited on (100) Si had very poor normal state
properties. SEM micrographs show the film was rough, and was separated into
isolated regions by an interconnected crack network. There are two primary
reasons why this film was heavily cracked. First, the large difference in thermal
expansion coefficients between YBa2Cu307_x and Si induced tensile stresses which
exceeded the fracture stress of YBa2Cu307_x. Second, cracking was more extensive
in this film than in the YBa2Cu307.x/Y-Zr02 film on (100) Si because the
YBa2Cu307.x film contained mixed orientations, and the non-(001) oriented grains


5
A Josephson junction is a special type of weak link junction in which the
relationship between the supercurrent, Is, and the gradient of the phase of the
superconducting wave function (A) is given by:5,6
Ia = Jc sin (A). (2-1)
where Ic is the maximum superconducting current which can travel through the
junction. For the rest of the discussion on devices, it will be assumed that the
weak links are Josephson junctions. The current vs. voltage behavior of a
Josephson junction is characterized by a zero resistance supercurrent which
persists when Ad> is constant, and a non-linear voltage which appears when AO
varies with time. The voltage, V, across a Josephson junction in this state is given
by:
2e7=_ft_li^) (2-2)
2tc dt
where e is the carrier charge and h is Planck's constant.
The I-V plot for a Josephson junction is given in figure 2-1,7 and virtually
all of the active superconducting devices utilize the behavior predicted by
equations 2-1 and 2-2. A brief overview of some of the devices which use
Josephson junctions, and their operating principles, is presented below.
Oscillators are devices which generate a repeating voltage vs. time waveform.
Superconducting oscillators, in which the output voltage is controlled by the
frequency of the supercurrent travelling through the junction (equation 2-2), can


87
35 nanoamps beam current, 132 volts emission voltage, and 40 60 microamps
emission current. Interdiffusion at the YBajCuaO^/barrier layer and barrier
layer/substrate interfaces was observed by ion sputtering a crater in the films, so
as to expose the barrier layer and substrate, then making a line scan across the
edge of the crater. With this method, errors induced by electrical charging and
Auger peak shifting in the insulating barrier layers were minimized.
Auger is often used to qualitatively measure the relative concentrations of
components within the top 10 of the surface. The type of element is
determined by the energy of electrons emitted as a result of the Auger process.
The Auger process is started by an incident electron beam with sufficient energy
to remove an inner shell electron, which creates a core hole. The ion energy is
reduced by filling the core hole with an electron from a more shallow energy
level, and emitting another electron from a shallow energy level. The energy of
the emitted electron is given by:113
Kinetic Energy = EA EB Ec (3-4)
where EA is the energy of the core level electron, EB is the energy of the shallow
level electron which fills the core hole, and Ec is the energy of the shallow level
electron which is emitted. Although Ea, Eb and Ec are all sensitive to the
chemical state of the atom, the time constant for Auger emission is short, so the
peaks are broad. Therefore the energies of electrons emitted via the Auger
process are less sensitive to the chemical state of the element than are electrons
emitted by other techniques, such as x-ray photoelectron spectroscopy. Hence


INTENSITY (ARBITRARY UNITS)
130
YBa2Cu307.x/Zr02
ON (100) SrTi03
A Y Ba2Cu3C>7.x
s
CO
o
o
CM
O
O
Mm
I
*+*++**
P
O
O

X
X
X
10
20 30 40 50
DIFFRACTION ANGLE (20)
60
Figure 4-25. X-ray diffraction pattern from a YBa2Cu307_x/Zr02 film on (100)
SrTi03.


145
Figure 4-37.
0 0 %
4
IR *
20KU X8600 |000 1.0U UFMSE
Scanning electron micrograph of a (001) YBa2Cu307x/(100) SrTi03
film on (1102) A1203.


142
during the deposition resulted in a highly (110) oriented SrTi03 film (figure 4-
35). YBa2Cu307.x films subsequently deposited on these barrier layers adopted
orientations which minimized lattice misfit at the YBajQ^O^/barrier layer
interface. YBa2Cu307_x films on (100) SrTi03 barrier layers were (001) oriented,
while (103) oriented YBa2Cu307.x films grew on (110) SrTi03 barrier layers. The
electrical properties of these films were sensitive to the YBajQ^O^ orientations
(figure 4-36). The normal state properties of an (001) YBa2Cu307_x /(100) SrTi03
film deposited on (1102) A1203 were similar to YBa2Cu307.x deposited directly on
(100) SrTi03, with a metallic normal state resistance, but a lower zero resistance
temperature (83 K). The Jc value for this film was 2.5xl06 amps/cm2 at 4.5 K.
SEM micrographs of the (001) YBa2Cu307_x/(100) SrTi03 film on (1102) A1203
show a large grained microstructure with clearly visible grain boundaries (figure 4-
37). By contrast, the surface of a YBa2Cu307.x film deposited on (100) SrTiOs
was featureless.
Millimeter-wave transmission data taken at 36 gigahertz shows that the
surface resistance of the YBa2Cu307.x/SrTi03 film deposited on (1102) A1203 was
approximately 10 milliohms at 4.2 K (figure 4-38), which is lower than the value
reported for a YBa2Cu307.x film grown on single-crystal LaA103 (50 milliohms).
Figure 4-39 shows a sharp drop in the phase angle as the film went from the
normal to the superconducting state. The imaginary part of the conductivity,


73
YBa2Cu307.x (13.6x 10"6/C). Chourasia et al." showed that in a YBajCi^O^
film deposited directly onto a Si substrate at room temperature, Si diffused into
YBa2Cu307.x near the interface, and formed a Si suboxide which depleted oxygen
from the CuO planes. After annealing the film at 860 C for 3 hours, the Si
diffusion was much more extensive, and Si02 was detected at the surface. The
authors concluded that oxidation of Si at the expense of CuO was the primary
reason that diffusion of Si into YBajQijO^ degraded the superconducting
properties of YBa2Cu307.x.
Better films have been grown directly on sapphire. YBa2Cu307_x films grown
on (1102) A1203 at 700 C showed T0 = 87 K and Jc values of 2xl06 amps/cm2
at 4.2 K.100 Attempts to deposit films at higher temperatures resulted in rapid
deterioration of the film because of interfacial reactions, probably BaO and CuO
reacting with A1203.
Because of the enormous technological advances which would result from
successful deposition of high quality YBa2Cu307.x films on Si and sapphire, a great
deal of research has been dedicated towards overcoming the interfacial reaction
problem. The most common approach has been to grow an intermediate barrier
layer prior to YBa2Cu307_x. The principal requirements of the barrier layer are
that it must be chemically inert to both the substrate and the YBa2Cu307.x film,
and should lattice match YBa2Cu307.x and the substrate. The materials which
have been most successful as barrier layers have been the materials which are also
the best substrate materials, such as SrTi03, Y-Zr02, and MgO. Although


12
X
Figure 2-4. Variation of magnetic flux density at the boundary of a
superconductor. Reference 4.


41
which case the film completely covers the substrate before another layer is
nucleated, and the film grows at a uniform rate normal to the substrate. Films
can also grow by a three-dimensional growth process, in which case the film atoms
agglomerate into islands and the islands coalesce to form the film. A third option
is for the film to grow by a mixture of two and three-dimensional growth. The
first growth mode, termed Frank-van der Merwe, is characterized by two
dimensional growth of the film. This occurs when the surface energy of the film
is less than that of the substrate and interface:
film ^ substrate + substrate-film interface ' '
In the Frank-van der Merwe growth mode, the surface free energy of the system
is lowered by replacing the substrate surface with a different surface which has a
lower surface energy. For most electronic applications, in which it is often
desirable to grow a thin film with the same structure and orientation as the
substrate, two dimensional growth is preferred because the adatoms will tend to
reside in potential minima along the substrate, thus preserving the pattern of the
substrate surface. The relative adatom-adatom or adatom-substrate bonding
energies also influence the film growth mode, with strong adatom-substrate
bonding favoring two dimensional growth. When the surface energy of the film is
high,


INTENSITY (arbitrary units)
Figure 4-5. Raman spectra for YBa2Cu307.x film deposited on (1102) LaA103.


11
than their low temperature counterparts, it is unlikely that high temperature
superconductors will be used as switching devices in the near future. While
devices which utilize fewer Josephson junctions, such as electromagnetic detectors
and SQUID's are more feasible, difficulties in creating reliable Josephson
junctions remain formidable.
The materials challenges presented by the second group of devices based on
superconducting materials, passive devices, are more likely to be surmounted in
the near future. The first commercial applications for high temperature
superconductors will probably emerge from this group.3 The largest applications
of passive superconductors will be as transmission lines, delay lines, and filters for
receiving and transmitting electromagnetic signals.
The advantages gained by replacing normal metals with superconducting
materials are derived from the magnetic field penetration depth:11
X(T) =
MO)
[!-(-£-) 3
TV,0.5
(2-7)
where A(T) is the temperature dependent magnetic penetration depth, A(0) is the
penetration depth near 0 K, T is the operating temperature, and T0 is the
highest temperature at which zero resistance is observed. The response of a
superconducting material to an applied magnetic field is shown in figure 2-4. For
a superconductor, A is a material property, so it is independent of the device
operating frequency.


70
cubic perovskite materials (such as SrTi03) is that the angle between the <100>
and <001> directions is not 90 . YBa2Cu307.x films with T0 values of 92, 89,
and 88 K have been deposited on single crystal LaGaO^88 PrGa03,89 and
YbFeCXj90 substrates, respectively. In each case, the YBa2Cu307_x films were
strongly (001) oriented. Although each of these substrates contain elements which
are known to suppress superconductivity (La, Pr, and Fe), interdiffusion was not
observed.
Despite a large lattice mismatch, highly (001) oriented YBajQ^O^ films with
T0 = 92 K have been deposited on UNb03 substrates.91 The unit mesh of Y-cut
LiNb03, onto which the YBa2Cu307_x films were grown, is 5.148 x 6.932 . The Jc
values for these films were 2x10s amps/cm2 at 77 K, and the reduction in Jc
(relative to YBa2Cu307.x films deposited on SrTi03) was attributed to the high
concentration of Li in the film. The diffusion of Li was so rapid that the
concentration of Li within the YBa2Cu307_x film film was peaked at the surface.
The second group of substrates are materials which are chemically inert to
YBa2Cu307.x, but do not lattice match nor do they have the perovskite lattice. Y-
Zr02 amd MgO have been widely studied because high quality YBa2Cu307.x films
have been grown on these materials. Despite the large YBa2Cu307.x/Y-Zr02
lattice mismatch of 5.95 % in the YBajQijO^ <110>/Y-ZrO2 <100> direction,
YBa2Cu307_x films with T0 = 88 K and Jc = 1X106 amps/cm2 at 77 K were
deposited on (100) Y-Zr02.92,93 Tietz et. al.94 proposed a model for


INTENSITY (ARBITRARY UNITS)
250 300 350 400 450 500 550 600
RAMAN SHIFT (cm'1)
Figure 4-6. Raman spectra from 500 and 1000 YBa2Cu307_x films on (1102) LaA103, and from bare LaAlO
substrates.


INTENSITY (ARBITRARY UNITS)
134
Figure 4-28. X-ray diffraction patterns from Y-Zr02, and YBa2Cu307.x/Y-Zr02
films on (1102) A1203.


271
80. T. Venkatesan, E.W. Chase, X.D. Wu, A. Inam, C.C. Chang, and F.K.
Shokoohi, Appl. Phys. Lett. 51, 243 (1988).
81. R.W.G. Wycoff, Crystal Structures (Interscience, New York, 1965) Vol. 2.
82. S. Geller and V.B. Bala, Acta Cryst. % 1019 (1956).
83. R.K. Singh, J. Narayan, A.K. Singh, and J. Krishnaswamy, Appl. Phys. Lett.
2271 (1989).
84. B. Roas, L. Schultz, and G. Endres, Appl. Phys. Lett. 51, 1557 (1988).
85. D.M. Hwang, T. Venkatesan, C.C. Chang, L. Nazar, X.D. Wu, A. Inam, and
M.S. Hegde, Appl. Phys. Lett. 4, 1702 (1989).
86. R.E. Muenchausen, S.R. Foltyn, X.D. Wu, R.C. Dye, N.S. Nogar, A.H.
Carim, F. Heidelbach, D.W. Cooke, R.C. Taber, and R.K. Quinn, in Proc.
SPIE Vol. 1394 (1990), ed. by R. Singh, J. Narayan, and D.T. Shaw. p. 221.
87. C.W. Nieh, L. Anthony, J.Y. Josefowicz, and F.G. Krajenbrink, Appl. Phys.
Lett. 6, 2138 (1990).
88. G. Koren, A. Gupta, E.A. Giess, A. Segmuller, and R.B. Laibowitz, Appl.
Phys. Lett. M, 1054 (1989).
89. M. Sasaura, M. Mukaida, and S. Miyazawa, Appl. Phys. Lett. 51, 2728
(1990).
90. R. Ramesh, A, Inam, W.A. Bonner, P. England, BJ. Wilkens, BJ. Meagher,
L. Nazar, X.D. Wu, M.S. Hegde, C.C. Chang, T. Venkatesan, and H.
Padamsee, Appl. Phys. Lett. 1138 (1989).
91. S.G. Lee, G. Koren, A. Gupta, A. Segmuller, C.C. Chi, Appl. Phys. Lett. 55L
1251 (1989).
92. D.P. Norton, D.H. Lowndes, J.D. Budai, D.K. Christen, E.C. Jones, K.W.
Lay, and J.E. Tkaczyk, Appl. Phys. Lett. 51, 1164 (1990).
93. D.M. Hwang, Q.Y. Ying, and H.S. Kwok, Appl. Phys. Lett. 51, 2429 (1991).
94.L.A. Tietz, C.B. Carter, D.K. Lathrop, S.E. Russek, R.A. Buhrman, and J.R.
Michael, J. Mater. Res. 4, 1072 (1989).


202
SrTi03 substrates shows the same H-0-5 dependence at 4.5 and 77 K (figures 5-2
and 5-3).
The Jc versus H behavior of YBa2Cu307_x/Zr02 films deposited on off-axis Y-
Zr02 and SrTi03 substrates are shown in figures 5-4 and 5-5. The shape of the
curves are similar to those reported by others for bulk and thin film samples with
Josephson weak links.123,124 The shape of Jc versus Ha curves for the YBa2Cu307_x
/Zr02 films, and lower Jc values (relative to the films with Y203 or Y-Zr02
barrier layers) indicates the Jc was limited by decoupling of Josephson weak links
at the grain boundaries. Since Zr02 is monoclinic, there could not be epitaxial
lattice matching at the YBa2Cu307.x/Zr02 interface. The Jc values of the
YBa2Cu307.x/Zr02 films deposited on (100) SrTi03 are close to the highest values
predicted for films with aligned (but not epitaxial) <100> and <010>
YBa2Cu307.x directions (2xl06 amps/cm2 at 4.5 K),125 and the Jc of the
YBa2Cu307.x/Zr02 film on Y-ZrOz is close to the maximum value predicted for
(001) films with random in-plane orientations (3 x10s amps/cm2 at 4.5 K).92 The
Jc and resistance versus temperature data indicate high-angle grain boundaries
were the primary reason for lower Jc values.
Presumably, the high Jc values observed in the YBa2Cu307.x/Y203 film on
SrTi03 could be attributed to epitaxial growth of both the Y203 and YBa2Cu307.x
films, which would limit the fraction of high-angle grain boundaries. Whether
highly (100) oriented Y203 was also deposited on Y-Zr02 substrates is not clear,
since the substrates were cut off-axis from the (100) planes, and diffraction peaks


234
film influences the stress in the subsequently deposited barrier layer and
YBa2Cu307.x films presents an interesting dichotomy. According to the continuum
mechanics model for thin films, one would expect the Al6Si2013 film to have a
negligible impact on the stress in the other films. The small expected effect of
the Al6Si2013 film would be to reduce the total tensile stress in the multilayered
film structure.
There were many similarities between YBa2Cu307_x/(YA103 or
Y203)/Al6Si2013 films deposited on (1102) A1203. Both barrier layer structures
were composed of the same components, hence differences in degradation
resulting from incorporation of different types of cations in the YBa2Cu307.x film
were minimized. Both YBa2Cu307.x films were predominately (001) oriented, yet
the resistance vs. temperature data for these films were vastly different. The
YBa2Cu307.x/YA103/Al6Si2013 film deposited on (1102) A1203 was metallic, with
a T0 = 75 K. The normal state resistance values and T0 for this film indicated
that incorporation of Al6Si2013 into the barrier layer structure slightly degraded
the performance of the YBa2Cu307.x film.
Degradation of the YBa2Cu307_x/Y203/Al6Si2013 film on (1102) A1203 was
much more extensive. The normal state resistance behavior of this film was
semiconducting, with a long transition to the superconducting state at T0 less than
39 K. An initial explanation for the degradation of the YBa2Cu307.x film was
that the growth of the Y203 was three dimensional, so pathways existed by which
Si could diffuse to the YBa2Cu307.x film. However, Raman spectra showed that


13
In microwave and millimeter-wave transmission lines, there is an electric field
induced in the superconducting line because of the inertia of superconducting
electron pairs to the electromagnetic field. A surface resistance results because
the normal-state electrons are excited by the electric field, and the surface
resistance is qualitatively a measure of how much of the electromagnetic signal is
lost as heat in the transmission line:6
Joule losses
Surface resistance = Rs =
(2-8)
where Hsurface is the magnetic field intensity inside the superconductor. In
superconductors, the depth to which a magnetic or electric field can extend is
limited by the penetration depth, and the surface resistance (Rs) is given by:6,12
(2-9)
where A is a constant, o) is the frequency of the electromagnetic signal, T is the
temperature, kB is Boltzmann's constant, and A(T) is the superconducting energy
gap.
By contrast, the surface resistance of normal metals is given by:13
(2-10)
where fiQ is the magnetic permeability of the metal, and o0 = al + ia2 is the
complex electrical conductivity, with,


272
95. S.M. Garrison, N. Newman, B.F. Cole, K. Char, and R.W. Barton, Sppl.
Phys. Lett. 50,2168 (1991).
96. Q. Li, O. Meyer, X.X. Xi, J. Geerk, and G. Linker, Appl. Phys. Lett. 55. 310
(1989).
97. J.T. Cheung, I. Gergis, M. James, and R.E. DeWames, Appl. Phys. Lett. 60.
3180 (1992).
98. T. Venkatesan, E.W. Chase, X.D. Wu, A. Inam, C.C. Chang, and F.K.
Shokoohi, Appl. Phys. Lett. 53, 243 (1988).
99. A.R. Chourasia, D.R. Chopra, A.H. Bensaoula, and P. Ruzakowski, J. Vac.
Sci. Technol. A 1& 115 (1992).
100. K. Char, D.K. Fork, T.H. Geballe, S.S. Laderman, R.C. Taber, R.D.
Jacowitz, F. Bridges, G.A.N. Connell, and J.B. Boyce, Appl. Phys. Lett. 56,
785 (1990).
101. A.E. Lee, C.E. Platt, J.F. Burch, R.W. Simon, J.P. Goral, and M.M. Al-
Jassim, Appl. Phys. Lett. 52, 2019 (1990).
102. E.J. Cukauskas, L.H. Allen, R.T. Holm, and G.K. Sherrill, Appl. Phys. Lett.
60, 389 (1992).
103. D.K. Fork, D.B. Fenner, G.A.N. Connell, J.M. Phillips, and T.H. Geballe,
Appl. Phys. Lett. 52, 1137 (1990).
104. H.M. O'Bryan and P.K. Gallagher, Ad. Cer. Mat. 2, 610 (1987).
105. D.K. Fork, F.A. Ponce, J.C. Tramontana, N. Newman, J.M. Phillips, and
T.H. Geballe, Appl. Phys. Lett. 5S, 2432 (1991).
106. E. Wiener-Avnear, G.L. Kerber, J.E. McFall, J.W. Spargo, and A.G. Toth,
Appl. Phys. Lett. 56, 1802 (1990).
107. S. Mima, H. Tsuge, T. Yoshitake, S. Matsubara, T. Satoh, Y. Miyasaka, and
N. Shohata, Proc. First Int. Symp. Superconductivity, p. 539 (1988).
108. X.D. Wu, A. Inam, M.S. Hegde, B. Wilkens, C.C> Chang, D.M. Hwang, L.
Nazar, T. Venkatesan, S. Miura, S. Matsubara, Y. Miyasaka, and N. Shohata,
App. Phys. Lett. 54, 754 (1989).


200
diffraction showed that the fully oxygenated YBa2Cu307.x film was primarily (001)
oriented, with small (100) peaks also present. Although the extent of in-plane
epitaxy cannot be determined from the 0-20 x-ray diffraction scans used in this
study, the high critical current density for this film (5xl07 amps/cm2 at 4.2 K) is
strong evidence that the YBa2Cu307.x grains were strongly coupled.
Comparisons between YBa2Cu307.x films deposited directly onto single-crystal
Y-Zr02 substrates with YBa2Cu3O7_x/Y-Zr02 films on Si show that the films
deposited on Si are cracked, whereas the YBa2Cu307.x films on the Y-Zr02
substrate were not cracked. SEM micrographs of Y-Zr02 films deposited on Si
do not show cracks, thus cracking occurs during YBa2Cu307.x deposition or the
cooling cycle. One approach to understand why cracking occurs in YBa2Cu307_x/
Y-Zr02 films deposited on Si is to look for microstructural differences between
these films and YBa2Cu307.x films deposited directly onto Y-ZrOz substrates.
Auger analysis shows the Y-Zr02 layer effectively prevented interdiffusion
between YBa2Cu307.x and Si, so degradation via Si interdiffusion into the
YBa2Cu307.x film was probably not the cause of cracking. X-ray diffraction
patterns from YBa2Cu307.x/Y-Zr02 films on Si versus YBa2Cu307.x films on Y-
Zr02 substrates show both are predominately (001) oriented, but the YBa2Cu307.x
film on Y-Zr02 also had minor (100) YBa2Cu307.x peaks. Based on the Auger
and x-ray data, the microstructures of the two films seem to be very similar.
Interpretation of the normal state resistance values of the two films in terms of
the Halbritter equation shows that after multiplying the resistance values of the


L ** Y2Cu205 + CuO
L + Y2Cu205 * CuO + A1203
Y2Cu205 CuO A1203
L CuO A1203
L A1203 + CuA102
(1180 C)
L ** CuA102 + CuO
(1040 C)
L ** AJ203 + CuO + CuAI02
AI203 CuO CuAI02
!s>
ON
Figure A-4. Continued.


209
oxygen readily diffuses through the yittria-rich grain boundaries, into the
YBa2Cu307.x grains, thus insuring the entire YBa2Cu307.x grain is fully oxygenated.
The combination of coherent YBa2Cu307_x/Y203 interfaces along with the high
oxygen mobility through Y203 minimizes degradation and oxygen depletion of
YBa2Cu307.x near the grain boundaries as long as the oxygen pressure is
sufficiently high during cool-down.
Dependence of T0 and Jc on A13Q3 Substrate Orientation
Variations in the electrical properties of YBa2Cu307.x/Y-Zr02 films deposited
on different orientations of A1203 substrates were also observed. YBa2Cu307.x/
Y-Zr02 films deposited on (1102) and (1210) A1203 substrates displayed metallic
resistances, but the transition to the superconducting state was sharper (83 K)
for the film deposited on (1102) A1203. The transition for the film deposited on
(1210) A1203 was broader, with a T0 of 80 K. The normal state electrical
properties for the YBa2Cu307.x/Y-Zr02 film on (0001) A1203 were much worse
than those of the films deposited on the other A1203 orientations, with resistance
values ~ 5 times higher than the YBa2Cu307.x/Y-Zr02 films deposited on (1102)
and (1210) A1203. There was also a very broad (89 77 K) transition to the
superconducting state. X-ray diffraction showed that the Y-Zr02 film deposited
on (0001) A1203 had a higher ratio of (200)/(lll) orientation than Y-Zr02 films
deposited on (1102) or (1210) A1203, and the YBa2Cu307.x/Y-Zr02 films with the
highest Jc values are typically observed when the Y-ZrOz layer is highly (100)


RESISTANCE (OHMS)
129
20
16
12
8
0
i r
YBagCu307_x/Zr02 ON
(100) SrTiOa
$
/
-6-
J L
50 100 150 200 250 300
TEMPERATURE (K)
Figure 4-24. Resistance versus temperature data for a YBajCh^O^/ZrOj film on
(100 SrTi03.


39
low temperatures will adopt the (100) orientation when deposited at higher
temperatures.15 Likewise, yittria-stabilized zirconia (Y-Zr02) films deposited on
(1102) A1203 grow with mixed (111) and (100) orientations at temperatures below
780 C, and are predominately (100) orientated when deposited at higher
temperatures.36
The relative surface free energies of the substrate, film surface, and interfacial
layer are some of the most important parameters which dictate the film
morphology, and whether film-substrate epitaxy is possible. In a thin film, the
surface energy can be a substantial portion of the total energy of the system. For
a planar interface between two phases, a and /3 37
(2-40)
adA = dUexcess- TdSexcessp
where a is the energy required to create a surface of area A, dUexcess and dSexcess
are the energy and entropy associated with creation of the new surface, and nx and
dnj are the chemical potentials and excess surface concentrations of species i at
the surface. The general expression for the surface energy is:
dU*" = dUtoua dUa -dU* ,
(2-41)
where
dUa = TdSa-PadVa PM*
(2-42)
and


131
similar film deposited on Y-Zr02. SEM micrographs of the YBa2Cu307.x/Zr02
film on (100) SrTi03 (figure 4-26) indicate the film is smooth, and the grain
boundaries are not well defined.
Comparisons between the electrical properties and microstructural features of
YBa2Cu307.x/Y-Zr02 films deposited on different orientations of A1203 helped to
clarify the types of defects which degraded superconductivity. Figure 4-27 shows
the resistance vs. temperature behavior for YBa2Cu307.x/Y-Zr02 films deposited
on the (1102), (1210), and (0001) faces of A1203. The YBa2Cu307.x/Y-Zr02 film
deposited on (1102) A1203 had the most metallic normal state resistivity and the
highest transition temperature (T0 = 83 K), while the YBa2Cu307_x/Y-Zr02 film
grown on (1210) A1203 was slightly less metallic and the T0 was 80 K.
The poorest electrical performance for YBa2Cu307.x/Y-Zr02 films deposited on
A1203 was observed on the (0001) face. The room temperature resistance values
for films on this plane were ~ 5 times as high as for YBa2Cu307.x/Y-Zr02 films
deposited on the (1102) and (1210) faces of A1203, the resistance vs. temperature
curve was less metallic, and T0 was depressed to 77 K.
Despite the varience in the orientations of the Y-Zr02 barrier layers, x-ray
diffraction data indicates the YBa^jO^ layers were highly (001) oriented for
each of the films (figures 4-28, 4-29, and 4-30). The Y-Zr02 layer deposited on
(1102) A1203 was primarily (200) oriented, with a significant (111) Y-Zr02 peak.
Although the diffraction peak from (111) is larger than the peak from (200) Y-
Zr02, the film is mostly (200) oriented because the relative x-ray intensity factor


226
Another factor which contributed to increased intragranular resistance was
incorporation of La into the YBa2Cu307.x film. For comparison, a YBa2Cu307.x
film deposited on (1102) A1203 using a YA103 barrier layer had lower resistances
and a higher transition temperature (T0 = 82 K) than did the YBa2Cu307.x films
grown on LaA103 barrier layers. Both barrier layers were amorphous, and this
indicates that La was more detrimental to the superconducting properties that
excess Y. It is difficult to determine whether the La is most damaging because it
segregates to the grain boundaries and forms insulating junctions through which
the superconducting charge carriers cannot penetrate, or if La diffuses into the
grains, replaces Y, and inhibits formation of the orthorhombic phase. Since the
Tonset for YBa2Cu307_x films deposited on (1102) A1203 using both LaA103 and
YA103 is ~ 89 K, it seems as though the temperature at which localized
superconductivity occurs within the grains is similar in both films. However, the
wide temperature region over which the resistance drops to zero in the
YBa2Cu307.x/LaA103 film on (1102) A1203 is indicative of poor intergranular
coupling. Thus we conclude that excess La is damaging to the YBa2Cu307.x film
because it segregates to the grain boundaries and inhibits superconducting charge
transport between grains. However, only (001) YBa2Cu307.x x-ray peaks were
observed in the film grown on the YA103 barrier layer, so the absence of
microcracking may also explain the difference in electrical properties.
YBa2Cu307_x/LaA103 films deposited on (1210) A1203 were even more degraded.
Resistance vs. temperature measurements show the film was semiconducting.


128
20KU XI 10 0 0 0077
Figure 4-23. Scanning electron micrograph of YBa2Cu307.x/Zr02 film deposited
on Y-Zr02.


46

where (p is the island rotation, f is the lattice misfit, and 6 is the rotation of misfit
dislocations. The maximum island rotation is obtained when 0 = 1, and hence = f. The result of island rotation is that films deposited on substrates in which
the lattice mismatch is severe will contain a much larger number of grain
boundaries than will films deposited on a more closely lattice matched substrate.
Several of the mechanisms responsible for film growth, and the energetic
parameters which dictate the orientation and microstructure of films, have been
presented in this section. Because several of the parameters are interdependent
and are difficult to observe experimentally, it is impossible to provide a
quantitative analysis of the effect of each variable. However, a general
understanding of the parameters necessary to promote two dimensional, epitaxial
growth provides considerable insight into what defects are likely to emerge from
various film processing techniques, and suggest methods by which film quality can
be improved.
Thermally induced stresses
Thermally induced cracking is a problem which plagues thin films deposited
on substrates. There is usually a difference in thermal expansion coefficients
between the film and substrate, so if the temperature of the film is changed after
the deposition, thermally induced stresses will be generated in the film and
substrate. The least desirable way for the film to relieve these stresses is by


26
of the sample is superconducting. The coexistance of normal and superconducting
regions (i.e. the mixed state) enables superconductivity to be maintained in much
higher magnetic fields in type 2 materials than are allowed in type l.4 All of the
commercially useful superconductors are type 2 materials.
Normal Superconducting
Magnetic
Flux
Density
Number of
Superelectrons
(a) Penetration depth and coherence range
Free
Energy
Density
Magnetic
Contribution
Electron-ordering
Contribution
(b) Contributions to tree energy
Free
Energy
Density
(c) Total free energy
Figure 2-8. Orgin of negative surface energy. Reference 4.


35
Nucleation and Epitaxial Growth
The superconducting properties of YBajQ^O^ are very dependent on
microstructure. A general section which describes nucleation and growth
processes is presented in order to establish a framework in which the electrical
properties of the YBa2Cu307.x films can be correlated with the growth processes.
Nucleation and growth of thin films is a large and complex subject, with many
variables which can potentially affect the microstructure and orientation of the
film. While it is difficult to obtain direct experimental evidence for many of the
parameters which determine growth mode, it is possible to speculate about the
forces which were operative during film growth by examining microstructural
properties such as surface morphology, grain orientation, and the distribution and
orientation of interfacial phases. Because superconducting and barrier layer films
have been deposited on several different substrates under a variety of growth
conditions, many of the growth mechanisms which determine film microstructure
can be deduced. In most deposition processes involving high temperature
superconductors in which the film is grown from vaporized constituents striking
the substrate, the experimental parameters which are varied include substrate
temperature, oxygen partial pressure, and the energies of the atoms and ions as
they strike the substrate. A general discussion of the fundamental processes
involved in film growth and how they are affected by growth conditions is


273
109. H. Myoren, Y. Nishiyama, N. Miyamoto, Y. Kai, Y. Yamanaka, Y. Osaka,
and F. Nishiyama, Jap. J. App. Phys. 26, L955 (1990).
110. K. Char, N. Newman, S.M. Garrison, R.W. Barton, R.C. Taber, S.S.
Laderman, and R.D. Jacowitz, Appl. Phys. Lett. 22, 409 (1990).
111. B.D. Cullity, Elements of X-ray Diffraction (2nd edition), (Addison-Wesley,
Reading, Mass., 1977).
112. J.L. Goldstein, D.E. Newbury, P. Echoin, D.C. Joy, C. Fiori, and E. Lifshin,
Scanning Electron Microscopy and X-rav Microanalysis. (Plenum Press, New
York, 1981).
113. PHI Model 660 Scanning Electron Microprobe Technical Manuel. Ver. 1.2.
(Perkin-Elmer, Eden Prarie, MN, 1987) chapter 4.
114. D.P. Woodruff and T.A. Delchar, Modem Techniques of Surface Science.
(Cambridge Press, Cambridge, 1986).
115. W.E. Beadle, J.C.C. Tsai, and R.D. Plummer. Quick Reference Manual for
Silicon Integrated Circuit Technology. (John Wiley and Sons, New York,
1985).
116. C.P. Bean, Phys. Rev. Lett. j£, 250 (1962).
117. C.P. Bean, Rev. Mod. Phys. 31 (1964).
118. Lake Shore Cryotronics, Inc., Technical Note 1/92 M2 (1992).
119. C. Thomsen, M. Cardona, B. Gegenheimer, and R. Liu, Physica C 151-155.
262 (1988).
120. R. Feile, Physica C m 1 (1989).
121. C. Thomsen, M. Cardona, B. Gegenheimer, R. Liu, and A. Simon, Phys.
Rev. B. 22, 9860 (1988).
122. U. Fano, Phys. Rev. 124, 1866 (1961).
123. R.L. Peterson and J.W. Ekin, Physica C 157. 325 (1989).
124. D.T. Shaw, Mat. Res. Soc. Bull., p. 39 (August, 1992).


YBa2Cu307_x/Y-Zr02 Films on Si, Y-Zr02, and LaA103 Substrates 108
YBa2Cu307_x/SrTi03 Films on A1203 Substrates 137
YBa2Cu307.x/LaA103 Films on Si and A1203 Substrates 154
YBa2Cu307.x/YA103 Films on Si and A1203 160
YBa2Cu307.x/Y203 Films on Si, Y-Zr02, and SrTi03 Substrates .... 173
YBa2Cu307.x/(YA103, LaA103, or Y203)/Al6Si2013 Films on
Si and A1203 180
5. DISCUSSION 196
Intergranular Versus Intragranular Effects 196
Effect of Y-Zr02, Y203, and Zr02 Barrier Layers on Jc Values .... 201
Dependence of T0 and Jc on A1203 Substrate Orientation 209
Estimate of Stress and Cracking Due to Differential
Thermal Expansion 211
Effects of Texture and In-Plane Alignment 212
Millimeter-Wave Properties 214
Effects of Surface Energy 215
Effects of Lattice Matching 217
Effects of Oxygen Pressure on SrTi03 Growth 219
LaA103 Barrier Layers 225
Y203 Barrier Layers 228
YA103 Barrier Layers 229
Al6Si2013 Barrier Layers 231
6. SUMMARY AND CONCLUSIONS 236
APPENDIX A. CALCULATION OF X-RAY ABSORPTION DEPTH
FOR YBa2Cu306^ 241
APPENDIX B. FLOW CHART OF THE Y Al Si Cu O SYSTEM . 243
REFERENCES 266
BIOGRAPHICAL SKETCH 276
v


120
4
41
20KU X5400 *0068 \.0U UFMSE
Figure 4-16. Scanning electron micrograph of YBa2Cu307x/Y-Zr02 film on
Y-Zr02.


228
have a higher thermal expansion coefficient than the (001) oriented grains. This
difference in expansion leads to increased tensile stresses and cracking.
Y3O3 Barrier Layers
YBa2Cu307.x/Y203 films deposited on (100) Si had a morphology very similar
to the YBa2Cu307_x/LaA103 films deposited on Si. These films had a very rough
surface and showed an interconnected crack network, which drastically limited the
ability of normal state or superconducting charge carriers to tunnel across the
cracks. In addition, the rough surface increased the surface area and thus
percolation distance for the hole pairs. Auger data show that Ba diffused through
the Y203 film and accumulated at the Y203/Si interface, which probably
contributed to the degradation. This behavior may have been influenced by the
high interfacial free energy of the Y-Si interface, which could have provided a
large driving force for Ba migration to and reaction with the substrate. This
implies that transport paths existed through the Y203 films. Prior studies in which
thin (< 40 ) metallic yittrium films were deposited onto Si (111) substrates by
electron beam evaporation showed that Si reacts with Y to form YSi2134. The
YSi2 grew as three dimensional islands, and large pinholes were observed after
post-annealing at 800 C. In our study, Y203 films deposited on (100) Si at 730
C then cooled to room temperature appeared to be smooth. However, we
postulate that during the subsequent superconductor deposition and extended
anneal necessary to form orthorhombic YBa2Cu307.x, the Y203 barrier layer


Cu20 Si02
2000
1600
1230
800
r
, |
Two Liquids
i '
J -1690
/
/

/
/
/
/
Si02 + Liquid
/
/
v /

\ /
v 1060
I
s*
CO
1
Cu20 + Si02
I 1
Cu20
25
50
75
1713
SiO-
Figure A-ll. Phase diagram of the CuO Si02 system. Reference 141.


APPENDIX A
CALCULATION OF X-RAY ABSORPTION DEPTH FOR YBa2Cu306J
Atomic weight of: yittrium = 88.9 g/mole
barium = 137.3 g/mole
copper = 63.5 g/mole
oxygen = 16.0 g/mole
Total weight = 88.9 + 2(137.3) + 3(63.5) + 6.5(16) = 658.2 g/mole.
Weight fraction of: yittrium = 88.9/658.2 = 0.135
barium = 274.6/658.2 = 0.417
copper = 127.0/658.2 = 0.290
oxygen = 104.0/658.2 = 0.158.
Mass absorption coefficient (u/p) of:
yittrium = 127.1
barium = 336.1
copper = 51.5
oxygen = 11.0
(u/p) for YBa2Cu307.x = 2wi(p/p)i
= 0.135(127.1) + 0.417(336.1) + 0.290(51.5) + 0.158(11.0) =
= 174.0
241


96
Baman_.SpggtrQ£CQpy
Raman spectroscopy was used to probe the intragranular microstructures of
the YBa2Cu307.x films. The intensities and positions of two of the Raman peaks
are sensitive to the oxygen content of YBa2Cu3O7.jp and were used to qualitatively
determine whether the films were oxygen deficient.119,120 Also, since each of the
Raman peaks are highly sensitive to the direction of the incident electric field, a
rough measure of the ratio of (001) vs. non-(OOl) oriented grains could be
determined. This feature was especially useful for films in which the grain sizes
were too small to produce x-ray diffraction peaks.
Raman scattering occurs when a photon is inelastically scattered by a
molecule, and the scattering is accompanied by either the absorption or emission
of a phonon with an energy equal to the difference in energies between the
incident and scattered photons:19
scattered photon ~ incident photon ^ phonon H)
When the incident light is in the visible region, the Raman spectra is dominated
by the change in polarizability as a molecule goes from its ground state to an
excited state. These polarization changes are caused by changes in the vibrational
energies of the molecule when it is excited by the electric field of the incident
light. The energies of the absorbed (or emitted) phonons are quantized, and
correspond to differences between the vibrational energy levels of the molecule.


222
ratio of the second layer is 1. We postulate that because the Sr and Ti species
reached the substrate surface as SrOx and TiOx species, incorporation of Sr and Ti
in the same plane was discouraged, so layer-by-layer growth was favored.
Evaluation of the (110) SrTi03 surface yields a different scenario. Both Sr
and Ti atoms are present in the same layer, and the fraction of oxygen atoms in
this layer is smaller than in the case of the (100) layer. In the (110) layer, both
the O/Ti and O/Sr ratios are 1, and the oxygen deficiency is alleviated by
interleaving layers composed entirely of oxygen. The structure of the (110)
surface is shown in figure 5-8.
Growth of (110) SrTi03 was favored at low oxygen pressures because the
oxygen/cation ratios were sufficiently low that a significant fraction of the Sr and
Ti species were not completely oxidized when they reached the substrate surface.
This enabled Sr and Ti to occupy the same plane, which favored growth of the
(110) orientation. Since SrTiOa is a compound with a limited range of non
stoichiometry, small deviations from the ideal SrTi03 can drastically alter the
chemical potential of the compound. The most prevalent type of defect in SrTi03
is oxygen vacancies.133 A plausible set of mechanisms which could have lead to
growth of the (110) phase were that the initial Sr-Ti-O oxygen deficient layer grew
because of the low oxygen/cation flux. Because the compound was oxygen
deficient at that point, the chemical potential for oxygen was greatly increased.
The oxygen chemical potential and oxygen deficiency was corrected for by growth


RESISTANCE (OHMS)
177
100 150 200
TEMPERATURE (K)
250 300
Figure 4-64. Resistance versus temperature data for a YBa2Cu307.x/Y203 film on
(100) SrTi03.


47
cracking, because cracks severely degrade the electrical properties of the films.
Thermally induced stresses have a significant impact on microstructural features
such as microcracking and grain size, and the microstructural processes by which
thermally induced stresses are relieved will be introduced.
Cracks are formed when the strain energy due to stresses in the film is greater
than the energy required to create new surfaces.45 The stress required to
propagate a crack is given by:
*Ef(Y,+Yp)]QS (2-55)
1 na
where Ef is the Young's modulus for the film, yE is the surface energy of the new
surface created by the crack, yp is the work of plastic deformation per unit area,
and a is the initial crack radius. A key point is that the fracture stress is
considerably lower at points where pre-existing flaws are present, such as at the
film-substrate interface.46
The model frequently used to calculate the stresses induced in thin films by
differences in the thermal expansion coefficients between the substrate and film(s)
begins with the assumption that the strain is entirely elastic.47 The dimensions of
the substrate are presumed to be unaffected by the presence of the film, and the
film is required to alter its shape in order to match the surface dimensions of the
substrate. Shear forces are generated at the film/substrate interface because of
the differing thermal expansion coefficients, and these shear forces cause the
substrate and film to bow slightly (figure 2-11). Assuming there is no plastic


INTENSITY (arbitrary units)
143
YBa2Cu307_x/SrTi03
on AI2O3(lT02)
CO
T-
o
A
A
V.
SrTi03 on AI2O3
(1T02). Deposited
at 40 mTorr 02.
A.
J V
A

smo3

AI2O3
10
CM
.O
O
O
CM
1
O'
.O
I CM
CM
4-A-
20
30
40
50
DIFFRACTION ANGLE (20)
60
Figure 4-35. X-ray diffraction patterns for SrTi03 and YBa2Cu307_x/SrTi03 films
deposited on (1102) A1203. SrTi03 films deposited at an oxygen
pressure of 40 mTorr.


RESISTANCE (OHMS)
186
120
100
80
60
40
20
0
0 50 100 150 200 250 300
TEMPERATURE (K)
T
o YBa2Cu307_x/YA10g/Al6Si2013 ON Si
YBa2CUg07_3t/LaA10g/Al6Si2013 ON Si
Figure 4-71. Resistance versus temperature data for: (O) YBa2Cu307.x/
YA103/Al6Si2013 film on (100) Si; () YBa^CW
LaA103/Al6Si2013 film on (100) Si.


79
Table 2-5. Thermal expansion of YBajCi^O**. Reference 104.
Thermal Expansion (xl06/C).
Dila-
tometer
average
<100>
<010>
<001>
Average
Orthorhombic
14.3
5.8
25.5
15.2
12.9
25-400 C
400-610 C
37.5
0.0
39.5
25.7
25
25-610 C
22.6
3.5
30.3
18.8
16.6
Tetragonal
11.5
17.0
13.3
10.9
25-800 C
SrTiOs barrier layers relative to superconducting films grown directly on A1203,
confirming that SrTi03 is an excellent barrier to A1 diffusion. Char et al.110 found
that growth of the YBa2Cu307.x/SrTi03 film on (1102) A1203 at 750 C produced
films with T0 = 86.5 K and Jc values of 2xl06 amps/cm2 at 74 K. Higher
YBa2Cu307.x deposition temperatures and hence better in-plane epitaxy, which
were made possible by the SrTi03 barrier layers, were credited as the cause for
improved electrical properties, relative to YBa2Cu307_x films deposited directly
onto (1102) A1203.


INTENSITY (arbitrary units)
Figure 4-52. Raman spectra from YBa2Cu307.x/LaAI03 films on (100) Si and (102) A1203.


122
8
T
YBa2Cu307_x/Y-Zr02
on (1102) LaAlOg
m
S
W
o
m
o
53
<
H
m
M
CQ
W
ftt
2
o
/
/
0 L
0
o
_i i_i 1 1 ; 1
50 100 150 200 250
TEMPERATURE (K)
Figure 4-17. Resistance versus temperature data for a YBa^QuO,_/Y
on (1102) LaA103.
300
Zr02 film


76
reductions in T0 and Jc as the YBa2Cu307_x film thickness increases result from
film cracking caused by tensile stresses in the YBajQ^O^ films. The cracking
seriously degrades the superconducting film properties as the YBa2Cu307.x film
thickness exceeds approximately 500 .
In addition to the stress created by the thermal expansion coefficient
differences between Si and YBa2Cu307.x, stresses in the YBa2Cu307_x films are
exacerbated because the thermal expansion coefficient of YBa2Cu307_x is highly
anisotropic.104 Table 2-5 tabulates the thermal expansion coefficients of
YBa2Cu307_x in three directions, and in different temperature regimes. The
differences in thermal expansion coefficents in the <100> and <010> directions
are caused by ordering of oxygen atoms along the <010> direction, and
formation of the orthorhombic phase. The largest thermal expansion coefficients
are observed in the <001> direction. For YBa2Cu307_x films in which the <001>
direction is normal to the substrate, nucleation of a grain with the <010>
direction along a given direction parallel to the substrate will also result in
nucleation of another grain with the <010> direction orthogonal to the first
grain. This is the method by which the system minimizes the stresses caused by
differing thermal expansion coefficients in the <100> and <010> directions. If
an (001) oriented YBa2Cu307.x film is deposited on a substrate (such as SrTi03 or
LaA103) in which the thermal expansion coefficient is close to the average
thermal expansion of YBa2Cu307_x in the <100> and <010> directions, film
cracking does not appear to be a problem. However, for substrates with lower


L ** Si02 + Al6Si2013
(1546 C)
L + AI203 ** Al6Si2013 4- Y4Si3012
A1203 Al6Si2013 Y4Si3012
L Al6Si2013 Y4Si3012
-L + Y4Si3012 ** Y2Si20
(1775 C)
L + Y4Si3012 ** Al6Si2013 + Y2Si207
Y4Si3012 Al6Si2013 Y2Si207
L Al6Si2013 Y2Si207
L ** Y2Si207 + Si02
(1660 C)
L ** AI6Si2013 + Y2Si207 + Si02
Al6^^2^13 Y2Si207 Si02
Figure A-l. Continued.


Effects of Oxygen Pressure on SrTi03 Growth
The oxygen pressure during growth of SrTi03 films on (1102) A1203 substrates
dramatically affected the orientation of the SrTi03 films. Films deposited at 200
mTorr 02 were highly (100) oriented, while films grown at 40 mTorr 02 were
(110) oriented. This behavior can be explained by correlating the surface
structures and lattice mismatches of the (100) and (110) SrTi03 films with the
deposition conditions which caused these orientations to flourish. Table 5-1 shows
that the minimum lattice mismatch occurs between the SrTiOa <110> and A1203
<012> directions. Starting with the premise that this relationship is the basis for
epitaxy between the SrTi03 film and (1102) A1203 substrate, it is equally
favorable for the film to adopt either the (100) or (110) orientations. The
principal difference between growth at the two different oxygen pressures was that
at 40 mTorr 02, the laser plume expanded from the target and encompassed the
substrate. Within the plume, the gas density is higher than the ambient, so the
flux of species from the target which reaches the substrate is greater than the
oxygen flux from the ambient.130,131,132 When the oxygen pressure was raised to
200 mTorr, the plume only extended approximately 3 cm, so it did not reach the
substrate, indicating the 02 gas density at the substrate was higher than the gas
density in the plasma. The (100) SrTi03 surface is comprised of alternating SrOx
and TiOx layers (figure 5-7). The O/Ti ratio in the first layer is 2, and the O/Sr


63
growth kinetics of the (100) and (010) oriented grains. Both of these orientations
have the <001> axis, which is the slow growth direction, parallel to the substrate,
but the <100> directions are at 90 to each other.
For YBa2Cu307.x films grown in-situ, the substrate temperature and lattice
mismatch at the film-substrate interface significantly affect the orientation of the
films.76 Films grown at 640 C on SrTi03 and LaA103 substrates were primarily
(001) oriented, whereas films on MgO, Y-ZrOz, and A1203 substrates, which were
also deposited at 640 C, grew with an (001) orientation. Increasing the growth
temperature to 720 C resulted in (001) oriented films on SrTi03 and LaA103
substrates. The variations in orientation were attributed to competition between
minimizing the surface energy of the film, which favors (001) orientation, and
minimizing structural coherence at the film-substrate interface during the early
stages of growth. Both SrTi03 and LaA103 have the perovskite structure, as does
YBajQ^Oy.jt, and the lattice matching between these substrates and YBa2Cu307.x
is reasonably close. At low growth temperatures, the reduced surface mobility
and possibility of film-substrate coherence is not as energetically favorable when
there is a large lattice mismatch between the film and substrate, which is the case
for YBa2Cu307.x films deposited on MgO, Y-Zr02, and A1203 substrates. Hence
the surface free energies are minimized by incoherent growth of the (001)
oriented grains.
Above 670 C, the tetragonal, oxygen depleted phase of YBa2Cu307_x is
stable, and at a typical growth temperatures of 750 C the non-superconducting


INTENSITY (arbitrary units)
YBa2Cu307.x /Y203/AlgSl20-j3
on (1102) Al203
250 300 350 400 450 500 550 600
RAMAN SHIFT (cm*1)
Figure 4-78. Raman spectra for YBa2Cu307_;x/YA103/Al6Si2C)13 and YBa2Cu307.J/Y203/Al6Si2013 on (1102) Al;


204
MAGNETIC FIELD (T)
Figure 5-2. Jc versus magnetic field intensity for YBa2Cu307.x/Y203 films on Y-
Zr02 and SrTi03 substrates. Data taken at 77 K. The solid line is
a fit of the experimental data to the flux creep model.


CHAPTER 6
SUMMARY AND CONCLUSIONS
YBa2Cu307_x is a promising material for low loss, low dispersion applications
such as electrical interconnects and microstrip lines on Si and A1203 substrates.
However, YBa2Cu307_x deposited directly onto Si or A1203 forms interfacial
phases which degrade both the YBa2Cu307.x film and substrate, the critical
current densities (Jc) of YBa2Cu307_x films deposited directly on (100) Y-Zr02
substrates are sensitive to the deposition conditions, and it is difficult to
reproducibly deposit films with high Jc values. To suppress interfacial
degradation, barrier layer films are deposited prior to growth of the YBa2Cu307.x
films. In this study, YBa2Cu307.x and barrier layer films were deposited in-situ
using a pulsed laser deposition process, with an excimer laser operating at a
wavelength of 248 nanometers. The barrier layers and YBa2Cu307_x films were
sequentially deposited at 730 C, without lowering the substrate temperature.
The orthorhombic, superconducting YBa2Cu307_x phase was formed during the
cooling process, and no post-annealing was required. Chemical reactions at the
substrate/barrier layer as well as at the barrier layer/YBa2Cu307_x interface
significantly influenced the normal state resistivities, the superconducting
transition temperatures, and Jc values of YBa2Cu307_x films. Optimal electrical
236


75
inductive measuring technique to intragranular conductivity, whereas the electrical
resistance technique is more sensitive to the presence of intergranular defects.
To date, Y-Zr02 has been the most successful barrier layer for YBa2Cu307.x
films grown onto Si substrates. 1 /m thick YBa2Cu307.x films with T0 = 82 K
have been grown at substrate temperatures of 650 C, using 0.1 ¡urn Y-ZrOa
barrier layers.83 The orientation of the Y-Zr02 layer greatly influences the
orientation of YBa2Cu307_x, and the best superconducting films are grown onto
(100) oriented Y-ZrOz barrier layers.103 Fork et al. showed that the ratio of
(200)/(lll) x-ray diffraction peaks from Y-ZrOz is strongly influenced by the
depostion temperature, and to a lesser extent, the oxygen partial pressure during
growth. The conditions under which highly (100) oriented Y-Zr02 barrier layers
were grown on Si were at a substrate temperature of 780 C and PQ2 = 7X10-4
Torr. Growth of (100) Y-Zr02 was also promoted by degreasing and cleaning the
silicon substrates in a flowing N2 hood, then transferring the substrates to the
deposition chamber via a N2 purged glove box, thus insuring a hydrogen
terminated Si surface. Thin YBa2Cu307.x films (305 ) deposited at 750 C, and
grown on 500 Y-Zr02 layers have T0 = 86 88 K, and Jc values of 2.2X106
amps/cm2 at 77 K. These values are comparable to those obtained from
YBa2Cu307.x films deposited on SrTi03 substrates. However, increasing the
YBa2Cu307_x film thickness to 1300 reduced the Jc to 1.5 x10s A/cm2 at 77 K.
Table 2-4 lists some of the more successful YBa^jO^/barrier layer structures
deposited on Si substrates, along with the film thicknesses. These data show that


84
experimental techniques, and the regimes in which they were used for this study,
is presented.
X-ray Diffraction
X-ray diffraction was initially used to verify that YBa2Cu307.x and barrier layer
films were being grown. Interfacial phases were also detected using x-ray
diffraction (XRD). A Phillips model APD 3720 x-ray diffractometer operating at
40 kilvolts and 20 milliamps was used to generate Cu Ka radiation of A = 1.54060
and 1.54439 . A graphite monochromater filtered out most of the Cu k/2
radiation. Peaks within the 20 range of 5 65 were detected, and the x-ray
detector was rotated at 3 per minute. Interplanar spacings were calculated using
the Bragg equation:111
nX = 2dsin0 (3-1)
where n is an integer, A is the photon wavelength, d is the interplanar spacing,
and is the angle (relative to the sample surface) at which the x-rays enter and
leave the sample. By matching the experimentally observed interplanar spacings
with those predicted by the Joint Committee of Powder Diffraction Standards
index, the phases and orientations of the films were determined. Although the
graphite monochromater eliminated over 99% of the Kfi radiation, samples which
produced extremely large Ka diffraction peaks, such as the single crystal
substrates, also produced measurable Kfi x-ray diffraction peaks.
The orientations of the YBa2Cu307_x films were highly dependent on the
orientations of the barrier layer films. To determine whether phase information


the superconducting electrons have the same momentum. When a transport
current is impressed on the system, the quantum mechanical relationship
21
m*vs=2mva=
2n
(2-21)
(where vs = velocity of superconducting electrons, and kj = k2 = kn for the
superconducting electrons) produces long range order in the momentum. There is
no resistance because the superconducting electrons are coupled together, and the
energy required to break a Cooper pair into excited electrons (quasiparticles) is
greater than twice the energy gap between superconducting and normal state
electrons (= 2A(T)).
The size of the superconducting energy gap is dependent on the fraction of
electrons which are superconducting, and this fraction varies with temperature.
Gorter and Casimer derived:
(2-22)
where n^/n is the fraction of superconducting electrons, T0 is the temperature at
which resistance vanishes in weak magnetic fields, and T is the temperature. The
variation of the superconducting energy gap (A(T)) with temperature is given by:22
(2-23)
where a is constant equal to 3.06 if weak electron-phonon coupling is assumed.
The variation of the superconducting gap with temperature is shown in figure 2-7,


199
which resulted from an inadequate supply of oxygen as the film was cooled below
the tetragonal-to-orthorhombic transition temperature (700 C). The tetragonal
regions would be normal junctions through which the superconducting charge
carriers had to pass, hence the amplitude of the superconducting wave function
was diminished and T0 lowered. The Halbritter equation was used to determine
the microstructural features most responsible for the normal state electrical
performance of the films. Since the resistance versus temperature curve does not
extrapolate to 0 ohms at 0 K for the 550 C film, this equation suggests
intergranular defects, such as grain boundaries were impeding the superconducting
charge carriers. Because the slope of the resistance versus temperature curve of
this film is larger than the 730 C film, the Halbritter equation suggests that
either intragranular defects were present, or the percolation pathway was
increased. Since the magnitude of the normal state resistance values for the 550
C film are similar to those of the 730 C film (figure 4-1), it is unlikely that the
percolation path lengths are significantly different, and we conclude that the film
was degraded by both intragranular and intergranular defects; intragranular
defects because the slope of the curve is higher than the ideal film, and
intergranular because the curve extrapolated to higher than 0 ohms at 0 K.
Scanning electron micrographs of the fully oxygenated films are virtually
featureless, except for the particulates which resulted from the laser deposition
process. This reinforces the conclusion that increased percolation distance was
not the cause of increased resistance in the oxygen depleted film. X-ray


28
Ch ^(Aft) | h 3(A 4ite dt2 dt
(2-28)
where C is the capacitance, h is Planck's constant, and Rn is the normal state
resistance of the junction. Ambegaokatar and Baratoff derived an expression
which shows that Jc is inversely proportional to the normal state resistance,22
(2-29)
where A(T) is the superconducting energy gap.
Coupling between the superconductor wavefunctions on each side of a
junction, and thus tunneling current density, are rapidly diminished by increasingly
thick insulator regions:23
(2-30)
where s is the insulator thickness (nm), a = 0.1 (eVy^/nm is the effective barrier
height, and the electron energy gap between valence and conduction bands in
the insulator. Many devices, such as detectors and oscillators, could be made
using superconductors if SIS junctions were more reproducible. However, since
superconducting areas within a distance smaller than the coherence length, £, of
the junction are often degraded, A(T) is not the energy gap for a perfect material
but is depressed because of defects near the junction. Thus Jc for a given junction
is difficult to control.
When two superconductors are separated by a normal metal (SNS structure),
superconductivity is induced in the normal metal via the proximity effect. The


16
dimensions for a VLSI microstrip line (thickness = 0.1 0.3 pm, width =1-3
/im, and length = 1-10 /im) dispersion necessitates that pulse lengths longer than
100 picoseconds must be used. However, a superconducting line of the same
dimensions would have negligible dispersion for pulses longer than 1
picosecond.15,16
Figure 2-5. Calculated phase velocities as a function of frequency for
superconducting and aluminum microstrip lines, at 77 K. Reference
15.
The real part of the propagation constant, a, is the attenuation and results
from Joule losses within the microstrip line. In a superconducting material, the
penetration depth is constant, so a (decibel/cm) is linearly proportional to the
number of wavelengths/cm, and thus the frequency. In normal metals, a is
dependent on the skin depth, which is proportional to or0-5, so a increases more
slowly with increasing frequencies in metallic microstrip lines than it does in
superconducting lines. However, at microwave and millimeter wave frequencies,


176
m > *
A
*.4|
.\ ^
4 *
20KU &600 5300 1.0UUFMSE
Figure 4-63. Scanning electron micrograph of a YBajQ^O? JY203 film on
Y-Zr02.


214
YBa2Cu307.x films deposited on LaA103 or SrTi03 substrates are featureless at
similar magnifications. This morphology may have caused the reduction in T0 and
Jc values for the YBa2Cu307.x/SrTi03 film on (1102) A1203.
Millimeter-Wave Properties
The millimeter-wave properties of YBa2Cu307.x/SrTi03 films were very good
(figures 4-38 and 4-39). The resistance at 36 GHz was 10'2 ohms, which was lower
than the value observed in a YBa2Cu307_x films deposited on LaA103 (5 X 10~2
ohms). There was a sharp change in the phase angle of the film as it went from
the normal to superconducting state. In the normal state, the phase angle was
close to 0, which is typical for normal conductors. At the transition temperature,
the phase angle of the current sharply dropped to 90 out of phase with the
voltage, which is expected for films in which the superconducting charge carriers
dominate the conductivity. There was no evidence that the SrTi03 barrier layer
affected the dielectric properties of the film or substrate, despite the fact that bulk
SrTi03 is lossy and has a dielectric constant of over 1500 at 77 K. Apparently,
film stresses inhibit the tetragonal-to-orthorhombic phase transition in SrTi03
films. Since the high dielectric constant is caused by the structural phase
transformation, inhibiting this transition prevents the dielectric constant from
increasing. Recent experiments in which the dielectric constant of a 4000
SrTi03 film was measured as a function of temperature show that the dielectric
constant is 300 from 4.2 to 300 K.129


Figure 4-69. Auger line scan across crater edge of a YBa2Cu307.x/Y203 film on (100) Si.


Figure 4-56. Auger line scan across crater edge of a YA103 film on (100) Si.


274
125. Y. Iijima, N. Tanabe, O. Kohno, and Y. Ikeno, Appl. Phys. Lett. j&Q, 769
(1992).
126. A. Catana, R.F. Broom, J.G. Bednorz, J. Mannhart, and D.G. Schlom, Appl.
Phys. Lett.m 1016 (1992).
127. Q.Y. Ying, C. Hilbert, N. Kumar, D. Eichman, M. Thomson, H.Kroger, and
D.M. Hwang, Appl. Phys. Lett. 5& 3036 (1991).
128. J.E. Blendell, D.K. Chiang, D.C. Cranmer, S.W. Freiman, E.R. Fuller, Jr., E.
Descher-Krasicka, W.L. Johnson, H.M. Ledbetter, L.H. Bennett, LJ.
Swartzendruber, R.B. Marinenko, R.L. Myklebust, D.S. Bright, and D.E.
Newbury, Adv. Ceram. Mat. 2, 512 (1987).
129. A. Walkenhorst, C. Doughty, X.X. Xi, S.N. Mao, Q. Li, T. Venkatesan, and
R. Ramesh, Appl. Phys. Lett. 5Q, 1744 (1992).
130. D.B. Geohegan, Appl. Phys. Lett. 5Q, 2732 (1992).
131. A. Gupta, B. Braren, K.G. Casey, B.W. Hussey, and R. Kelly, Appl. Phys.
Lett. 52, 1302 (1991).
132. P.E. Dyer, A. Issa, and P.H. Key, Appl. Phys. Lett. 52, 186 (1990).
133. V.E. Henrich, in Surface and Near-Surface Chemistry of Oxide Materials,
ed. J. Nowotny and L.C. Dufour (Elsevier, Amsterdam, 1988) p. 40.
134. M.P. Siegal, JJ. Santiago, and W.R. Graham, Mat. Res. Soc. Symp. Proc.
Vol. 198, p. 589 (1990).
135. M.J. Hyatt and D.E. Day, J. Am. Ceram. Soc. JQ, C-283 (1987).
136. N.A. Toropov, I.A. Bondar. F.Y. Galakhov, X.S. Nikogosyan, and N.V.
Vinogradova, in Phase Diagrams for Ceramists. 1969 Supplement, edited by
E.M. Levin, C.R. Robbins, and H.F. McMurdie (American Ceramic Society,
Columbus, 1969) p. 96.
137. N.A. Toropov and LA. Bondar, in Phase Diagrams for Ceramists. 1969
Supplement, p. 107.
138. A. Staronka, H. Pham, and M. Rolin in Phase Diagrams for Ceramists. 1975
Supplement, edited by E.M. Levin and H.F. McMurdie (American Ceramic
Society, Columbus, 1975) p. 133.


99
non-(001) orientations. Correlations between the height of the 440 cm"1 peak and
oxygen content have not been established, so this peak was only used to identify
the presence of non-(OOl) oriented grains.
Millimeter-Wave Transmission Measurements
The amplitude and phase angle of millimeter-wave signals transmitted through
a YBa2Cu307.x/SrTi03 film deposited on (1102) A1203 were measured using the
millimeter-wave power transmission technique. From this data, surface resistance
and complex conductivity for the YI^Q^O^ film were calculated.
A 36 gigahertz signal, operating in the transverse electric (01) mode, was
propagated in a rectangular waveguide, at a power level of 16 milliwatts. The
substrate was clamped to the end of the waveguide with two waveguide flanges,
with the film side facing the incoming signal. The magnitude and phase of the
transmitted power were measured as a function of temperature with a Hewlett-
Packard 8510 Network Analyzer. Changes in the fraction of power transmitted
through the film/substrate combination, and the phase of the transmitted power,
were dominated by changes in the resistance of the YBa2Cu307_x film as it was
cooled from the normal to the superconductive states. A complete description of
the formulas used to calculate the transmitted power, phase angle of the signal,
surface resistance, and complex conductivity are presented in reference 12.


INTENSITY (arbitrary units)
Figure 4-75. Raman spectra of YBa2Cu307.x/YA103/A]6Si2013 and YBa2Cu307.x/LaA103/Al6S2013 films on (100) Si.


268
32. J. Venables and G.L. Price, in Epitaxial Growth. Part B, edited by J.W.
Matthews (Academic Press, New York, 1975) p. 387.
33. D. Walton, Philos. Trans., 167 (1962).
34. R. Kern, in Current Topics in Materials Science. Vol. 3. edited by E. Kaldis
(North-Holland, Amsterdam, 1979) p. 319.
35. K.I. Chopra, Thin Film Phenomena. (Krieger Publishing, Malabar, 1985).
36. H. Schmidt, K. Hradil, W. Hosier, W. Wersing, G. Gieres, and RJ. Seebock,
App. Phys. Lett. 2, 222 (1991).
37. R.A. Swalin, Thermodynamics of Solids. 2nd ed. (John Wiley and Sons, New
York, 1972).
38. L.E. Murr, Interfacial Phenomena in Metals and Alloys. (Addison-Wesley,
London, 1975).
39. R. Kern in Current Tonics in Materials Science. Vol. 3. edited by E. Kaldis
(North-Holland, Amsterdam, 1979) p. 187.
40. J.H. van der Merwe in Recent Developments in the Theor of Epitaxy, edited
by R. Vanselow and R. Howe (Springer-Verlag, Berlin, 1984) p. 365.
41. A. Zur and T.C. McGill, J. Appl. Phys. 5, 378 (1984).
42. D.A. Porter and K.E. Easterling, Phase Transformations in Metals and
Alloys. (Van Nostrand, New York, 1981) p. 148.
43. C.A.B. Ball and J.H. van der Merwe, in Dislocations in Solids, edited by
F.R.N. Nabarro (North-Holland, Amsterdam, 1983) p. 123.
44. J.W. Matthews, in Epitaxial Growth. Part B. edited by J.W. Matthews
(Academic Press, New York, 1975)
45. W.D. Kingery, H.K. Bowen, and D.R. Uhlmann, Introduction to Ceramics.
2nd ed. (John Wiley and Sons, New York, 1976) p. 786.
46. M.F. Doerner and W.D. Nix, CRC Crit. Rev. in Sol. State and Mat. Sci. 14,
225 (1988).
47. P.H. Townsend, D.M Barnett, and T.A. Brunner, J. Appl. Phys. 2,4438
(1987).


generate electromagnetic radiation at milliwatt-level powers,8 at frequencies up to
1012 Hz.
Figure 2-1. The current versus voltage characteristic of a Josephson junction.
Reference 7.
Radiation detectors are based on the principle that electronic radiation will
induce AC currents in the Josephson junction.7 Regardless of the orgin, currents
add algebraically, hence the DC critical current will be suppressed by the
radiation induced AC currents. In the finite voltage regime, the mixture of AC


RESISTANCE (OHMS)
150
50
40
30
20
10
0
i i i i i
o YBa2Cua07_x/SrTi0g on (1210) Al20a
YBa2Cu307_x/SrTi0g on (0001) A1203
0 50 100 150 200 250 300
TEMPERATURE (K)
Figure 4-41. Resistance versus temperature data for YBa2Cu307.x/SrTi03 films
on: () (0001) A1203; (O) (1210) A1203.


INTENSITY (ARBITRARY UNITS)
191
Figure 4-74. X-ray diffraction patterns from YBa2Qi307_x/YA103/Al6Si2013 and
YBa2Cu307.x/LaA103/Al6Si2013 films on (100) Si.


RESISTANCE (OHMS)
133
YBa2Cu307_x/Y-Zr02 FILMS DEPOSITED
ON DIFFERENT ORIENTATIONS OF A10,
A 3
Figure 4-27. Resistance versus temperature data for YBa2Cu307.x/Y_-Zr02 films
on different orientations of A1203: (O) (1102); () (1210); (A)
(0001).


165
YBajCi^Oy.x/YAlC^ films deposited on Si and (1102) A1203, the resistance was
normalized by multiplying the resistance values of the YBa2Cu307.x/YA103 film
on Si by 0.125. The resistance versus temperature curve of this film was slightly
less metallic than for YBa2Cu307.x /YA103 on (1102) A1203, and had a lower
transition temperature (T0 = 78 K). In this study, the YBa2Cu307.x/YA103 film
on Si had the highest T0 of all the films deposited on Si substrates. SEM data
(figure 4-54) show that the YBa2Cu307_x/YA103 film on Si is composed of
interconnected submicron grains, with much fewer cracks as compared to samples
of YBa2Cu307.x/Y-Zr02 on Si. A cross-sectional view (figure 4-55) shows the
film/substrate interface is sharp, but there is not a well defined interface between
the YBa2Cu307.x and YA103 layers. AES data of YA103 deposited onto Si (figure
4-56) show segregation of A1 towards Si and formation of an interfacial Al-Si-O
phase. X-ray diffraction (figure 4-57), however, does not indicate any Al-Si-O
phase. AES analysis of the YBa2Cu307.x/YA103 film on Si (figure 4-58) show
very little interdiffusion between YBa2Cu307.x, YA103, or Si, indicating YA103 was
an effective barrier to chemical diffusion. SEM of YA103 films deposited on Si at
730 C (figure 4-59) show the films are featureless, hence the morphology
observed in the YBa2Cu307.x/YA103 film must have emerged during or after the
YBa2Cu307.x deposition.
YA103 was also an effective diffusion barrier between YBa2Cu307.x and
(1102) A1203. The only x-ray diffraction peaks from YA103 deposited on (1102)
A1203 are from the substrate (figure 4-60), indicating the film was amorphous.


221
Sr-Ti-0
LAYER
0.39 nm
OXYGEN
LAYER
0.39 nm
SURFACE STRUCTURE OF (110) SrTi03
O = Oxygen
= Titanium
= Strontium
ooooo
8*8888

> iO oo o o
0-550 nnn If
*o oooo
k >\
oo oo oo oo
oooooooo
^ lOO oo oo oo
0.550 nm
*0 0 oo oo oo
K H
Figure 5-8. Surface structure of (110) SrTi03.


8
CQ
3

o
w
u
s
CO
II
CQ
w

6
4
2
0
i i
O YBaaCus07_x ON (10Z) LaA10a. CHAMBER
FILLED WITH 0a AFTER DEPOSITION.
YBagCua07_x ON (ll02) LaA10#. FILM
COOLED TO 660 C BEFORE FILLING . '
CHAMBER WITH 0-. . '
* /
0 50 100 150 200 250 300
TEMPERATURE (K)
Figure 4-1. Resistance versus temperature data for YBajQ^O^ films deposited
on (102) LaA103. (O) Oxygen chamber filled with oxygen
immediately after the deposition, while the substrate temperature was
730 C. () Film cooled to 550 C before filling the chamber with
oxygen, and the lower T0 (79 K) was attributed to an incomplete
tetragonal-to-orthorhombic phase transition.


RESISTANCE (OHMS)
174
Figure 4-61. Resistance versus temperature for a YBajQ^O^/YjC^ film
on Y-Zr02.


Figure 4-70. Scanning electron micrograph of a Y203 film on (100) Si.


82
vacuum or reducing the substrate temperature. This holder did not allow
continuous rotation of the target during deposition, but the target was moved
slightly every 800 pulses to expose a new surface to the laser radiation. The
deposition temperatures were measured by a thermocouple spot-welded to the
heater block. The barrier layer films were approximately 1000 thick, and unless
specified otherwise, were grown in 40 mTorr 02. The YBa2Cu307_x films were
2000 3000 thick, and were always deposited at PD2 = 200 mTorr. After the
YBa2Cu307.x films were deposited, the chamber was filled with oxygen, and the
temperature was maintained at 730 750 C for 20 minutes in order to facilitate
the tetragonal-to-orthorhombic transition. The films were cooled to 450 C over
60 minutes, held at 450 C for 45 minutes at approximately 300 Torr of oxygen to
ensure complete oxygenation, then cooled to room temperature.
All of the substrates used in this study were single crystal. The SrTi03 and Si
substrates were cut parallel to the (100) planes, LaA103 was cut parallel to the
(1102) planes, and A1203 was cut parallel to the (1102), (1210), or (0001) planes.
The Y-Zr02 substrates were cut 5 12 degrees from the (100) planes, as
determined by Laue back-diffraction patterns. By depositing the YBa2Cu307.x
films on off-axis Y-Zr02 substrates, we increased the tendency for high-angle
grain boundaries to form in the barrier layer and superconducting films. This
feature enabled us to determine whether the barrier layers would passivate the
YBa2Cu307.x films from the defects introduced by these substrates, and allow
growth of superconducting films with high Jc values.


Table 5-1. Continued.
Substrate
Barrier Layer
T0 (K)
Jc (amps/cm2)
(100) Si
LaA103
< 46

(1102) A1203
LaA103
58

(1210) A1203
LaA103
< 30

(100) Si
yaio3
78

(1102) A1203
yaio3
82
lxlO4
off-axis Y-Zr02
y2o3
90
lxlO7
(100) SrTi03
y2o3
88
1.8 X107
(100) Si
y2o3
< 20

(100) Si
Y A103/Al6Si2013
61.5

(100) Si
LaA103/ Al6Si2013
59

(1102) A1203
YA103/Al6Si2013
75

(1102) A1203
Y203/AlfjSi2013
< 39



INTENSITY (ARBITRARY UNITS)
Figure 4-43. X-ray diffraction patterns from SrTi03 and YBa2Cu307_x/SrTi0
films on (1210) A1203.


215
Effects of Surface Energy
The good superconducting properties of YBa2Cu307.x/SrTi03 films on (1102)
A1203 were largely due to the ease with which (100) oriented SrTi03 films could
be deposited on (1102) A1203 substrates. Comparisons between the orientations
of SrTi03 films deposited on (1102), (0001), and (1210) A1203 indicates that
lattice mismatch, excess AHf for SrTi03, and phenomena intrinsic to the pulsed
laser deposition process synergistically affected the microstructures of the SrTiOa
films.
Reduction of interfacial free energy by epitaxial growth is a strong driving
force for highly textured film growth. Table 5-1 lists the smallest in-plane lattice
mismatches which can occur between the low index directions of SrTi03, LaA103,
(100) MgO, and (1102), (1210), or (0001) orientations of A1203. The minimum
lattice mismatch between SrTi03 and any cut of A1203 is 7.4%. By comparison,
the minimum lattice mismatch between LaA103 and A1203 is only 4.5%.
Experimentally, we observe that highly oriented SrTi03 can be grown on (1102)
and (0001) A1203, whereas LaA103 films deposited on (1102) A1203 are
amorphous. The cause of the behavior was attributed to the more negative excess
AHf for SrTi03. For the reaction


74
excellent quality YBa2Cu307_x films have been deposited on single crystal LaA103
substrates, LaA103 has not been reported to be a good barrier layer material for
growth of YBa2Cu307_x on either A1203 or Si substrates. Attempts to grow
LaA103 films on a variety of substrates at 760 C showed that epitaxial films
could be deposited on SrTi03 and LaA103 substrates, while attempts to grow
LaA103 films on Si, A1203, and MgO substrates resulted in amorphous films.108
This indicates that matching both the lattice parameters and crystal structures at
the film-substrate interface are important parameters which critically influence the
crystallinity and orientation of the YBa2Cu307.x films.
Experiments in which LaA103 and YBa2Cu307.x were simultaneously deposited
onto an MgO substrate showed that T0 gradually dropped as the fraction of
LaA103 in the film increased.101 YBa2Cu307.x films which contained no LaA103
had a T0 = 87 K, while the transition temperature dropped to 77.5 and 30 K
for YBa2Cu307.x films containing 9 and 13 mole percent LaA103, respectively.
When the transition from Tonset to T0 was determined by measuring the magnetic
response of the film to a magnetic field (inductive response), the transition
remained fairly sharp (< 5 K) for all the films. However, when measured by the
electrical resistance technique, the transitions became much broader as the
fraction of LaA103 in the films increased. Apparently, the formation of LaA103
at the grain boundaries decreased coupling of the superconducting wave function
between the grains. The discrepancies in the widths of the temperature ranges
over which the transitions occured was attributed to the higher sensitivity of the


CRITICAL CURRENT DENSITY (A/cm2)
Figure 4-18. Critical current density versus magnetic field intensity for a YBa2Cu307.x/Y-Zr02 film deposited on
(102) LaA103.
to
OJ


INTENSITY (ARBITRARY UNITS)
136
Figure 4-30. X-ray diffraction patterns from Y-Zr02, and YBa2Cu307.x/Y-Zr02
films on (1210) A1203.


31
of 5 meV was assumed.25 However, the Ambegaokatar-Baratoff expression is also
valid for dirty SNS junctions, provided that significant changes in the phase of the
wavelength only occur near or in the normal layer; hence the observed gap could
represent the reduced order parameter inside the normal layer. The authors
concluded that grain boundaries were dirty SNS junctions because Jc(0) was a
factor of 10 less than values predicted by the SIS model. On the other hand, Jc vs
B(T) for different orientations of the magnetic field indicated intragranular Jc
values were limited by flux creep across the grains, and intragranular Jc values
were determined by the density and pinning energies of the flux-pinning sites.
Because the sensitivity of electrical properties in the superconducting and
normal states to film microstructure is very pronounced, a great deal of
information about the film microstructure can be obtained from the resistance vs.
temperature data. Zero resistance occurs when there is a continuous pathway in
which the wave function of the superconductor is phase ordered along the entire
path. If the wave function is strongly coupled between grains, the transition from
the normal to the superconducting state is abrupt, and the film T0 is close to the
transition temperature observed in bulk samples. However, if the film
superconductivity is localized, meaning there are isolated superconductor volumes
such as grains separated by barriers at which the intergranular coupling is weak,
or if the coupling strength varies randomly at the different grain boundaries, the
transition to the superconducting state will be percolative and there will be a
finite resistance at temperatures below Tonset. Deviations from the ideal resistivity


FLOW CHART FOR Al Cu Si O
A1,Q, SiO, CuO ALO, TERNARY CuO SiQ.
v
Figure A-2. Flow chart of the Al Cu Si O system.


208
Despite similar resistance versus temperature behavior for all the films, there
were large variations in Jc. The T0 for each of the films was 87-90 K, and the
resistance versus temperature curves extrapolated close to 0 ohms at 0 K. This
shows that the normal state intergranular resistances were neglible compared to
the intragranular resistances. However, superconducting films deposited on Y203
or Y-Zr02 barrier layers had high Jc values, but the Jc values of films on Zr02
barrier layers were lower, and were limited by poor intergranular coupling of the
superconducting hole pairs. SEM micrographs show that films deposited with
Y203, Y-Zr02, or Zr02 barrier layers had different grain sizes and surface
roughnesses, but there was no correlation between the features observed by SEM
and the Jc values. We believe the YBa2Cu307.x films with Y203 or Y013Zr0 87O194
barrier layers grew with yittria-rich grain boundaries, which were flux pinning sites
at the grain boundaries, similar to the suggestion of Catana et al.126
Effective flux-pinning sites result from steep gradients in the superconducting
order parameter between superconducting and non-superconducting regions of the
film. The ability of the YBa2Cu307_x/Y203 interface to be a flux-pinning site can
be traced to two properties. First, Catana et al.126 showed that the YBa2Cu307.x
/Y203 interface was coherent, and the intersection of a growing YBa2Cu307.x
grain with a Y203 inclusion resulted in epitaxial growth of YBa2Cu307.x around
the inclusions, with formation of edge dislocations and other defects near the
Y203 inclusions. Second, Ying et al.127 reported that the mobility of oxygen
through Y203 is very high. This means that as the YBajQ^O^ films are cooled,


FLOW CHART FOR Y Cu A1 0
Y,Q, CuO Y,Oo ALO, TERNARY CuO ALO,
L ~ Y203 + Y4A1209
(1940 C)
L + Y4A1209 YA103
(1875 C)
L YA103 + Y3A15012
(1850 C)
L ~ Y3A15012 + A1203
(1760 C)
L + Y203 * Y2Cu2Os
L + Y203 ** Y2Cu205 + Y4A1209
Y203 Y2Cu205 Y4A1209
L Y2Cu205 Y4A1209
Figure A-4. Flow chart of the Y Cu Al O system.


216
SrO + Ti02 = SrTi03
the system lowers its AHf ^ .c by 47.4 kilojoules/(gram atom) of oxygen by
adopting the SrTi03 structure, compared to the extrapolated AHf m .c which a
mixture of SrO + TiOz phases would have. Although the AHf for LaA103 has not
been measured, it is likely that the excess AHf is much smaller than that of
SrTi03, since the crystal structures of La^ and A1203 are both hexagonal, and
the structure of LaA103 is basically an alternating series of LaOx and A10x layers.
Hence we conclude that the excess AHf for SrTi03 was a prime driving force for
crystallization of the film, whereas the lower excess AHf for LaA103 caused the
film to remain amorphous.
Table 5-2. Lattice mismatch between SrTi03 or LaA103 films, and A1203 or MgO
substrates. Mismatches were calculated by using the value of ardm for
the direction indicated in the vertical column, and calculating the
minimum mismatch which could be achieved by matching afllm with
one of the low index directions of the substrate surface, and using the
lattice parameter of this direction for a,.ubstrate.
(1102) A1203
(1210) A1203
(0001) ai2o3
(100) MgO
<100> SrTi03
+ 11.5
-5.3
-18.1
-7.1
<110> SrTi03
+ 7.4
+15.1
+ 15.8
+ 31.2
<100> LaA103
+ 8.3
-8.0
-20.4
-9.8
<110> LaA103
+ 4.5
-17.5
+ 12.6
+ 27.6


Table 2-4. Superconducting transition temperatures and critical current densities for YBa2Cu307.x/barrier layer films
on silicon and LaA103 substrates.
Substrate
Barrier layer and
Thickness ()
YBa2Cu30,x
thickness (A)
Transition
temperature
(K)
Critical current
density
(amps/cm2)
Reference
Si
Y-Zr02 (1000)
10,000
82

83
Si
BaT i03/MgAl204
(3500/5000)

70

107
Si
BaTi03/MgAl204
(3500/750)
1000
86-87
6xl04 at 77 K
108
Si
Yp/^-ZrOj
(100/900)
600
82-84
lxlO6 at 77 K
109
Si
Y-Zr02
(500)
130
86-88
2xl06 at 77 K
103
Si
Y-Zr02
(500)
1350

1x10s at 77 K
103
LaA103

1300
88-90
5xl06 at 77K
73
o
00


85
about the barrier layer or interfacial phases buried beneath the superconducting
film could be obtained, the x-ray attenuation depths for the various films were
calculated. Assuming the x-ray intensity decreases exponentially as it enters the
sample, the attenuation for each compound was calculated using the expression:
-f =exp[-(-£)p] (3-2)
h P
where I/I0 is the fraction of the incident x-ray intensity which penetrates to a
depth = t, (p/p) is the mass absorption coefficient for each phase, and p is the
density of the phase, (p/p) was calculated from the weighted fractions of the
mass absorption coefficients for the individual elements:
(tW = wi(-) +w2(t) + w3() +-+w(t)
(3-3)
P Pi P 2 P 3 P
where wn is the weight fraction of element n in the phase, and (p/p)n is the mass
absorption coefficient for element n. For the calculation, the penetration depth at
which I/Iq was equal to 0.368 was taken to be the absorption depth (= t). Of the
films examined in this study, YI^CujO** had the smallest absorption depth (=9
pm). Since the YBa2Cu307.x films were typically 3000 thick, and the barrier
layer films were 1000 thick, we concluded that diffraction data was obtained
from all of the films in the multilayered structure. The calculation of the x-ray
absorption depth for YBajQ^O^ is presented in appendix A.


218
SURFACE MESHES OF (0001), (1102) AND (1Z10)
ORIENTATIONS OF Al203
O = Oxygen
0 = Aluminum
(1210) A
1.299 nm
>
0.824 nm
*0.0 0.0-0 o
0.0 0.0.0
*0 o O- o om o
o o o o' o
Figure 5-6. Surface meshes of (0001), (1102), and (1210) orientations of A1203.


INTENSITY (arbitrary units)
Ln
Figure 4-12. Raman spectra from YBa2Cu307.)/Y-Zr02 and YBa2Cu307./YA103 films deposited on (100) Si.


182


18
Microstructure/Superconductor Relations
The microstracture of a superconducting film has a large influence on the
electrical and magnetic properties, and understanding how the superconducting
properties are impacted by film microstructure will dictate the processing
procedures used to fabricate the films. In YBajQ^O^ the superconducting and
normal-state properties are anisotropic with respect to crystallographic direction,18
so film orientation is a critical parameter. The T0 and Jc of the film are
dependent on superconducting properties such as coherence length (£), and
penetration depth (A), which, in turn, are influenced by orientation, grain size,
and grain boundaries. Clearly, microstructural considerations play a large role in
defining the electrical and magnetic properties of the film. To understand the
relationship between microstructure and superconductive behavior, it is instructive
to review how superconductivity evolves. Although the physical processes
responsible for superconductivity in YBa2Cu307_x have not been uncovered, most
of the properties are well described by the microscopic Bardeen-Cooper-Schrieffer
(BCS) theory and the phenomenological equations which predate BCS, so the
properties of YBa2Cu307.x will be described in terms of conventional
superconductors.
In a normal metal, there is a repulsive energy between electrons, and the
energy levels are described by Fermi-Dirac statistics:19,20


224
on (110) SrTi03 barrier layers did not become superconducting until 77 K. This
discrepancy was attributed to microcracking in the YBa2Cu307_x film deposited on
(1102) A1203. Since the thermal expansion coefficient of YBa2Cu307.x is largest in
the <001> direction, a (103) oriented YBa2Cu307_x film will be subjected to
larger thermal stresses than an (001) oriented YBa2Cu307_x film, so microcracking
is more probable.
SrTi03 films deposited on (0001) A1203 at 200 mTorr were strongly (111)
oriented, and YBa2Cu307.x films deposited on this structure were primarily (001)
oriented, with minor fractions of (100), (013), and (113) phases also present. The
non-metallic normal state resistance was attributed to high-angle grain boundaries
resulting from the polycrystalline nature of the film, and the reduced T0 was
caused by thermally induced microcracks, similar to the case in which the
YBa2Cu307.x/Y-Zr02 film was deposited on (0001) A1203.
(1210) A1203 was the only plane of A1203 on which highly oriented SrTi03
barrier layers were not grown. X-ray diffraction shows only weak (110) SrTi03
peaks from films deposited at 200 mTorr onto (1210) A1203, and the principal
peak from a YBa2Cu307.x film deposited on this structure is the (103) orientation.
The resistance vs. temperature behavior of this film was semiconducting, with a T0
= 52 K. The poor electrical performance of this film resulted from several
factors, including high angle grain boundaries, microcracking, and judging from
the weak intensities of the x-ray peaks, a large fraction of amorphous Y-Ba-Cu-O
constituents.


RESISTANCE (OHMS)
112
Figure 4-9. Resistance versus temperature data for a YBa2Cu307.x film deposited
on a Y-Zr02 substrate, and a YBa2Cu307_x/Y-Zr02 film on (100) Si.
Resistance values for the YBa2Cu307.x/Y-Zr02 film on Si are
multiplied by 0.061.


110
20KU X4000 3080 1.0U UFMSE
Figure 4-8. Scanning electron micrograph of a Y-Zr02 film on (100) Si.
Raman spectra of the YBa2Cu307.x/Y-Zr02 film deposited on Si (figure 4-12)
show an asymmetric peak at 335 cm'1, and a smaller peak at 500 cm'1. The
spectra from a YBa2Cu307_x/YA103 film on Si is also shown in this graph, and the
similarities between these two spectra and the spectra obtained from YBa2Cu307_x
films deposited on (1102) LaA103 indicates that the intragranular regions of the
YBa2Cu307.x/(Y-Zr02 or YA103) films deposited on Si were undamaged.
Deposition of Y-Zr02 barriers on single-crystal Y-ZrOz substrates prior to the
YBa2Cu307_x films significantly improved the normal state electrical behavior and
Jc values, as shown by the resistance versus temperature behavior in figure 4-13.


INTENSITY (ARBITRARY UNITS)
Figure 4-31. Raman spectra of YBa2Cu307.x/Y-Zr02 films on (102) and (0001) orientations of A1203.
w
00


267
17. K.B. Bhasin, J.D. Warner, R.R. Romanofsky, V.O. Heinen, and CM.
Chorey, NASA Tech. Mem. 103235.
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Physics-Volume 2 (Addison-Wesley, London, 1964).
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Ql, 219 (1988).
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Rev. Lett. 01, 2476 (1988).
26. J. Halbritter, Int. J. Mod. Phys. 2, 719 (1989).
27. C.B. Eom, A.F. Marshall, S.S. Laderman, R.D. Jacowitz, and T.H. Geballe,
Science 249. 1549 (1990).
28. RJ. Cava, BJBatlogg, C.H. Chen, E.A. Rietman, S.M. Zahurak, and D.
Werder, Phys. Rev. B. 20, 5719 (1987).
29. C.J. Jou and J. Washburn, J. Mater. Res. 4, 795 (1989).
30. S.W. Chan, D.M. Hwang, R.Ramesh, S.M. Sampere, L. Nazar, R. Gerhardt,
and P. Pruna, in High Tc Superconducting Thin Films: Processing.
Characterization, and Applications, edited by R. Stockbaur, (AIP Proc. 200.
New York, 1989) p. 172.
31. G.M. Rosenblatt, in Treatise on Solid State Chemistry. Yol. 6A. edited by
N.B. Hannay (Plenum Press, New York, 1976) p. 179.


269
48. W.D. Nix, Met. Trans A 2Q, 2217 (1989).
49. J. Vilms and D. Kerps, J. Appl. Phys. 2, 1536 (1982).
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1098 (1988).
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A.W. Kleinsasser, R J. Gambino, B. Bumble, and M.F. Chisholm, Appl.
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Appl. Phys. Lett. i4, 1365 (1989).
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63. J.H. Brannon, IEEE Circuits and Devices, 19 (Sept., 1990).


49
SIR
Tensile
ESS
Compressive
Figure 2-11. Cross-sectional side view of a two-layered film on a substate,
showing how elastic film stresses are accomodated by the substrate.
Reference 47.
between the zero strain plane and the film-substrate interface. The second
requirement is that the sum of moments about the axis which runs through
the zero strain plane must equal zero. Continuum mechanics requires that the
sum of moments induced by various mechanical forces acting on a common axis
must cancel each other in order for the axis to remain stationary. In our case, the
axis running through the zero-strain plane is common to all the forces acting on
the films and substrate.


Figure 4-57. X-ray diffraction pattern of a YBa2Cu307_x/YA103 film on (100) Si.
INTENSITY (ARBITRARY UNITS)


66
YBa2Cu3C>7
Chain
Plane
o-3,
Figure 2-16. Crystal structure of YBajCujO^ illustrating the CuO chains and
Cu02 planes.


180
Comparing Jc data from YBa2Cu307x films deposited on Y-Zr02 substrates using
Zr02, Y-Zr02, or Y203 barrier layers, we conclude that Y203 concentrations
greater than 13% in the barrier layers are crucial to the formation of YB^CujO^
films with high Jc values.
YBa2Cu307.x/Y203 films on Si had very poor electrical properties, with a
semiconducting normal state resistance and a long transition from Tonset = 90 K
to T0 less than 20 K (figure 4-67). SEM micrographs (figure 4-68) show an
interconnected crack network through the YBa2Cu307.x/Y203 film on Si, and the
film surface was very rough. AES data (figure 4-69) show that Ba has diffused
from the film to the substrate. Y203 films deposited on (100) Si at 730 C (figure
4-70) are smooth and featureless, which indicates the microstructural changes in
figure 4-68 occur during the YBajQ^O^ deposition and oxygenation.
YBa7Cu?Q7 x/(YA10?1LaAlQ3. or Y7Q,)/AbSi7On_Films_Qn i
-an3 Substrates
A series of experiments were performed in which 800 thick Al6Si2013 barrier
layers were deposited on Si or (1102) A1203 substrates prior to growth of the
second barrier layer (LaA103, YA103, or Y203) and the YBa2Cu307_x film. For Si
substrates, the role of the Al6Si2013 was to create a layer with a slightly larger
thermal expansion coefficient than Si. On (1102) A1203, the lower thermal
expansion coefficient of Al6Si2013 (relative to A1203) altered the stress in the
second barrier layer and YBa2Cu307.x films. Because the electrical resistance of
the YBa2Cu307.x films were very sensitive to the extent of cracking within the


32
and superconducting transition of a YBa2Cu307_x film can be attributed to
microstructural defects such as high angle grain boundaries, interfacial phases at
the grain boundaries, and randomly oriented grains. An expression for the normal
state resistivity, p, of a single crystal YBajQ^O^ film free of microstructural
defects was proposed by Halbritter:26
p (7) = a '7*+ pt, (2-M)
where the superscript i denotes intragranular effects, a' is a parameter which
contains the temperature dependence of intragranular normal state resistivity, and
p0L is the intragranular resistivity exptrapolated to T = O K. For single crystal
YBajQtyO^ films, Halbritter reported p0L' = 0 and a1 = 0.5 ohmXcm/K.
However, many YBa2Cu307.x films, especially those grown on substrates which are
reactive or do not lattice match YBa2Cu307.x, are polycrystalline and contain grain
boundaries or other types of weak links which increase the percolation distance of
the electrical conduction network. Using a technique similar to Eom et al.,27
deviations from the ideal resistance vs. temperature curve for YBa2Cu307.x can be
attributed to specific types of microstructural defects. Halbritter expanded the
equation to separate the normal state resistivity into intergranular (e.g. grain
boundary or cracks) and intragranular contributions caused by electron scattering
from other electrons, lattice ions, or defects within the grain:


INTENSITY (ARBITRARY UNITS)
162
YBa2Cu307.x/LaAI03 _
on AI2O3(1210) i
LaAI03 on
AI2O3(1210)
O
ic3
/
YBa2Cu307x
LaAI03
ai2o3
+ Al203kp
/
20 30 40 50
DIFFRACTION ANGLE (20)
60
Figure 4-51. X-ray diffraction patterns from LaA103 and YBa,Cu,07_v/LaA10,
films on (1210) A1203.


197
Table 5-1. Superconducting transition temperatures (zero resistance) and critical
current densities for YBa2Cu307_x/barrier layer films on different
substrates. Dashed lines (---) indicate no data were taken.
Substrate
Barrier Layer
T0 (K)
Jc (amps/cm2)
(1102) LaA103
none
88
5X107
off-axis Y-Zr02
none
86
6.8X103
(100) SrTi03
none
89

(100) Si
Yo.i3Zr0870194
76

off-axis Y-ZrOz
Yo.i3Zr0i8701>94
87
1.5 X107
(1102) LaA103
Yo.i3Zr0g70194
88
3X107
(1102) A1203
Yo.i3Zr0870194
83
1.9 xlO4
(1210) A1203
^o.i3^r0870194
80

(0001) ai2o3
^0.13^r0.87^1.94
77

off-axis Y-Zr02
Zr02
89
9X104
(100) SrTi03
Zr02
88
3.6 xlO6
(1102) A1203
(100) SrTi03
83
2.5 xlO6
(1102) A1203
(110) SrTi03
77
5X103
(1210) A1203
SrTi03
52

(0001) A1203
SrTi03
51



BIOGRAPHICAL SKETCH
The author was born November 7, 1959, in St. Louis, Missouri. He received
a Bachelor of Science degree in ceramic engineering from the University of
Missouri at Rolla in 1982, and a Master of Science in ceramic engineering from
the University of Illinois at Urbana-Champaign in 1985.
276


147
Figure 4-39. Phase angle versus temperature for a YBa2Cu307_K/SrTi03 film on
(1102) A1203. Data taken at 36 gigahertz.


83
A variety of techniques were used to characterize the superconducting and
barrier-layer microstructures. Film orientation and interdiffusion at the barrier
layer/substrate and YBa2Cu307.Jt/barrier layer interfaces critically affected the
electrical properties of the YBa2Cu307_x films, and evaluation of the interfacial
reactions and diffusion phenomena which promoted various types of
microstructures were required in order to correlate film microstructure with
electrical performance. Several analytical techniques, including x-ray diffraction
(XRD), scanning Auger electron spectroscopy (AES), scanning electron
microscopy (SEM), and Raman spectroscopy were used to evaluate the film
microstructures. Because the beam/sample interactions and detection techniques
were different for each of the measurement techniques, various microstructural
features could be examined. By understanding the mechanisms by which the data
was generated, the sample volume which was probed by each technique, and the
factors which were likely to reduce the validity of the data (such as electron
charging in AES), a complementary set of data were obtained which uncovered
many of the microstructural features which influenced the superconducting
properties of the films. Similarity, electrical and magnetic measurements provided
essential information about the film microstructures, and the suitability of the
YBa2Cu307_x/barrier layer/substrate combinations for various devices. A basic
understanding of how the techniques work, and the potential sources of error are
essential in order to assess the data. In this section, a description of the various


212
mechanism. At lower temperatures, plastic deformation becomes more difficult,
so the thermal stresses become more elastic. From this analysis, we conclude that
estimation of the thermal stresses via the continuum mechanics model is too high
because it neglects plastic deformation in the high temperature regime. The
temperature below which elastic deformation becomes appreciable is a function of
the bonding strength at the YBajQ^O^/substrate (or barrier layer) interface.
While the model cannot be taken too literally, it predicts that if thermally induced
elastic stresses exceed the fracture strength of the film, cracking will occur. This
was experimentally observed in YBa2Cu3Q7_x films deposited on Si, and to a lesser
extent on (0001) A1203.
The slightly broadened transition from Tonset (88 K) to T0 (80 K) in the
YBa2Cu307.x/Y-Zr02 film deposited on (1210) A1203 resulted from the
polycrystalline Y-Zr02 barrier layer. The lack of an x-ray diffraction pattern from
Y-ZrOz suggests the film was fine grained and randomly oriented. We conclude
that the YBa^Q^O^ film must have had random in-plane orientation, and T0 was
lowered by poor coupling across high-angle grain boundaries.
Effects of Texture and In-Plane Alignment
The best electrical properties for YBa2Cu307.x films deposited on A1203
substrates were obtained using SrTi03 barrier layers. YBa2Cu307.x/SrTi03 films
on (1102) A1203, in which both films were deposited at 200 mTorr oxygen,
showed a metallic resistance vs. temperature behavior which extrapolated to 0


270
64. J.T. Cheung and H. Sankur, CRC Crit. Rev. in Sol. State and Mat. Sci. 15,
63 (1988).
65. D. Bhattacharya, Thesis, University of Florida (1991).
66. R.K. Singh and J. Narayan, Phys. Rev. B. 41, 8843 (1990).
67. T. Venkatesan, X.D. Wu, A. Inam, and J.B. Wachtman, Appl. Phys. Lett. 52,
1193 (1988).
68. R.E. Muenchausen, K.M. Hubbard, S. Foltyn, R.C. Estler, N.S. Nogar, and
C. Jenkins, Appl. Phys. Lett. 56, 578 (1990).
69. J.P. Zheng. Z.O. Huang, D.T. Shaw, and H.S. Kwok, Appl. Phys. Lett. 54,
280 (1989).
70. C. Girault, D. Damiani, J. Aubreton, and A. Catherinot, Appl. Phys. Lett.
54, 2035 (1989).
71. J.D. Jorgensen, M.A. Beo, D.G. Hinks, L. Soderholm, KJ. Volin, R.L.
Hitterman, J.D. Grace, I.K. Schuller, C.U Segre, K. Zhang, and M.S.
Kleefisch, Phys. Rev. B 56, 3608 (1987).
72. C.J. Jou and J. Washburn, J. Mater. Res. 4, 795 (1989).
73. R.K. Singh, J. Narayan, and A.K. Singh, J. Appl. Phys. 52, 3452 (1990).
74. G. Koren, E. Polturak, B. Fisher, D. Cohen, and G. Kimel, Appl. Phys. Lett.
55, 2330 (1988).
75. R.K. Singh and J. Narayan, J. Appl. Phys. 52, 3785 (1990).
76. C.B. Eom, J.Z. Sun, K. Yamamoto, A.F. Marshall, K.E. Luther, T.H.
Geballe, and S.S. Laderman, Appl. Phys. Lett. 55, 595 (1989).
77. R.H. Hammond and R. Bormann, Physica C 162-164. 703 (1989).
78. R. Liu, C. Thomsen, W. Kress, M. Cardona, B. Gegenheimer, F.W. de
Wette, J. Prade, A.D. Kulkami, and U. Schroder, Phys. Rev. B 52, 7971
(1988).
79. G. Burns, F.H. Dacol, C. Feild, and F. Holtzberg, Sol. St. Comm. XL 367
(1991).


RESISTANCE (OHMS)
159
Figure 4-49. Resistance versus temperature data for YBa2Cu307.x/LaA103 films
deposited on: () (1102) A1203; (O) (1210) A1203.


Figure 4-46. Auger line scan across crater edge of a YBa2Cu307.x/LaA103 film on (100) Si.
CT\


(2-43)
dUp = TdSp P W1 + 2^ dhf .
Substituting the expressions for dU and dl/ into equation 2-42, we obtain
adA^dlJto^-TdS^-Yt^flni+P'dV+ptdV* (2-44)
Assuming the volume of the system is constant, the Helmholtz free energy is
obtained:
pee** = s**dT+'E¡ pfa + odA ,
(2-45)
and
dFexces3.
(2-46)
For a crystal which is freely grown from its supersaturated vapor (i.e. no
substrate effects on nucleation, growth energetics, or kinetics), the equilibrium
shape is given by:38
1 2 3 t
X2 X3 Xt
= constant
(2-47)
where a{ is the surface energy of the ith face and A¡ is the distance from the center
of the crystal to the face. This result implies that a crystal will grow so as to
minimize its surface energy, and in the case of thin film growth on a substrate,
orientations with the lowest surface energies are the most energetically favorable
to grow.
The morphologies which a growing film adopts generally fall into one of the
following three categories.39 Film growth can be classified as two dimensional, in


Figure 4-59. Scanning electron micrograph of a YA103 film on (100) Si.


237
properties were observed when the YB^Ci^O^ films were predominately (001)
oriented, and free of microcracks.
The orientation and lattice parameters of the barrier layers dictated the
orientation of the YBa2Cu307.x films. For SrTi03 barrier layers deposited on
(1102) A1203 substrates, the SrTi03 texture was heavily influenced by the oxygen
pressure during deposition. SrTi03 films deposited at 200 mTorr 02 were (100)
oriented, and the subsequently deposited YBajCujO^ film grew with an (001)
orientation. When the oxygen pressure during the SrTi03 deposition was dropped
to 40 mTorr, the SrTi03 layer adopted the (110) orientation, and the YBa2Cu307.x
film was (103) oriented. In both cases, the YBa2Cu307_x film grew with the
orientation which minimized lattice mismatch at the YBa2Cu307.x/SrTi03
interface.
The normal state and superconducting properties of the (001) and (103)
YBa2Cu307.x films were dramatically different. The normal state resistance of the
(001) YBa2Cu307.x/(100) SrTi03 film deposited on (1102) A1203 was metallic,
with a sharp transition to the superconducting state at 83 K. The Jc was 2.5 X106
amps/cm2 at 4.5 K, and the microwave resistance was 10 milliohms at 36 GHz
and 65 K. These values indicate the YBa2Cu307.x film had good in-plane
epitaxy. By contrast, the normal state electrical resistance of the (103)
YBa2Cu3O7.x/(110) SrTi03 film on (1102) A1203 was only slightly metallic, with a
transition to the superconducting state at 77 K. At 4.5 K, the Jc was 6.8 X103
amps/cm2. The primary source of degradation in this film was microcracking,


97
The phonon energies are characteristic of particular vibrational motions, and
therefore can be used to identify the molecular environment in which the
vibrations occur.
In our experiments, a 3 milliwatt Ar+ laser was used to generate unpolarized
541.5 nanometer radiation, which struck the sample perpendicular to the sample
surface. The intensity and wavelength of the reflected light was measured with a
Jobin-Yvon photomultiplier, operating in the photon counting mode. In
YT^CuaQy.jj, the most prominant Raman spectra are at 335, 440, and 500
cm'1. The line at 335 cm'1 results from out-of-phase bending of the 0(2) and
0(3) atoms (see figure 2-16), and the scattering intensity is greatest when the
incident electric field is parallel to the <100> or <010> directions. The 440 cm'
1 line is caused by in-plane bending of the 0(2) and 0(3) atoms, and the line at
500 cm'1 is caused by the Cu(l) 0(4) vibration parallel to <001>. The 440 cm'1
and 500 cm*1 scattering intensities are greatest when the incident electric field is
in the <001> direction.121
A great deal of microstructural information regarding the oxygen content of
the film, as well as disorder and substitution on the Cu(l) sites, and the relative
fractions of (001) oriented vs (100) and (010) oriented grains was obtained by
comparing the relative peak heights and widths of of these films. Oxygen
deficiency in YBa2Cu307.x manifests as a decrease in the frequency of the 500 cm'1
peak, with a shift from 505 cm'1 for x = 0, to approximately 480 cm'1 for x = 1.
The 500 cm'1 peak is strongly influenced by the 0(1) atoms, and oxygen vacancies


MICROSTRUCTTJ RE/ELECTRICAL PROPERTY CORRELATIONS FOR
YBa2Cu307.x/BARRIER LAYER FILMS DEPOSITED ON A1203, SILICON, AND
YITTRIA-STABILIZED ZIRCONIA SUBSTRATES
By
CARL HENRY MUELLER
A DISSERTATION PRESENTED TO THE GRADUATE SCHOOL
OF THE UNIVERSITY OF FLORIDA IN PARTIAL FULFILLMENT
OF THE REQUIREMENTS FOR THE DEGREE OF
DOCTOR OF PHILOSOPHY
UNIVERSITY OF FLORIDA
1992

ACKNOWLEDGEMENTS
Several people have helped me complete this thesis. I would like to thank my
advisor, Professor Paul Holloway for being an excellent role model and showing me
how a scientist solves problems. I am especially grateful for his guidance during the
early stages of the project, and for letting me choose my own directions as the project
matured. I was fortunate to have very talented people on my committee, and
Professors Abbaschian, Anderson, Connell, and DeHoff each provided insights into
materials properties which were important to this project and will be valuable
throughout my career.
I am grateful for the help I recieved from co-workers within our group, especially
Kelly Truman and Ludie Hampton. Much of the data was collected at the Major
Analytical Instrumentational Center at the University of Florida, and Eric Lambers,
Wayne Aeree, and Richard Crockett did a superb job of keeping the instruments in
top condition, which made the data collection and interpretation possible.
I would like to thank Drs. Kul Bhasin, Felix Miranda, Mark Stan, and Crystal
Cubbage of NASA Lewis Research Center for their help. A large portion of this
thesis would not have been possible without their help.
n

My family provided a strong emotional base which allowed my to complete the
program. I would like to thank my parents, Gerhard and Lillian Mueller, for
stressing the importance of finishing a project. I would also like to thank my
brothers, Don, Lloyd, and Keith, my sister Jan, and their families for their
encouragement and for helping to keep things in perspective.
Finally, I would like to thank the friends who made my stay in Gainesville so
enjoyable. I will miss the Sunday afternoon soccer games.
in

TABLE OF CONTENTS
ACKNOWLEDGEMENTS ii
ABSTRACT vi
CHAPTER
1. INTRODUCTION 1
2. LITERATURE REVIEW 4
Devices 4
Microstracture/Superconductor Relations 18
Weak link Behavior 27
Nucleation and Epitaxial Growth 35
Thermally Induced Stresses 46
YBa2Cu307.x Film Growth 53
Barrier Layer Technology 67
3. EXPERIMENTAL TECHNIQUES 81
Film Growth by Laser Deposition 80
X-ray Diffraction 84
Scanning Electron Microscopy 86
Scanning Auger Electron Spectroscopy 86
Electrical Resistance Measurements 90
Critical Current Density 94
Raman Spectroscopy 96
Millimeter-Wave Transmission Measurements 99
4. RESULTS 101
YBa2Cu307.x on (1102) LaA103 101
IV

YBa2Cu307_x/Y-Zr02 Films on Si, Y-Zr02, and LaA103 Substrates 108
YBa2Cu307_x/SrTi03 Films on A1203 Substrates 137
YBa2Cu307.x/LaA103 Films on Si and A1203 Substrates 154
YBa2Cu307.x/YA103 Films on Si and A1203 160
YBa2Cu307.x/Y203 Films on Si, Y-Zr02, and SrTi03 Substrates .... 173
YBa2Cu307.x/(YA103, LaA103, or Y203)/Al6Si2013 Films on
Si and A1203 180
5. DISCUSSION 196
Intergranular Versus Intragranular Effects 196
Effect of Y-Zr02, Y203, and Zr02 Barrier Layers on Jc Values .... 201
Dependence of T0 and Jc on A1203 Substrate Orientation 209
Estimate of Stress and Cracking Due to Differential
Thermal Expansion 211
Effects of Texture and In-Plane Alignment 212
Millimeter-Wave Properties 214
Effects of Surface Energy 215
Effects of Lattice Matching 217
Effects of Oxygen Pressure on SrTi03 Growth 219
LaA103 Barrier Layers 225
Y203 Barrier Layers 228
YA103 Barrier Layers 229
Al6Si2013 Barrier Layers 231
6. SUMMARY AND CONCLUSIONS 236
APPENDIX A. CALCULATION OF X-RAY ABSORPTION DEPTH
FOR YBa2Cu306^ 241
APPENDIX B. FLOW CHART OF THE Y Al Si Cu O SYSTEM . 243
REFERENCES 266
BIOGRAPHICAL SKETCH 276
v

Abstract of Thesis Presented to the Graduate School
of the University of Florida in Partial Fulfillment of the
Requirements for the Degree of Doctor of Philosophy
MICROSTRUCTURE/ELECTRICAL PROPERTY CORRELATIONS FOR
YBa2Cu307.x/BARRIER LAYER FILMS DEPOSITED ON A1203, SILICON, AND
YITTRIA-STABILIZED ZIRCONIA SUBSTRATES
By
CARL HENRY MUELLER
December, 1992
Chairperson: Professor Paul Holloway
Major Department: Materials Science and Engineering
YBa2Cu307.x and barrier layer films were deposited on single-crystal silicon (Si),
A1203, yittria-stabilized zirconia (Y-Zr02), SrTi03, and LaA103 substrates. A pulsed
laser deposition process was used to deposit the films at a substrate temperature of
730 750 C, and the films were cooled in an oxygen ambient. The films were
characterized using resistance versus temperature, critical current density (Jc), x-ray
diffraction (XRD), scanning electron microscopy (SEM), Auger electron spectroscopy
(AES), and Raman spectroscopy.
Growth of barrier layers on Si and A1203 substrates prior to the superconductor
suppressed chemical interdiffusion between the superconductor and substrate. For
(1102) A1203, the best barrier layer was a SrTi03 film deposited at 200 mTorr of
oxygen. The YBa2Cu307.x film had a zero resistance temperature of 83 K, and the
vi

Jc was 2.5 X106 amps/cm2 at 4.5 K. The surface resistance was 102 ohms at 36
gigahertz.
On silicon substrates, YBa2Cu307.x degradation is aggrevated by thermal stresses
created by the difference in thermal expansion coefficients between YE^Q^O^ and
Si (13.2 versus 3.8 xlO^C, respectively), which causes microcracking in the
YBa2Cu307.x films. Cracking and interdiffusion were minimized by depositing a
YA103 barrier layer prior to YBa2Cu307_x. The thermal stresses were relieved by
viscoelastic relaxation in the YBa2Cu307_x film, and the T0 was 78 K.
The Jc values of YBa2Cu307.x films on Y-Zr02 substrates were increased by
depositing Y-Zr02 or Y203 barrier layers. YBa2Cu307.x/Y203 films on Y-Zr02
substrates had Jc values of 9x10s and lxlO7 amps/cm2 at 77 and 4.5 K. The Jc of
YBa2Cu307.x films deposited on a Y-Zr02 substrate without a barrier layer was
6.8 x 103 amps/cm2 at 4.5 K. The higher Jc values were attributed to pinning of the
magnetic flux by excess Y203 at high-grain boundaries.
vii

CHAPTER 1
INTRODUCTION
There are several commercial and military application for thin films which are
superconducting at temperatures above 77 K. Since most semiconducting
devices perform optimally near 77 K, the potential for high-speed devices with
low-attenuation superconducting interconnects is promising1. In another
electronics area, passive microwave and millimeter-wave devices patterned into
superconducting thin films dramatically outperform the resonators and filters
presently being used.2 As the technology for depositing superconducting films on
A1203 substrates improves, this performance will be further enhanced. Potentially,
the largest applications for superconducting films are in power applications.3
Transmission cables, motors, and high field magnets which are too costly to
operate at 4.5 K would be economically feasible at 77 K. Progress in
fabricating superconducting wires has been difficult because the wires currently
being fabricated suffer from brittleness and low critical current density (Jc) values
(3X104 amps/cm2 at 77 K). A better technique for making superconducting
cables may be to deposit superconducting films on fibers such as yittria-stabilized
zirconia (Y-Zr02) which are mechanically strong and have a similar thermal
expansion coefficient as the superconducting film.
1

2
For each of these applications, Jc values greater than ~ 10s amps/cm2 are
required.3 The sensitivity of Jc to interfacial phases and high-angle grain
boundaries in YBa2Cu307.x is one of the most significant problems, and it has
inhibited commercialization of YBa2Cu307_x films. To date, YBa2Cu307.x films
with Jc values above 10s amps/cm2 at 77 K have only been deposited on single
crystal substrates such as SrTi03, LaA103, and Y-Zr02 which nearly lattice match
YBa2Cu307_x. The ability to deposit films with high Jc values on randomly
oriented or polycrystalline substrates would be a tremendous boost towards
commercialization, since this would allow a variety of substrates with different
shapes and sizes to be used.
In this thesis, the literature is reviewed in chapter 2 in order to provide a
background for this work. The experimental techniques used to deposit and
characterize the films are described in chapter 3. The data original to this thesis
is presented in chapter 4, and is discussed in detail in chapter 5.
Chapter 2 shows how superconductivity evolves. The temperature at which
the resistance disappears (T0) and Jc are closely tied to the microstructure, and an
overview of the microstructural phenomena which most critically affects the
superconducting properties is introduced. The laser deposition process is
described, and the mechanisms which enable multicomponent films to grow with
the same stoichiometry as the target are presented. Chapter 2 concludes with a
survey of the work directed towards depositing YBa2Cu307_x films on various
substrate materials.

Chapter 3 describes the film growth technique, and the methods used to
characterize the films. Each of the characterization tools relies on different
physical phenomena to probe the film microstructures, so different aspects of the
microstructures were uncovered. A brief description of the operating principles of
each of the probes is given in this chapter.
The experimental results are presented in chapter 4. The sections are
arranged so as to compare the effectiveness of each barrier layer to induce growth
of optimal quality YBajCujQ^ films on different substrates. The text explains
what information the data is providing, and points out the most notable features
in each figure.
Chapter 5 compares and contrasts the data in order to explain the observed
phenomena. The primary objective of this thesis is to uncover the microstructural
features which were primarily responsible for the normal state and
superconducting properties. By using a variety of experimental techniques, we
arrived at a more clear understanding of film microstructures than if only one or
two techniques were used; thus we were able to correlate the film microstractures
with the electrical properties.

CHAPTER 2
LITERATURE REVIEW
Devices
Much of the interest in high temperature superconductivity is due to the large
number of applications which could benefit from replacing normal metals and
conventional solid state electronic devices with superconducting materials.
Basically, applications which could utilize high temperature superconducting thin
films can be divided into two categories: active devices which require one or
more weak link junctions, and passive devices which utilize the intrinsic properties
of superconductors for enhanced performance.3
In superconductors, zero resistance and macroscopic quantization occur
because the wave functions of the electron or hole pairs are coherently coupled to
each other. This coherence enables calculations of the phase difference between
two different points of the superconductor to be made.4 Weak-link junctions are
thin insulating or normal metal regions which separate two superconducting
volumes, and the lower density of superconducting charge carriers causes a
gradient in the phase of the superconducting electron (or hole) wave function
across the weak link.
4

5
A Josephson junction is a special type of weak link junction in which the
relationship between the supercurrent, Is, and the gradient of the phase of the
superconducting wave function (A) is given by:5,6
Ia = Jc sin (A). (2-1)
where Ic is the maximum superconducting current which can travel through the
junction. For the rest of the discussion on devices, it will be assumed that the
weak links are Josephson junctions. The current vs. voltage behavior of a
Josephson junction is characterized by a zero resistance supercurrent which
persists when Ad> is constant, and a non-linear voltage which appears when AO
varies with time. The voltage, V, across a Josephson junction in this state is given
by:
2e7=_ft_li^) (2-2)
2tc dt
where e is the carrier charge and h is Planck's constant.
The I-V plot for a Josephson junction is given in figure 2-1,7 and virtually
all of the active superconducting devices utilize the behavior predicted by
equations 2-1 and 2-2. A brief overview of some of the devices which use
Josephson junctions, and their operating principles, is presented below.
Oscillators are devices which generate a repeating voltage vs. time waveform.
Superconducting oscillators, in which the output voltage is controlled by the
frequency of the supercurrent travelling through the junction (equation 2-2), can

generate electromagnetic radiation at milliwatt-level powers,8 at frequencies up to
1012 Hz.
Figure 2-1. The current versus voltage characteristic of a Josephson junction.
Reference 7.
Radiation detectors are based on the principle that electronic radiation will
induce AC currents in the Josephson junction.7 Regardless of the orgin, currents
add algebraically, hence the DC critical current will be suppressed by the
radiation induced AC currents. In the finite voltage regime, the mixture of AC

7
Josephson supercurrent with induced normal state AC currents will produce sum
and difference harmonic currents within the junction. When the ac Josephson
frequency and radiation frequency are harmonics of each other, there is a mixing
harmonic, and a step in the current vs. voltage spectrum appears at zero
frequency and at the mixing frequency. By varying the voltage and thereby the
supercurrent frequency across the junction, a large range of frequencies can be
sampled via the conversion factor:
1 microvolt = 484 MHz. (2-3)
Surface Quantum Interference Devices (SQUID's) are used to detect weak
magnetic fields, and operate on the principle that the number of flux lines which
pass through a Josephson junction must be quantized.4,9 If the magnetic field
being detected is not a quantized number of fluxons, the Jc of the junction will
change so that the magnetic flux from the magnetic field plus the flux generated
by the current through the junction is quantized (figure 2-2). The magnetic
intensity of a fluxon, is given by:
= = 2.07 x 1015 Webers (2-4)
2e
Thus the I V profile of a Josephson junction is modulated by the magnetic field.
In practice, the magnetic sensitivity of a dc SQUID is significantly improved by
placing two Josephson junctions in parallel (figure 2-3), which makes the loop
defined by the film plus the two junctions the area through which the modulating

8
magnetic flux travels. This dramatically increases the area over which the flux is
measured. For two junctions placed in parallel,
1,(4 = /c(0)cos^l (2-5)
0
The high sensitivity of SQUID devices make them attractive for a variety of
biomedical, geological, and military applications.
l*oA HoA
APPLIED MAGNETIC FIELD STRENGTH, Ha
Figure 2-2. Periodic variations in the critical current density with increasing
magnetic field strength. Reference 4.

9
I
I
I
I
Figure 2-3. Schematic of a two-junction SQUID. Reference 4.
Because Josephson junctions can be changed from superconducting to
semiconducting and vice-versa by varying the current or magnetic field, they have
been used as switching devices in digital electronics.9 The intrinsic switching
speed of a Josephson junction is:

10
Switching speed (2-6)
2n A (7)
where A(T) is the temperature dependent superconducting energy gap. The
intrinsic switching speed for niobium-based Josephson junctions is 0.22
picoseconds. As an example of the increased performance which can be achieved
by replacing semiconducting switching devices with superconducting Josephson
junctions,10 a four-bit data processor made with GaAs transistors had a clock
speed of 72 MHz, and a power dissipation of 2.2 watts. A processor which
performed the same functions using Josephson junctions had a clock speed of 770
MHz and dissipated 5 milliwatts. More recently, superconducting electronics were
used to fabricate a four-bit shift register using 3 pm linewidths, which operated at
9.6 GHz and dissipated 40 microwatts. By comparison, devices made from GaAs
or Si which used 0.5 pm linewidths, dissipated approximately 100 milliwatts. The
lower power dissipation associated with Josephson switching devices is an
important advantage, since heat generation and removal limit the density and
bandwidth of circuits based on semiconducting electronics.
Superconducting devices employing Josephson junctions have great potential
and, in theory, could completely change the materials and operating principles on
which high speed electonics are currently based. However, reproducible and
reliable Josephson junctions are difficult to fabricate. Efforts to create high
speed, commercially acceptable switching devices based on low temperature
superconducting electronics have not been successful. Because the high
temperature superconductors are much more sensitive to microstructural defects

11
than their low temperature counterparts, it is unlikely that high temperature
superconductors will be used as switching devices in the near future. While
devices which utilize fewer Josephson junctions, such as electromagnetic detectors
and SQUID's are more feasible, difficulties in creating reliable Josephson
junctions remain formidable.
The materials challenges presented by the second group of devices based on
superconducting materials, passive devices, are more likely to be surmounted in
the near future. The first commercial applications for high temperature
superconductors will probably emerge from this group.3 The largest applications
of passive superconductors will be as transmission lines, delay lines, and filters for
receiving and transmitting electromagnetic signals.
The advantages gained by replacing normal metals with superconducting
materials are derived from the magnetic field penetration depth:11
X(T) =
MO)
[!-(-£-) 3
TV,0.5
(2-7)
where A(T) is the temperature dependent magnetic penetration depth, A(0) is the
penetration depth near 0 K, T is the operating temperature, and T0 is the
highest temperature at which zero resistance is observed. The response of a
superconducting material to an applied magnetic field is shown in figure 2-4. For
a superconductor, A is a material property, so it is independent of the device
operating frequency.

12
X
Figure 2-4. Variation of magnetic flux density at the boundary of a
superconductor. Reference 4.

13
In microwave and millimeter-wave transmission lines, there is an electric field
induced in the superconducting line because of the inertia of superconducting
electron pairs to the electromagnetic field. A surface resistance results because
the normal-state electrons are excited by the electric field, and the surface
resistance is qualitatively a measure of how much of the electromagnetic signal is
lost as heat in the transmission line:6
Joule losses
Surface resistance = Rs =
(2-8)
where Hsurface is the magnetic field intensity inside the superconductor. In
superconductors, the depth to which a magnetic or electric field can extend is
limited by the penetration depth, and the surface resistance (Rs) is given by:6,12
(2-9)
where A is a constant, o) is the frequency of the electromagnetic signal, T is the
temperature, kB is Boltzmann's constant, and A(T) is the superconducting energy
gap.
By contrast, the surface resistance of normal metals is given by:13
(2-10)
where fiQ is the magnetic permeability of the metal, and o0 = al + ia2 is the
complex electrical conductivity, with,

14
and
i =
1 + art
2_2
(2-11)
o<**
1 + (2-12)
ax and o2 are the conduction and displacement currents, and x is the mean time
between electron-phonon interactions. In normal metals, the skin depth (<5) is the
distance into the metal which an electric field can penetrate, and is a function of
the electrical conductivity:
v T7 l
(2uo0pw)-5
where c is the speed of light (3xl010 cm/sec) and fi is the magnetic permeability.
The dependence of the skin depth on electrical conductivity and hence
frequency means the dielectric constant of the metal changes with frequency. The
parameters pertinent to microwave signal transmission are contained within the
propagation constant, y,14 where
y = a +j p (2-14)
a is the attenuation constant, and is the wavenumber of the microstrip line. /?
is given by:

15
P = fi (2-15)
Kg
where Ag is the wavelength of the electromagnetic signal as it propagates in the
microstrip line. The speed at which electrical signals propagate through the metal
is given by the phase velocity, vp, where
(2-16)
H jiQux and e = e0e where fi0 and e0 are the magnetic permeability and
dielectic permittivity of free space, and (i00)'05 is equal to the speed of light.
Given that nt and er are the relative permeability and dielectric constant of the
microstrip material, the phase velocity of an electromagnetic signal propogating in
the metal is given by:
(ery
0.5
(assuming pr = 1)
(2-17)
Changes in the electrical conductivity, and hence dielectric constant of the metal
as a function of frequency, will also result in phase velocities which vary with
frequency. This phenomena, termed dispersion, causes electrical pulses composed
of various Fourier components to become spread out as they propagate along a
transmission line, and is the primary reason why normal metal conductors are
inadequate for long transmission lines. Dispersion in metallic interconnect lines
makes them inadequate for delay times greater than 1 microsecond, or for
wideband lines carrying short pulses (figure 2-5). For example, assuming typical

16
dimensions for a VLSI microstrip line (thickness = 0.1 0.3 pm, width =1-3
/im, and length = 1-10 /im) dispersion necessitates that pulse lengths longer than
100 picoseconds must be used. However, a superconducting line of the same
dimensions would have negligible dispersion for pulses longer than 1
picosecond.15,16
Figure 2-5. Calculated phase velocities as a function of frequency for
superconducting and aluminum microstrip lines, at 77 K. Reference
15.
The real part of the propagation constant, a, is the attenuation and results
from Joule losses within the microstrip line. In a superconducting material, the
penetration depth is constant, so a (decibel/cm) is linearly proportional to the
number of wavelengths/cm, and thus the frequency. In normal metals, a is
dependent on the skin depth, which is proportional to or0-5, so a increases more
slowly with increasing frequencies in metallic microstrip lines than it does in
superconducting lines. However, at microwave and millimeter wave frequencies,

17
the superconducting penetration depth is much less than the normal metal skin
depth, and it is only at (o > 1012 Hz that attenuation in a superconducting
stripline approaches a for normal metals (figure 2-6).16,17
In semiconductor devices, the reduced attenuation and Joule heating provided
by superconducting interconnect lines would permit interconnect lines to be scaled
down considerably, thereby reducing chip to chip propagation delays.
Figure 2-6. Surface resistances of YBajQigOy.* (dotted line) and copper (solid
line) films deposited on LaA103 substrates. Reference 17.

18
Microstructure/Superconductor Relations
The microstracture of a superconducting film has a large influence on the
electrical and magnetic properties, and understanding how the superconducting
properties are impacted by film microstructure will dictate the processing
procedures used to fabricate the films. In YBajQ^O^ the superconducting and
normal-state properties are anisotropic with respect to crystallographic direction,18
so film orientation is a critical parameter. The T0 and Jc of the film are
dependent on superconducting properties such as coherence length (£), and
penetration depth (A), which, in turn, are influenced by orientation, grain size,
and grain boundaries. Clearly, microstructural considerations play a large role in
defining the electrical and magnetic properties of the film. To understand the
relationship between microstructure and superconductive behavior, it is instructive
to review how superconductivity evolves. Although the physical processes
responsible for superconductivity in YBa2Cu307_x have not been uncovered, most
of the properties are well described by the microscopic Bardeen-Cooper-Schrieffer
(BCS) theory and the phenomenological equations which predate BCS, so the
properties of YBa2Cu307.x will be described in terms of conventional
superconductors.
In a normal metal, there is a repulsive energy between electrons, and the
energy levels are described by Fermi-Dirac statistics:19,20

19
(2-18)
e ** +1
where f(E) is the probability that a given electron state is occupied, E is the
energy, Ep is the Fermi energy, k is Boltzman's constant, and T is the
temperature. In normal metals, electrical conductivity is possible because there
are empty electronic states at energies greater than Ep, into which electrons can
hop and thus move through the lattice. Current is transported through a material
because the applied electric field raises the electron energy level of one terminal
relative to the other, thus increasing the probability that electrons will hop from
the high potential to the lower potential region of the sample. Electrical
resistance arises from processes which transfer kinetic energy to the crystal lattice
by electron-electron and electron-phonon collisions.
The mechanisms for current transport in superconductors are different from
those observed in normal metals. Superconductivity is a cooperative phenomena
involving many electron or hole pairs, and is possible because of electron (hole)-
phonon coupling, which creates an attractive force between electrons (holes). The
potential energy caused by an electronic transition from an initial state k, to
another state k', is given by:6

20
T,, N 4ttc2 %
K(g,a>)= :-J1-h
q2 + k2
o)2-o)2q
].
(2-19)
where V (q, between the k and k* states, aq is the phonon frequency at wave-vector q, and k is
the wave number of superconducting electrons at the Fermi surface. During an
electron transition between two states (k to k'), a phonon may be absorbed. If (o
< ft)q, and the transition lowers the potential energy of the system, the electron-
phonon interaction creates an energy gap in the energy vs. momentum spectrum,
and the most favorable way for two electrons with p > Pf (p = momentum at the
Fermi level) to lower their energy below 2Ef is to form a bound state, in which
electrons with equal and opposite momenta combine to form Cooper pairs. The
wavefunction for a Cooper pair is given by:4
- (M0)
p h
where W is the amplitude of the wavefunction (also known as the "order
parameter"), multiplied by the travelling wave expression in which P = electron
momentum, r = position, h is Planck's constant, and | *P |2 is the density of
superconducting electrons. Cooper pairs obey Bose-Einstein statistics, hence it is
energetically favorable for all the pairs can have the same momentum. Because
all the superconducting pairs have the same momentum, and thus the same
wavelength, superposition of the coherent waves results in another wave with the
same wavelength. Superconductors possess macroscopic quantization because all

the superconducting electrons have the same momentum. When a transport
current is impressed on the system, the quantum mechanical relationship
21
m*vs=2mva=
2n
(2-21)
(where vs = velocity of superconducting electrons, and kj = k2 = kn for the
superconducting electrons) produces long range order in the momentum. There is
no resistance because the superconducting electrons are coupled together, and the
energy required to break a Cooper pair into excited electrons (quasiparticles) is
greater than twice the energy gap between superconducting and normal state
electrons (= 2A(T)).
The size of the superconducting energy gap is dependent on the fraction of
electrons which are superconducting, and this fraction varies with temperature.
Gorter and Casimer derived:
(2-22)
where n^/n is the fraction of superconducting electrons, T0 is the temperature at
which resistance vanishes in weak magnetic fields, and T is the temperature. The
variation of the superconducting energy gap (A(T)) with temperature is given by:22
(2-23)
where a is constant equal to 3.06 if weak electron-phonon coupling is assumed.
The variation of the superconducting gap with temperature is shown in figure 2-7,

22
1.0
0.8
<
£ 0.4
0.2
0
T/Tc
Figure 2-7. The superconducting energy gaps of lead, tin, and indium versus
temperature. Reference 4.
and shows the gap is a maximum at T = O K. To illustrate the fraction of the
ideal superconducting bandgap which commercial high temperature
superconducting devices will be expected to operate at, we assume T0 to be 90
K, and the device operating temperature to be 77 K. These values correspond
to the temperature at which zero resistance is achieved in YBa2Cu307.x, and the
boiling point of liquid nitrogen. Using equation 2-23, we see that A(T)/A(0)
drops to 0.38 at T/Tc = 0.86, and disappears at T = T0. Increasing the
superconducting energy gap at 77 K is one of the prime motivations for studying
higher Tc materials such as the bismuth and thallium-based superconductors.
A parameter which critically affects film and device quality is the coherence
length (£), which is the decay length for the wave function created by the
formation of Cooper pairs,7
5 =
hvF
71A
(2-24)

23
where vF is the velocity of electrons at the Fermi level and A is the size of the
superconducting energy gap. An equivalent definition is that the coherence length
is the minimum distance over which the density of superconducting electrons can
vary, and hence the minimum distance between superconducting and
nonsuperconducting regions. The coherence length (£) is much shorter in
YBa2Cu307.x than it is in conventional superconductors such as Nb3Sn, and is
highly dependent on crystallographic direction (table 2-1). The short coherence
length in YBajQ^O^ (table 2-2) is primarily due to the small Fermi velocity of
the superconducting electrons, and an important consequence of the short
coherence length is that microstructural defects, such as grain boundaries,
impurity atoms, dislocations, or chemically unstable surfaces which create
imperfect or disordered regions of similar size to the coherence length, can
significantly alter the superconducting wave function, especially in the <001>
direction. By contrast, microstructural defects are considerably less degrading to
the order parameters of conventional superconductors.
Table 2-1. Normal state parameters of conventional metals (such as Nb3Sn) and
YBa2Cu307_x. Reference 7.
Parameters
Conventional
Metals
YBa2Cu307 x
II (001)
1 (001)
m*
1-1.5 mc
5 mc
25 mc
Er (eV)
5-10
0.3
0.3
kF (cm)1
108
5xl07
5xl07
vF (cm/sec)
1-2x108
107
2xl06

24
A final parameter which is important for understanding the magnetic and
electrical transport properties of superconductors is the magnetic penetration
depth, A, which is a measure of how deeply a magnetic field penetrates into a
superconductor. The depth to which a magnetic field can penetrate into a
superconductor is limited because the superconductor will generate circular eddy
currents, which create an internal magnetic field that nullifies the applied field. A
is defined as4
f~B(x)dx=XB0 (2-25)
where B0 is the applied magnetic field, and B(x) decays exponentially as it enters
the superconductor (figure 2-4). By virtue of the Maxwell equation,21
Vx£ = + 4nJ, (2-26)
dt
A is the maximum depth at which a transport current (J) can flow, hence it is
analogous to the skin depth in normal metals.
Table 2-2. Superconducting parameters of conventional metals (such as Nb3Sn)
andYBa2Cu307_x. Reference 7.
Parameter
Conventional
Metal
YBa2Cu307 x
II (001)
X (001)
T0 (K)
< 23
95
95
2A/kBT0
< 4.4
5-8
2-3.5
A/Ep
lo-4
2x1o1
lxlO'1
to (A)
lOMO4
15
7

25
The penetration depth varies with superconducting electron density, and hence
temperature. The Gorter-Casimir expression for A is:11
A(7) =
T 4 -
(2-7)
where A is the penetration depth at 0 K, and varies with film quality. For a
highly-(001)-oriented, 3000 thick YBa2Cu307_x film grown on LaA103, A0 = 1800
was measured.17 Films with larger fractions of non-(001) oriented grains, or
which contain high-angle grain boundaries, had larger penetration depths.
The way in which a superconductor responds to a magnetic field is dependent
on the Ginzburg-Landau parameter k, defined as4
k=0.966 (2-27)
*0
The magnetic field may be applied, or may be induced by the transport current (a
self-field). Superconductors with k < ]2 are classified as Type 1, while those
with k > ] 2 are Type 2. In both types, there is a magnetic contribution which
increases the free energy density, AG, while the electron ordering associated with
the formation of Cooper pairs lowers AG. If the AG associated with electron
ordering occurs over a shorter distance from the surface than the magnetic
contribution ( as in Type 2 superconductors), there will be a minimum in the free
energy density at the surface, and it becomes energetically favorable to form an
interface between the normal and superconducting regions (figure 2-8). Thus
there are isolated circular regions in the sample which are normal, while the rest

26
of the sample is superconducting. The coexistance of normal and superconducting
regions (i.e. the mixed state) enables superconductivity to be maintained in much
higher magnetic fields in type 2 materials than are allowed in type l.4 All of the
commercially useful superconductors are type 2 materials.
Normal Superconducting
Magnetic
Flux
Density
Number of
Superelectrons
(a) Penetration depth and coherence range
Free
Energy
Density
Magnetic
Contribution
Electron-ordering
Contribution
(b) Contributions to tree energy
Free
Energy
Density
(c) Total free energy
Figure 2-8. Orgin of negative surface energy. Reference 4.

27
Weak Link Behavior
Because of the short coherence length, microstructural defects significantly
affect the electrical properties of YBa2Cu307_x. Defects are regions in which the
amplitude of the superconducting wave function is depressed, and the density of
superconducting electrons is lowered. Since these defects reduce the
superconducting order parameter, reductions in T0 and Jc result if the transport
current is forced to travel through the defects.6
Defects are classified as either superconductor-insulator-superconductor (SIS)
or superconductor-normal-superconductor (SNS), depending on the electrical and
magnetic properties of the junction.4 For the case in which two superconductors
are separated by a thin insulating layer, there is a finite probability that
superconducting electron pairs will tunnel through the insulator and a
superconducting current will be maintained. Because of the insulating layer, the
wave function is no longer continuous between superconductors, but there is a
phase difference, AO, between the wavefunction at each of the interfaces which
determines the maximum supercurrent, Is, which can pass through the insulator:5
/5=icsinA (2-1)
If the transport current exceeds ic, AO is no longer constant with time, and a
voltage appears across the junction. The currents which result from the time
varying AO and subsequent voltage across the insulator are given by:4

28
Ch ^(Aft) | h 3(A 4ite dt2 dt
(2-28)
where C is the capacitance, h is Planck's constant, and Rn is the normal state
resistance of the junction. Ambegaokatar and Baratoff derived an expression
which shows that Jc is inversely proportional to the normal state resistance,22
(2-29)
where A(T) is the superconducting energy gap.
Coupling between the superconductor wavefunctions on each side of a
junction, and thus tunneling current density, are rapidly diminished by increasingly
thick insulator regions:23
(2-30)
where s is the insulator thickness (nm), a = 0.1 (eVy^/nm is the effective barrier
height, and the electron energy gap between valence and conduction bands in
the insulator. Many devices, such as detectors and oscillators, could be made
using superconductors if SIS junctions were more reproducible. However, since
superconducting areas within a distance smaller than the coherence length, £, of
the junction are often degraded, A(T) is not the energy gap for a perfect material
but is depressed because of defects near the junction. Thus Jc for a given junction
is difficult to control.
When two superconductors are separated by a normal metal (SNS structure),
superconductivity is induced in the normal metal via the proximity effect. The

29
proximity effect entails diffusion of superconducting electron pairs and normal
electrons across the interface to create a weak superconducting layer in the
normal metal. There is a coherence length for superconducting pairs in the
normal metal which is given by:23
hVw.
normal metal *
(2-31)
where vF is the Fermi velocity of the normal metal, and the electron mean free
path is assumed to be longer than the coherence length ("clean limit").
Conversely, if the electron mean free path is shorter than the coherence length
("dirty limit"), the coherence length is given by:
* uAt
b /
where 1 is the electron mean-free path. The critical current across an SNS
junction is given by:
(2-32)
, (2-33)
where B is a constant, and a is the normal metal thickness.
The short coherence length in YBa2Cu307.x is the primary reason why
microstructural defects have a significant effect on the superconducting properties.
Dimos et al. conducted a series of experiments designed to determine the
dependence of Jc and T0 on temperature and magnetic field, and reveal the types
of weak links responsible for reduced superconducting properties.24 For epitaxial

30
YBa2Cu307.x films grown on (100) SrTi03 substrates, where two substrates were
sintered together so as to deliberately produce a misorientation angle in the
film yet insure the <001> directions in the YBajQ^O^ film on both sides of the
sintered SrTi03 junction were parallel, variations in Jc as a function of grain
misorientation showed a dramatic drop as the grain boundary angle was increased.
Inside the grains, Jc values of 4xl06 amps/cm2 at 5 K were measured. However,
across the grain boundaries Jc values dropped from 4x 106 amps/cm2 at 0
misorientation to O.lxlO6 amps/cm2 at 12.5 misorientation.
Further increases in resulted in similar or lower Jc values, and did not
follow a systematic trend. The authors observed that J^/Jc8 was approximately
proportional to 1/ for < 20, and proposed a model in which the dislocation
spacing was the predominant factor in determining Jc. The dislocation array acts
as a partial barrier to superconducting electrons, or perhaps as an easy path for
flux flow, since the superconducting order parameter is depressed at the
dislocation cores. Although various studies have observed significantly larger Jc
values (> 5xl06 amps/cm2 at 77 K) in epitaxial (OOl)-oriented films as opposed
to polycrystalline films, this study was the first to indicate that larger Jc would be
observed in (001) films with in-plane epitaxy and low grain boundary angles,
compared to to (001) oriented films with random in-plane orientation.
Later experiments by Mannhart et al. on epitaxial, (OOl)-oriented YBa2Cu307_x
films grown on (100) SrTi03 substrates determined that the values of
intergranular Jc as a function of temperature could be fitted by the
Ambegtaokatar-Baratoff equation for a SIS junction in which an energy gap A(0)

31
of 5 meV was assumed.25 However, the Ambegaokatar-Baratoff expression is also
valid for dirty SNS junctions, provided that significant changes in the phase of the
wavelength only occur near or in the normal layer; hence the observed gap could
represent the reduced order parameter inside the normal layer. The authors
concluded that grain boundaries were dirty SNS junctions because Jc(0) was a
factor of 10 less than values predicted by the SIS model. On the other hand, Jc vs
B(T) for different orientations of the magnetic field indicated intragranular Jc
values were limited by flux creep across the grains, and intragranular Jc values
were determined by the density and pinning energies of the flux-pinning sites.
Because the sensitivity of electrical properties in the superconducting and
normal states to film microstructure is very pronounced, a great deal of
information about the film microstructure can be obtained from the resistance vs.
temperature data. Zero resistance occurs when there is a continuous pathway in
which the wave function of the superconductor is phase ordered along the entire
path. If the wave function is strongly coupled between grains, the transition from
the normal to the superconducting state is abrupt, and the film T0 is close to the
transition temperature observed in bulk samples. However, if the film
superconductivity is localized, meaning there are isolated superconductor volumes
such as grains separated by barriers at which the intergranular coupling is weak,
or if the coupling strength varies randomly at the different grain boundaries, the
transition to the superconducting state will be percolative and there will be a
finite resistance at temperatures below Tonset. Deviations from the ideal resistivity

32
and superconducting transition of a YBa2Cu307_x film can be attributed to
microstructural defects such as high angle grain boundaries, interfacial phases at
the grain boundaries, and randomly oriented grains. An expression for the normal
state resistivity, p, of a single crystal YBajQ^O^ film free of microstructural
defects was proposed by Halbritter:26
p (7) = a '7*+ pt, (2-M)
where the superscript i denotes intragranular effects, a' is a parameter which
contains the temperature dependence of intragranular normal state resistivity, and
p0L is the intragranular resistivity exptrapolated to T = O K. For single crystal
YBajQtyO^ films, Halbritter reported p0L' = 0 and a1 = 0.5 ohmXcm/K.
However, many YBa2Cu307.x films, especially those grown on substrates which are
reactive or do not lattice match YBa2Cu307.x, are polycrystalline and contain grain
boundaries or other types of weak links which increase the percolation distance of
the electrical conduction network. Using a technique similar to Eom et al.,27
deviations from the ideal resistance vs. temperature curve for YBa2Cu307.x can be
attributed to specific types of microstructural defects. Halbritter expanded the
equation to separate the normal state resistivity into intergranular (e.g. grain
boundary or cracks) and intragranular contributions caused by electron scattering
from other electrons, lattice ions, or defects within the grain:

33
p oo E (psi)t* pi) <2-35)
where ^(p#,) is the sum of the resitivities associated with all grain boundaries, and
(a'T + p0Ll) describes the intragranular resistivity. The percolation parameter, p,
accounts for increases in the percolation distance of the transport current resulting
from microstructural defects including grain boundaries and cracks. When p^ is
small compared to p(a1 + p0L), intergranular defects have little effect on p, and
separation of inter and intragranular defects is difficult. If the current pathway of
lowest resistance is through the grain boundaries, the net effect of a change in p^
on R vs. T is to shift the curve to higher resistance values, without changing the
slope (p values) of the curve. Extrapolation of the curve to 0 K will result in
higher values of ^(p#,) when grain boundary resistance becomes larger, but not so
large as to significantly alter the current pathways. However, when ^ip#,) is
comparable to or greater than the intragranular resistivity, the percolation
distance of the transport increases as the current selects pathways which avoid
grain boundaries or other microstructural defects such as cracks, thereby
increasing p.
The same defects which increase normal state resistivity are responsible for
depressions in T0 and J,..2829 Intragranular Josephson junctions resulting from
oxygen disorder or twinning locally depress the order parameter, and thus the T0
within the grain. The most damaging defects are usually intergranular, and are
observed at high angle (> 20 ) grain boundaries where dislocated regions often
have significant oxygen disorder.25 At high angle grain boundaries,

34
superconducting electrons are preferentially transported through the grain
boundaries at microbridges where the oxygen disorder is minimal (figure 2-9).
Interfacial impurity phases increase the separation between superconducting
regions, further degrading the interface and making it more difficult for
superconducting electrons to tunnel through the grain boundary. Because the
coherence length of superconducting electrons is longer in the <100> and <010>
than in the <001 > direction, T0 and Jc in (001) oriented films are not as sensitive
to grain boundaries as are films with other orientations. However, high angle
grain boundaries are highly dislocated regions where impurity phases tend to
accumulate, so the order parameter is depressed in (001) oriented films with
random in-plane orientation. By contrast, low angle grain boundaries are
generally free of interfacial phases and do not depress superconductivity. In
addition, 90 grain boundaries, which usually result from twinning or adjacent
(010), (100) and (001) oriented grains, generally do not contain significant
amounts of impurity phases, and hence do not reduce T0 or Jc values.30
Figure 2-9. Intergranular defect with one microbridge. Cuprate regions near the
grain boundary are normal conducting, and the superconducting order
parameter is diminished in these regions. Reference 26.

35
Nucleation and Epitaxial Growth
The superconducting properties of YBajQ^O^ are very dependent on
microstructure. A general section which describes nucleation and growth
processes is presented in order to establish a framework in which the electrical
properties of the YBa2Cu307.x films can be correlated with the growth processes.
Nucleation and growth of thin films is a large and complex subject, with many
variables which can potentially affect the microstructure and orientation of the
film. While it is difficult to obtain direct experimental evidence for many of the
parameters which determine growth mode, it is possible to speculate about the
forces which were operative during film growth by examining microstructural
properties such as surface morphology, grain orientation, and the distribution and
orientation of interfacial phases. Because superconducting and barrier layer films
have been deposited on several different substrates under a variety of growth
conditions, many of the growth mechanisms which determine film microstructure
can be deduced. In most deposition processes involving high temperature
superconductors in which the film is grown from vaporized constituents striking
the substrate, the experimental parameters which are varied include substrate
temperature, oxygen partial pressure, and the energies of the atoms and ions as
they strike the substrate. A general discussion of the fundamental processes
involved in film growth and how they are affected by growth conditions is

36
presented in order to establish a framework by which film microstructure evolves.
In the specific case of YBa2Cu307.x and barrier layer growth, correlations between
growth conditions and the underlying growth mechanisms will lead to an
understanding of how film quality can be optimized.
The crystallinity and orientation of a film can be manipulated by changing the
growth rate. In order to grow a film, more adatoms must stick to the surface than
are evaporated, hence there must be a flux causing a supersaturation of atoms or
ions reaching the surface and forming stable nuclei. The flux of atoms impinging
on the surface ( atoms ^ P(T)
cm 2 x sec (2nmkT)0S
(2-36)
(
where P(T) is the vapor pressure at the substrate surface, m is the mass of the
impinging species at the substrate surface, T is the substrate temperature, and k is
Boltzmann's constant. To achieve epitaxy, it is essential that the atom be able to
move freely across the substrate until it reaches a potential minimum. The jump
frequency, a), of an adatom on a substrate is given by:
(2-37)
where v = number of jump attempts/sec (typically 1013/sec), and ED is the
activation energy for surface diffusion. The mean stay time for an adatom is:32

37
(2'38)
where Ed is the energy for desorption of an adatom from the substrate. The
mean distance an adatom diffuses before being desorbed is given by:
<2-39)
where is the single jump distance. This result shows that the surface diffusion
length increases with decreasing substrate temperature because the mean stay
time is increased. However, epitaxy is not necessarily improved by lowering the
substrate temperature because the jump frequency is lowered, hence the rate at
which equilibrium is reached is lowered. When the number of adatoms diffusing
across the surface at a given time approaches the density of surface sites, energy is
lost by adatom-adatom collisions, and adatom-substrate interactions are
diminished, thereby decreasing epitaxy. For the remainder of this study, we
assume that the flux of impinging species was low enough, and the substrate
temperature sufficiently high that the number of adatoms diffusing along the
surface was negligible compared to the number of surface sites, hence the
adatoms were able to reach their lowest energy configuration and equilibrium
conditions prevailed.
Until the nucleus reaches a critical size, it is more likely that the atoms will
dissociate and join a different nucleus or desorb, thus sub-critical nuclei do not
participate in the film growth process. In a real deposition system, the flux of

38
atoms striking the substrate consists of clusters of atoms as well as single atoms or
ions. Although it is difficult to determine what is the smallest stable nucleus size,
it is logical to expect that larger clusters striking the surface are more likely to
remain and form stable nuclei than are smaller clusters or atoms, since there are
more overlayer-substrate bonds formed. In addition, increasing the substrate
temperature is expected to increase the critical nucleus size, so the distribution of
clusters or atoms which eventually form stable clusters is further skewed towards
the higher cluster sizes.33
The relationship between critical cluster size and substrate temperature is
important because the orientation of the entire film can be heavily influenced by
the size of the original clusters from which stable nuclei are grown. Theoretical
and experimental data on face centered cubic metallic films deposited on NaCl
substrates indicates that for a critical cluster size of 3, the atoms will adopt a
triangular arrangement, and the film will grow with a (111) orientation.34
Similarily, when the smallest critical nucleus size is 4 atoms, the nucleus will
arrange itself in a square or rectangular mesh, and (100) oriented growth is
favored. In general, the first monolayer will grow with the surface mesh initiated
by the stable cluster, and subsequent layers will maintain this orientation and
grow with the closest packed planes parallel to the substrate. Based on this
analysis, face centered cubic films in which the critical cluster sizes are 3 or 4
atoms are expected to grow with (111) and (100) orientations, respectively. It has
been observed that several metallic films which grow with a (111) orientation at

39
low temperatures will adopt the (100) orientation when deposited at higher
temperatures.15 Likewise, yittria-stabilized zirconia (Y-Zr02) films deposited on
(1102) A1203 grow with mixed (111) and (100) orientations at temperatures below
780 C, and are predominately (100) orientated when deposited at higher
temperatures.36
The relative surface free energies of the substrate, film surface, and interfacial
layer are some of the most important parameters which dictate the film
morphology, and whether film-substrate epitaxy is possible. In a thin film, the
surface energy can be a substantial portion of the total energy of the system. For
a planar interface between two phases, a and /3 37
(2-40)
adA = dUexcess- TdSexcessp
where a is the energy required to create a surface of area A, dUexcess and dSexcess
are the energy and entropy associated with creation of the new surface, and nx and
dnj are the chemical potentials and excess surface concentrations of species i at
the surface. The general expression for the surface energy is:
dU*" = dUtoua dUa -dU* ,
(2-41)
where
dUa = TdSa-PadVa PM*
(2-42)
and

(2-43)
dUp = TdSp P W1 + 2^ dhf .
Substituting the expressions for dU and dl/ into equation 2-42, we obtain
adA^dlJto^-TdS^-Yt^flni+P'dV+ptdV* (2-44)
Assuming the volume of the system is constant, the Helmholtz free energy is
obtained:
pee** = s**dT+'E¡ pfa + odA ,
(2-45)
and
dFexces3.
(2-46)
For a crystal which is freely grown from its supersaturated vapor (i.e. no
substrate effects on nucleation, growth energetics, or kinetics), the equilibrium
shape is given by:38
1 2 3 t
X2 X3 Xt
= constant
(2-47)
where a{ is the surface energy of the ith face and A¡ is the distance from the center
of the crystal to the face. This result implies that a crystal will grow so as to
minimize its surface energy, and in the case of thin film growth on a substrate,
orientations with the lowest surface energies are the most energetically favorable
to grow.
The morphologies which a growing film adopts generally fall into one of the
following three categories.39 Film growth can be classified as two dimensional, in

41
which case the film completely covers the substrate before another layer is
nucleated, and the film grows at a uniform rate normal to the substrate. Films
can also grow by a three-dimensional growth process, in which case the film atoms
agglomerate into islands and the islands coalesce to form the film. A third option
is for the film to grow by a mixture of two and three-dimensional growth. The
first growth mode, termed Frank-van der Merwe, is characterized by two
dimensional growth of the film. This occurs when the surface energy of the film
is less than that of the substrate and interface:
film ^ substrate + substrate-film interface ' '
In the Frank-van der Merwe growth mode, the surface free energy of the system
is lowered by replacing the substrate surface with a different surface which has a
lower surface energy. For most electronic applications, in which it is often
desirable to grow a thin film with the same structure and orientation as the
substrate, two dimensional growth is preferred because the adatoms will tend to
reside in potential minima along the substrate, thus preserving the pattern of the
substrate surface. The relative adatom-adatom or adatom-substrate bonding
energies also influence the film growth mode, with strong adatom-substrate
bonding favoring two dimensional growth. When the surface energy of the film is
high,

42
^film ^ substrate + substrate -film Interface
(2-49)
the surface energy is minimized by the formation of three dimensional islands,
and the islands coalesce to form the film. This is known as the Volmer-Weber
growth mode, and is characterized by strong adatom-adatom bonding. The
orientations of these films are less influenced by the substrate than in the case of
two dimensional growth, and generally have more grain boundaries.
The third mode of film growth, Stranski-Krastanov, occurs when
Gfilm & substrate + ^substrate -film interface
(2-50)
The Stranski-Krastanov mode contains features of both of the previous growth
modes, and is often characterized by two dimensional growth of the first
monolayer, followed by three dimensional growth of the rest of the film.
Assuming the film growth rate, surface energy, and adatom-substrate bond
strengths are such that two dimensional film growth predominates, the film will try
to minimize orsubstratc.nira interface by growing epitaxially on the substrate, thus
eliminating the energy required to create an additional surface at the interface.
However, if the film has a different chemical composition than the substrate, the
lattice parameter of the film will be different and there will be a strain induced in
the film as it attempts to adopt the lattice spacings of the substrate. Once the
elastic limit of the film is exceeded, the strain is relieved via formation of
dislocations at the interface and the film is no longer completely epitaxial with the
substrate.

43
There are a variety of methods by which the film will attempt to minimize
differences in lattice parameters with the substrate. Starting with the definition of
lattice misfit ( = f),40
lattice misfit (=/) = -(tbn.?ub^te (2-51)
substrate
where afllm and asubstrate are the lattice parameters of the film and substrate. It has
been observed that the orientation which a film adopts with respect to the
substrate is primarily a fimcton of lattice misfit, as well as the adatom-film and
adatom-adatom bonding strengths. If the lattice parameters of the film and
substrate are very close (generally < 1%) and both have the same crystal
structure, the film will align itself so the film directions in the plane of the
interface match those of the substrate. For larger mismatches, the film will adopt
an in-plane orientation such that the maximum number of lattice points will be
commensurate with the substrate.41 Theoretical and experimental data indicate
that the unit cell over which the film minimizes misfit and becomes commensurate
with the substrate may exceed 500 , hence the correlation between film and
substrate directions is not always obvious or simple to deduce. One commonly
observed method of minimizing strains induced by dissimilar surface meshes is to
strain the lattice of the film so as to match row spacings in one of the directions.42
The Nishiyama-Wassermand and Kurdjumov-Sachs matchings are shown in figure
2-10. The difference between the two orientations is the Kurdjumov-Sachs

44
orientation achieves better row matching by rotating the film surface mesh with
respect to the substrate mesh.
R
Figure 2-10. Matching of atomic rows with epitaxial configurations of dissimilar
rhombic meshes. The substrate unit cell PQRS is drawn in solid
lines, and the overlayer unit cells in dashed lines, (a) Nishiyama-
Wassermand orientation, with overlayer PABC matching rows
parallel to PR, and overlayer PERF matching rows parallel to SQ.
(b) Kurdjumov-Sachs orientation; after rotating the overlayer mesh
relative to the substrate, overlayer PABC matches rows parallel to
PQ and PS. Reference 42.
As the lattice mismatch continues to increase, elastic strain within the film
exceeds the shear stress limit, and strain energy is released by the formation of
misfit dislocations.43 Misfit dislocations usually consist of edge dislocations with
the glide direction parallel to the interface. The elastic strain which can be
accomodated before misfit dislocations form is a function of the film-substrate
bonding, with strong bonding leading to larger elastic strain accomodation and a

45
reduced tendency to form misfit dislocations. The strain energy per unit length of
an edge dislocation line is given by:
EgL-ln.*
(2-52)
4ji(1-v) r0
where G is the shear modulus of the film, b is the Burger's vector of the
dislocation, v is Poissons's ratio, R and r0 denote the outer and inner radii of the
strain field over which the dislocation acts. In bulk materials, R can be
approximated as the distance between parallel misfit dislocations, * b/f. In thin
films, the area over which a strain field acts is limited by the film thickness, and if
the film thickness is less than b/f thick, the strain energy of a dislocation can be
substantially reduced.
In films used for electrical applications, one of the most damaging aspects of
misfit dislocations is that they create grain boundaries which degrade the
electrical performance of the films. For pure edge dislocations, the angle between
two adjacent grains, 0, is given by:
sin 0 =
D
(2-53)
where D is the distance between dislocations. Misfit dislocations also reduce the
film-substrate epitaxy. During the initial stages of film growth, islands are highly
mobile and wil try to attain epitaxy with the substrate. If the substrate-film
bonding is weakened by misfit dislocations, there will be a tendency for the
islands to rotate slightly about the ideal in-plane epitaxial positions.52 The extent
of island rotation is given by:

46

where (p is the island rotation, f is the lattice misfit, and 6 is the rotation of misfit
dislocations. The maximum island rotation is obtained when 0 = 1, and hence = f. The result of island rotation is that films deposited on substrates in which
the lattice mismatch is severe will contain a much larger number of grain
boundaries than will films deposited on a more closely lattice matched substrate.
Several of the mechanisms responsible for film growth, and the energetic
parameters which dictate the orientation and microstructure of films, have been
presented in this section. Because several of the parameters are interdependent
and are difficult to observe experimentally, it is impossible to provide a
quantitative analysis of the effect of each variable. However, a general
understanding of the parameters necessary to promote two dimensional, epitaxial
growth provides considerable insight into what defects are likely to emerge from
various film processing techniques, and suggest methods by which film quality can
be improved.
Thermally induced stresses
Thermally induced cracking is a problem which plagues thin films deposited
on substrates. There is usually a difference in thermal expansion coefficients
between the film and substrate, so if the temperature of the film is changed after
the deposition, thermally induced stresses will be generated in the film and
substrate. The least desirable way for the film to relieve these stresses is by

47
cracking, because cracks severely degrade the electrical properties of the films.
Thermally induced stresses have a significant impact on microstructural features
such as microcracking and grain size, and the microstructural processes by which
thermally induced stresses are relieved will be introduced.
Cracks are formed when the strain energy due to stresses in the film is greater
than the energy required to create new surfaces.45 The stress required to
propagate a crack is given by:
*Ef(Y,+Yp)]QS (2-55)
1 na
where Ef is the Young's modulus for the film, yE is the surface energy of the new
surface created by the crack, yp is the work of plastic deformation per unit area,
and a is the initial crack radius. A key point is that the fracture stress is
considerably lower at points where pre-existing flaws are present, such as at the
film-substrate interface.46
The model frequently used to calculate the stresses induced in thin films by
differences in the thermal expansion coefficients between the substrate and film(s)
begins with the assumption that the strain is entirely elastic.47 The dimensions of
the substrate are presumed to be unaffected by the presence of the film, and the
film is required to alter its shape in order to match the surface dimensions of the
substrate. Shear forces are generated at the film/substrate interface because of
the differing thermal expansion coefficients, and these shear forces cause the
substrate and film to bow slightly (figure 2-11). Assuming there is no plastic

48
deformation in the film or substrate, nor any slippage at the film/substrate
interface, stress in the film (af) can be obtained by measuring the radius of
curvature of the film/substrate combination:48
ar
EX
6(1 ~vs)Rtf
(2-56)
where ts and tf are the thicknesses of the substrate and film, vs is the Poisson's
ratio of the substrate, and R is the radius of curvature of the film/substrate
combination. There are two basic requirements imposed by continuum mechanics
which this model satisfies.49 First, the sum of the forces per unit width acting on
the film and substrate must be zero. The force per unit width is defined as (ft,
where i denotes either the substrate or one of the film layers. As an example of
how the stresses are distributed in the film and substrate, consider the case in
which the film stresses are tensile (figure 2-11). These film stresses will cause the
substrate to bow slightly, so the top of the substrate (adjacent to the film) will be
in compression. The magnitude of the forces will vary within the substrate, with
the maximum compressive stress observed at the top, and the maximum tensile
stress at the bottom of the substrate. Near the middle, there will be a plane
through which the stress (and therefore strain) is zero, and this is called the zero
strain plane. In the films, the stress is assumed to be constant throughout the
thickness of the each film layer. In this example, the tensile eft forces contributed
by the film and the portion of the substrate below the zero strain plane must
balance the compressive (ft contribution of the upper part of the substrate,

49
SIR
Tensile
ESS
Compressive
Figure 2-11. Cross-sectional side view of a two-layered film on a substate,
showing how elastic film stresses are accomodated by the substrate.
Reference 47.
between the zero strain plane and the film-substrate interface. The second
requirement is that the sum of moments about the axis which runs through
the zero strain plane must equal zero. Continuum mechanics requires that the
sum of moments induced by various mechanical forces acting on a common axis
must cancel each other in order for the axis to remain stationary. In our case, the
axis running through the zero-strain plane is common to all the forces acting on
the films and substrate.

50
Several features regarding the stress distribution in a multilayered film are
implicit in the continuum mechanics model. Since the oftf force for each layer of
film is significantly less than adopt the surface dimensions of the relaxed substrate if there is no plastic
deformation at the film-substrate interface, or between any of the film layers. A
consequence of the small ot force of the film relative to that of the substrate is
that the stresses induced in each layer result almost entirely from interactions
between the film and substrate. According to this model, the stress in any given
layer of a multilayered film is independent of the sequence in which the films
were deposited, and the thermally induced stresses in each of the layers is given
by:46
J4oirat)T (2-57)
where Ef and vf are the Young's modulus and Poisson's ratio for the film, a{ and
as are the thermal expansion coefficients for the film and substrate, and AT is the
change in temperature to which the film-substrate combination is subjected. A
key assumption of this model is that the chemical composition of each layer is
uniform. Since the composition is uniform, the thermal expansion coefficient
within each layer is constant, so the dimensions of the top of each layer will be
the same as the bottom. This assumption, coupled with the assumption of no
plastic deformation at any of the interfaces, leads to the conclusion that it is not
possible to reduce the thermally induced stresses in the outermost film by

51
depositing intermediate layers in which the thermal expansion coefficients
gradually bridge the gap between the substrate and outermost film.
Film stresses can significantly affect microstructural features such as grain size
and porosity in thin films. Using an energy minimization argument, Chaudhari50
showed that tensile stresses can significantly reduce the average grain size of a
film. Because the atomic packing density is lower in a region containing grain
boundaries than in a crystalline region, tensile stresses are reduced in films with
small grain sizes. A quantitative energy minima for the film is given by the
expression:
strain boundary
'boundary
(2-58)
where do is the initial grain size, d is the final grain size, e0 is the initial strain in
the film, and a is a normalized distance parameter used to compare the atomic
density of the grain boundary region with that of a grain. If the boundary region
has the same atomic density as the grain, then a is zero. If the boundary has a
monolayer of atoms missing, then a is 1. /? is a geometrical factor used to
characterize the shape of the grains. For grains with a square cross-section, ¡5 =
2. y is the grain boundary energy. For films subjected to tensile stresses, there is
some combination of strain and grain boundary energies for which the overall
energy of the film is a minimum. For films subjected to compressive stresses,
there is no overall energy minima, and the total film energy decreases as the grain
size increases.

52
In brittle materials, film cracking is caused by the motion and accumulation of
dislocations, which creates regions of high elastic energy and ultimately fracture.
The film microstructure can significantly alter the magnitude of stress required to
move dislocations and initiate plastic deformation.46,47,51 Numerous experiments
have shown that the stress at which plastic deformation is initiated is greater in
thin films than in bulk materials. This is because dislocations are pinned at the
film-substrate interface, thus increasing the stress required to move dislocations
through the film. The minimum stress required to move a dislocation in a film is
given by:48
a, = A [ ] [ ln(^-)]
f l2n(l-Vf)hi l(Vf+v4) b
(2-59)
where A is a geometrical constant which accounts for the angle between the
applied stress and Burgers vector, b is the Burgers vector of the dislocation, vf and
vs are the elastic shear moduli of the film and substrate, and h is the film
thickness. For brittle films, the stresses required to move dislocations are
associated with the onset of microcracking. Equation 2-59 shows that the stress
required to move dislocations is inversely proportional to the film thickness, and is
the reason why there is often a critical film thickness which, if exceeded, causes
cracks to form.
Depositing a second layer on top of the initial film significantly increases the
stress in the initial layer. This seems to contradict the elastic continuum model,
which asserts that the stress in a given layer is only a function of the differences in

53
thermal expansion coefficients between the each layer and the substrate, and is
independent of the other layers. However, the top layer creates an additional
interface which pins dislocations. Hence plastic deformation in the initial layer is
suppressed, and the yield strength is increased.
The elastic continuum model is a popular model for predicting film stresses in
the elastic limit. However, the assumptions of perfect bonding at each of the
interfaces, and chemical homogenity within each layer are not realistic in most
cases. Microstructural features such as chemical bonding at interfaces, grain size,
and film thickness will affect the stress at which the mode of deformation changes
from elastic to plastic, and must be taken into account when interpreting the
stresses required to form cracks.
YBaoCugO-, Film Growth
Shortly after the discovery of superconductivity in YBa2Cu307.x, the difficulty
of growing stoichiometric YBa2Cu307.x thin films by conventional methods such as
sputtering52,53 and electron beam evaporation54 became apparent. Depositing a
film with the correct Y:Ba:Cu ratios is a difficult process using these techniques.
Pulsed laser deposition became a popular technique for growing superconductor
films because stoichiometric films could be readily grown from a target with the
same composition.55,56 Pulsed lasers have generated stoichiometric films over a
wide range of pulse energies, laser wavelengths, and pulse durations. Nd-YAG
lasers operating at 1064, 532, and 355 nm and pulse energies ranging from 0.6 to

54
3.0 J/cm2 have generated YBa2Cu307.x films with the correct stoichiometry.57,58,59
YBa2Cu307.x films were also grown using a pulsed C02 laser (10.6 /im), but a
Y-enriched target was required to produce a stoichiometric film.60 By increasing
the pulse energy density of a C02 laser, stoichiometric YBa2Cu307.x films were
grown from a stoichiometric target, but the film morphology and superconducting
properties were seriously degraded by large globules in the film.61 The best films
have been grown using pulsed excimer lasers which operate at ultraviolet
wavelengths. The main problems associated with laser deposition is the presence
of fragments in the films which are ejected from the target, and nonuniform film
thicknesses.62 Shorter wavelength excimer lasers are popular because the
particulate size is smaller in these films than in films grown with longer
wavelength lasers.57
The ease with which stoichiometric films can be grown, and the presence of
particulates in the films both result from the laser-target interactions. There are
two basic ablation mechanisms: thermal and electronic.63 Thermal processes
result from rapid heating of the target surface and subsequent evaporation and
sublimation from the surface. Evaporation occurs when the laser power density
(Q0) exceeds the minimum power density necessary for evaporation, Q^64
r> 0.5
where pQ is the target density, D is the thermal diffusivity, U is the sublimation
energy, and r is the duration of the laser pulse.

55
Assuming the optical absorption depth (a1) of the target is small compared to
thermal diffusion length, the relationship a(2Dt)0^ > > 1 is valid. The
temperature rise of the target surface layer (where the thickness, t, of the surface
layer heated by the laser is defined as t = (2DT)0,5), can be approximated by
comparing the energy absorbed during the laser pulse with the thickness of target
surface heated by the pulse:
Ar=(i-)-
it
cvPo(2Dty
0.5
(2-61)
where R is the target reflectivity, I is the power density (W/cm2), t is the duration
of the laser pulse, and Cv is the specific heat.
Since the thickness over which laser energy is absorbed in the target is
inversely proportional to the absorption coefficient, target materials with high
absorption coefficients will attain higher surface temperatures because the energy
is confined to a smaller volume. The absorption coefficient is also a function of
the laser wavelength (A), and the absorption coefficient for YBa2Cu307.x increases
as A decreases.65 Congruent evaporation from a multicomponent target occurs
because the components cannot segregate over a region greater than the thermal
diffusive region (2DT)0-5 during the time over which the target is irradiated.
Because the temperature within the thermal diffusive region is high enough to
evaporate all the components, the entire region is evaporated and the film
stoichiometry is the same as that of the target.

56
Electronic mechanisms are also operative during laser deposition.66 Photons
with energies greater than the first ionization potential energies (7.726, 5.512, and
6.38 eV for Cu, Ba, and Y, respectively) will excite the target atoms, thereby
breaking bonds and causing ejection of ions. Experiments in which YBa2Cu307.x
film morphology and electrical properties of films grown at different wavelengths
showed conclusively that smoother films with fewer and smaller particulates, as
well as lower normal state resistivities and higher Jc values, were obtained in films
grown at shorter wavelengths (figure 2-12).57
The smaller particulate sizes which are observed when YBa2Cu307_x films are
deposited using shorter wavelength lasers is largely attributed to the higher
absorption coefficients in YBa2Cu307.x targets with decreasing A. The absorption
coefficients are 1.2 x 10s, 1.5 x 10s, and 1.7 x 10s cm'1 at 1064, 532, and 355 nm,
respectively.57 A higher absorption coefficient results in a thinner layer at the
surface into which the laser energy is coupled, thus creating a hotter
plume with finer fragments. However, the improved microstructures which result
from short wavelength radiation cannot be completely attributed to the slightly
larger YBa2Cu307.x target absorption coefficients. Strong absorption by
photofragments in short wavelength radiation, and subsequent fragmentation into
smaller particles, probably contributes to the smooth morphology of films grown
at shorter wavelengths.57
A great deal of effort has been made to understand the mechanisms by which
material is transferred from a YBa2Cu307_x target to the substrate via laser

57
ablation, and to optimize the film growing process. The film thickness across the
substrate varies as cosn0, (n ~ 4),67 which indicates the ejected particle distribution
from the target is highly peaked. In addition, the cation ratios in the film are no
longer stoichiometric at lateral distances greater than 20 degrees from the target
Figure 2-12. Resistivity as a function of temperature for three YBa2Cu307.x films
deposited on (100) SrTi03 substrates by a Nd:YAG laser and it's
second and third harmonics. Reference 57.

58
normal.66 Apparently the highly peaked, forward-directed component is primarily
responsible for stoichiometry in the growing film. One explanation for this
behavior is that the laser generated plasma results in rapid evaporation from the
target surface.65 The gas is initially at high pressures because the rate of
evaporation
is greater than the rate at which atoms and ions can leave the target surface. The
plasma expands into the vacuum, creating a supersonic molecular beam. Time-of-
flight measurements indicate that mean kinetic energies for Cu(l), Y(l), and
Ba(l) are 41.3, 43.4, and 47.9 eV, respectively for particles generated by a
193 nm laser.69 Time resolved optical spectroscopy measurements showed that at
oxygen partial pressures less than 1.3X10"4 Torr, the velocities of neutral and
ionized atomic species, as well as the diatomic species (such as YO, BaO, CuO)
were approximately 106 cm/sec. Increasing the oxygen pressure to 10'2 Ton-
reduced the velocities of the atomic and diatomic species to approximately
5x10s cm/sec.70
Although the ability to deposit a stoichiometric film is an extremely important
parameter for growing superconductor films with low surface resistivities and high
Je, there have been other advances in film growth techniques which have
significantly improved the film quality. Optimization of the growth temperature
and oxygen pressure during deposition, as well as the oxygen pressure and cooling
rate after the deposition, have generated YBa2Cu307.x films with nearly ideal
superconducting properties. Initially, amorphous films with the correct

59
stoichiometry were grown in vacuum (typically < 10"5 Torr) onto unheated
substrates. The films were then post annealed in flowing 02 at 850 900 C,
thereby forming the tetragonal YBa2Cu307_x phase.56
During cooling from the high temperature tetragonal phase, YBa2Cu307_x
undergoes large structural changes. Understanding what these changes are, and
how they are affected by temperature and oxygen pressure are critically important
for optimizing growth of YBa2Cu307.x films. The temperature at which
YBa2Cu307_x transforms from the non-superconducting tetragonal phase to the
superconducting orthorhombic phase is dependent on oxygen pressure.71 In 100%
oxygen, the transition occurs near 700 C. The oxygen content and change in
oxygent content as a function of temeprature are shown in figure 2-13 for bulk
YBa2Cu307.x in 100 % oxygen. If the oxygen pressure is lowered to 20% oxygen,
the tetragonal s* orthorhombic transition temperature is lowered to 670 C, and
in 2% oxygen the transition is depressed to 620 C. Ordering of the oxygen
atoms into one-dimensional Cu-O chains along the <010> direction is the
primary mechanism responsible for the tetragonal it orthorhombic transition.
The dramatic increase in the <010>, and decrease in <100> lattice parameters
as the orthorhombic phase is formed are shown in figure 2-14.
The temperature at which the transition occurs affects the kinetics of the
phase change. The transition occurs via a nucleation and growth process, in which
the ordering of oxygen along <010> and lengthening of the <010> lattice
parameter begins at a grain boundary or free surface.72 If the transition

temperature is suppressed by decreasing the oxygen pressure, growth of the
orthorhombic phase will be slow, and rapid cooling of the sample will result in
60
7.0
6.9 h
6.8
6.7
x
O
w 6.6
O
CO 6*5
>-
- 6.4
x
O
6.3
6.2
6.1
6.0
o,
Orthorhombic
<
I
Tetragonal
>-
\
\ T.
\
\
(b)
dOx
dT
\
\
\
\
\ -

l
J
200 400 600 800
TEMPERATURE (C)
1000
Figure 2-13. Total oxygen content, and change in oxygen content as a function of
temperature for YBajQ^O^. Reference 72.

61
Figure 2-14. The [100] and [010] lattice parameters of YBa2Cu307.x versus
temperature for a bulk sample heated in 100% oxygen. Reference
71.

62
incomplete growth of the orthorhombic phase. Rapidly cooled YBa2Cu307.x
samples will be comprised of orthorhombic nuclei surrounded by an oxygen
depleted, tetragonal phase matrix. YBa2Cu307.x films prepared by heating an
amorphous film to 850 900 C in oxygen, then slowly cooling through the
tetragonal-to-orthorhombic transition, have T0 values of approximately 85 K on
nonreactive substrates such as SrTi03, and T0 * 75 K on the more reactive A1203
substrates.55,56 Post-annealed films on (100) SrTi03 substrates are polycrystalline,
but are preferentially textured with the (001) plane parallel to the substrate.
Major improvements in Jc values were accomplished by depositing the films
in-situ at elevated temperatures and controlled oxygen pressures. Epitaxial
YBa2Cu307.x films on (100) SrTi03 and (100) Y-Zr02 substrates were grown at
substrate temperatures ranging from 500 650 C and 200 mTorr 02, and
subsequently annealing the films at 450-500 C in 760 Torr 02 for 60 minutes.73,74
In-situ films grown on (100) SrTi03 and (100) Y-Zr02 had Jc values of 5X106
amps/cm2 and lxlO6 amps/cm2, respectively, at 77 K; both had T0 values of 90
K. In addition, films grown in-situ have much lower room-temperature
resistivities (160 ^ohmXcm) than post-annealed films (~ 1 milliohmXcm). This
behavior was attributed to the predominance of (001) orientation in the in-situ
film, whereas the post-annealed films contained a mixture of (001), (100) and
(010) oriented grains. The microstructures of in-situ films also formed an
epitaxial orientation with the substrate, whereas a "basket-weave" structure was
observed in post-annealed films.75 The basket-weave structure arises from the

63
growth kinetics of the (100) and (010) oriented grains. Both of these orientations
have the <001> axis, which is the slow growth direction, parallel to the substrate,
but the <100> directions are at 90 to each other.
For YBa2Cu307.x films grown in-situ, the substrate temperature and lattice
mismatch at the film-substrate interface significantly affect the orientation of the
films.76 Films grown at 640 C on SrTi03 and LaA103 substrates were primarily
(001) oriented, whereas films on MgO, Y-ZrOz, and A1203 substrates, which were
also deposited at 640 C, grew with an (001) orientation. Increasing the growth
temperature to 720 C resulted in (001) oriented films on SrTi03 and LaA103
substrates. The variations in orientation were attributed to competition between
minimizing the surface energy of the film, which favors (001) orientation, and
minimizing structural coherence at the film-substrate interface during the early
stages of growth. Both SrTi03 and LaA103 have the perovskite structure, as does
YBajQ^Oy.jt, and the lattice matching between these substrates and YBa2Cu307.x
is reasonably close. At low growth temperatures, the reduced surface mobility
and possibility of film-substrate coherence is not as energetically favorable when
there is a large lattice mismatch between the film and substrate, which is the case
for YBa2Cu307.x films deposited on MgO, Y-Zr02, and A1203 substrates. Hence
the surface free energies are minimized by incoherent growth of the (001)
oriented grains.
Above 670 C, the tetragonal, oxygen depleted phase of YBa2Cu307_x is
stable, and at a typical growth temperatures of 750 C the non-superconducting

64
YBa2Cu307.x phase is formed. At 750 C, the perovskite lattice is
thermodynamically stable at oxygen pressures greater than 150 mTorr.77 Below
this pressure, YBa2Cu307.x decomposes into its component oxides, hence in-situ
growth is dependent on oxygen pressure, and is usually performed at a P02 of
approximately 200 mTorr (figure 2-15). The tetragonal-to-orthorhombic transition
temperature is a function of oxygen pressure, with a maximum temperature of 700
C. This transition is usually induced by backfilling the vacuum chamber to 10 -
760 Torr 02 after the deposition, and slowly cooling to 450 C. The film is kept
at 450 C for approximately 30 minutes to insure oxygenation of the Cu(l) atoms,
and ordering of the 0(1) atoms in the <010> direction. In the tetragonal phase,
0(1) and 0(5) sites are randomly occupied by O atoms, whereas in the
orthorhombic phase the 0(1) sites are completely full and the 0(5) sites are
empty (figure 2-16).78,79
YBa2Cu307_x films deposited in-situ at 650 750 C have higher Jc values than
post-annealed films. In-situ films have a higher ratio of (001)/(100) oriented
grains, and have better in-plane epitaxy in the <100> and <010> directions. It
has been established that the penetration depth is increased, and Jc values are
lowered by weakly coupled grains separated by high-angle grain boundaries or
non-superconducting interfacial phases. Hence the improved electrical properties
of films grown in-situ results from the reduction of weak links at the grain
boundaries.

65
TEMPERATURE (C)
Figure 2-15. Oxygen partial pressure versus temperature plot showing the critical
stability line for YBajQ^O^ at y = 6.0. Reference 77.

66
YBa2Cu3C>7
Chain
Plane
o-3,
Figure 2-16. Crystal structure of YBajCujO^ illustrating the CuO chains and
Cu02 planes.

67
Barrier Layer Technology
Because of the increased performance which could be realized by replacing
metallic interconnects and microstrip lines with superconductors, significant effort
has been expended towards finding deposition techniques which enable
YBa2Cu307.x films with high T0 and Jc values to be deposited on silicon and A1203
substrates. Sapphire is an attractive substrate material for microwave applications
because it is relatively inexpensive, mechanically strong, and has a much lower
dielectric loss tangent (tan <5) than other substrate materials (table 2-3).17
Table 2-3. Dielectric properties of substrate and barrier layer materials.
Reference 17
Material
Dielectric
constant
Loss
tangent
Sapphire (A1203)
9.4
1x10*
Silicon (Si)
12
1 x 103
Y-Zr02
27
6x 104
LaA103
25
5.8 x 104
SrTi03
305
MgO
9.65
4x 104

68
Devices fabricated on Si substrates would benefit from the lower attenuation and
signal distortion which superconducting interconnects can provide. Because
silicon and sapphire are chemically reactive with YBajQ^O^ and attempts to
grow YBa2Cu307.x films directly on these substrates result in interfacial phases
which damage both the substrate and film,80 substantial efforts to grow
intermediate barrier layers on the substrate prior to YBa2Cu307_x film deposition
have been made. In order to understand why some barrier layers are successful
and thus obtain the expertise necessary to design better barrier layer structures, it
is helpful to characterize YBa2Cu307_x films grown on inert single crystal
substrates. The first class of substrates are materials which are chemically inert to
YBa2Cu3O7.jp lattice match reasonably well, and have the perovskite structure.
SrTi03 is cubic with 2^ = 3.905 at 25 C,81 and LaA103 is rhombohedral with 2lq
= 7.586 and and included angle of 901'. However, LaA103 undergoes a cubic
rhombohedral transformation at 430 C, and at 650 C LaA103 is cubic with 2^ =
3.818 ,82 so in-situ films are grown on a cubic perovskite LaA103 substrate.
YBa2Cu307_x films grown at 650 C on (100) SrTi03 substrates show T0 at 89 K
and Jc = 7.5xl06 amps/cm2 at 77 K. Cross-sectional transmission electron
microscopy indicates the interface is atomically flat and abrupt, and the film does
not contain any secondary phases.83,84,85 The films grew with the < 001 > direction
normal to the substrate, the films are heavily faulted, with staggered,
discontinuous (001) layers. Close to the substrate, Y or Ba combine with Cu and
O to form perovskite subshells which are epitaxial with the substrate.

69
Perpendicular to the substrate, each perovskite subshell contains both Y and Ba,
but Y and Ba segregate into different domains along the interface. Away from
the interface, the (001) layers become continuous, but very wavy. Increasing the
substrate temperature to 720 780 C reduced the number of defects. Films
grown at the higher temperatures were completely (001) oriented, with less wavy
(001) layers, and Jc values of 2.2 xlO6 amps/cm2 at 77 K.
YBa2Cu307.x films with T0 = 90 K and Jc = lxlO6 amps/cm2 have been
deposited on (1102) LaA103 at 750 C.86 These films were highly (001) oriented
during the first 4000 of their growth. At greater thicknesses however, the
growth mode abruptly changed, and the dominant growth mode was with the
(100) and (010) planes parallel to the substrate. Computer simulations based on
nucleation densities and growth rates of (001) and (lOO)-oriented grains help
explain the experimentally observed film microstructures.87 Assuming the
nucleation density of (OOl)-oriented grains is 109/cm2, and for (100) and (010)
oriented grains is 2xl08/cm2, and the growth rate in the <100> and <010>
directions is 10 times larger than in the <001> direction, a microstructure
emerges in which the film surface is initially dominated by (OOl)-oriented grains.
As the film thickens, the (100) and (OlO)-oriented grains, which have a high
growth rate normal to the substrate, coalesce and the film is covered by these
orientations.
YBa2Cu307.x films have also been deposited on several substrates which have
distorted perovskite structures. The difference between these substrates and the

70
cubic perovskite materials (such as SrTi03) is that the angle between the <100>
and <001> directions is not 90 . YBa2Cu307.x films with T0 values of 92, 89,
and 88 K have been deposited on single crystal LaGaO^88 PrGa03,89 and
YbFeCXj90 substrates, respectively. In each case, the YBa2Cu307_x films were
strongly (001) oriented. Although each of these substrates contain elements which
are known to suppress superconductivity (La, Pr, and Fe), interdiffusion was not
observed.
Despite a large lattice mismatch, highly (001) oriented YBajQ^O^ films with
T0 = 92 K have been deposited on UNb03 substrates.91 The unit mesh of Y-cut
LiNb03, onto which the YBa2Cu307_x films were grown, is 5.148 x 6.932 . The Jc
values for these films were 2x10s amps/cm2 at 77 K, and the reduction in Jc
(relative to YBa2Cu307.x films deposited on SrTi03) was attributed to the high
concentration of Li in the film. The diffusion of Li was so rapid that the
concentration of Li within the YBa2Cu307_x film film was peaked at the surface.
The second group of substrates are materials which are chemically inert to
YBa2Cu307.x, but do not lattice match nor do they have the perovskite lattice. Y-
Zr02 amd MgO have been widely studied because high quality YBa2Cu307.x films
have been grown on these materials. Despite the large YBa2Cu307.x/Y-Zr02
lattice mismatch of 5.95 % in the YBajQijO^ <110>/Y-ZrO2 <100> direction,
YBa2Cu307_x films with T0 = 88 K and Jc = 1X106 amps/cm2 at 77 K were
deposited on (100) Y-Zr02.92,93 Tietz et. al.94 proposed a model for

71
YBa2Cu307.x growth on Y-Zr02 substrates which asserts that YBa2Cu307.x grows
in a manner which permits matching of the oxygen sublattices. This model also
proposes that the large lattice misfits are accomodated by 90 degree boundaries
and stacking faults. Norton et. al.92 reported that YBa2Cu307_x films with T0 = 90
K and Jc = 11,000 amps/cm2 at 77 K were grown on a polished, randomly
oriented Y-ZrOz substrate at 680 C. The films were strongly (001) oriented, and
the electrical properties of the films were more sensitive to the substrate
temperature during growth than were films grown on SrTi03 and LaA103.
Increasing the growth temperature to 730 C resulted in YBa2Cu307.x films with
Jc values of 1000 amps/cm2. The drop in Jc was attributed to increased chemical
interaction at the Y-Zr02/YBa2Cu307_x interface, with subsequent formation of
BaZr03. Presumably, BaZr03 diffused through the grain boundaries and reduced
intergranular conduction.
The lower Jc values for YBajQijO^ films on polycrystalline Y-Zr02 relative
to single crystal Y-ZrOz was attributed to the presence of high angle grain
boundaries. Garrison et al.95 demonstrated that by altering the deposition
conditions, YBa2Cu307.x films could be deposited such that matching of the
YBa2Cu307.x <100>/Y-ZrO2 <100> or <110> directions could be induced. When one
or the other of these orientations was dominant, YBa2Cu307_x films with Jc = 106
amps/cm2 at 77 K were observed. However, if both orientations were present in
the same film, the Jc values were only 102 104 amps/cm2 at 77 K. This

72
behavior was attributed to the high angle grain boundaries which resulted from
the mixed in-plane orientations of the YBa2Cu307_x films.
Films grown on (100) MgO substrates also showed T0 values of 89 K and Jc
= lxlO6 A/cm2 at 77 K.96,97 Similar to Y-Zr02, there is a large YBa2Cu307.x/
MgO lattice mismatch (5.9 and 4.2% in the YBa2Cu307.x <100> and <0io>/MgO<100>
directions, respectively). YBa2Cu307.x films grown onto MgO at 670 C were
predominantly (001) oriented. Unlike Y-Zr02, there was no evidence of
interfacial reactions at the YBa2Cu307.x/Mg0 interface. Similar to the case for
Y-Zr02, it was speculated that YBa2Cu307.x grows on MgO in a manner which
permits matching of the oxygen sublattices. The model also proposed that the
large lattice misfits are accomodated by 90 boundaries and stacking faults.
The third group of substrate materials is characterized by materials which
chemically react with YBa2Cu307_x, thus making growth of high quality
YBa2Cu307_x directly onto these substrates very difficult. Unfortunately, the two
substrate materials most widely used for electronic applications-(lOO) silicon for
integrated circuits and A1203 (sapphire) for micro and millimeter-wave electronics,
belong to this group. A 1.5 /m thick YBa2Cu307.x film grown in-situ at 700 C
onto a Si substrate showed T0 = 70 K, and an interfacial reaction layer ~ 0.5 fim
thick was observed.98 Auger electron spectroscopy detected Si at the film surface,
indicating Si diffusion through grain boundaries and microcracks. The poor film
quality was also attributed to microcracks generated by the large difference in
thermal expansion coefficients for Si (3.8xlO^/C) and

73
YBa2Cu307.x (13.6x 10"6/C). Chourasia et al." showed that in a YBajCi^O^
film deposited directly onto a Si substrate at room temperature, Si diffused into
YBa2Cu307.x near the interface, and formed a Si suboxide which depleted oxygen
from the CuO planes. After annealing the film at 860 C for 3 hours, the Si
diffusion was much more extensive, and Si02 was detected at the surface. The
authors concluded that oxidation of Si at the expense of CuO was the primary
reason that diffusion of Si into YBajQijO^ degraded the superconducting
properties of YBa2Cu307.x.
Better films have been grown directly on sapphire. YBa2Cu307_x films grown
on (1102) A1203 at 700 C showed T0 = 87 K and Jc values of 2xl06 amps/cm2
at 4.2 K.100 Attempts to deposit films at higher temperatures resulted in rapid
deterioration of the film because of interfacial reactions, probably BaO and CuO
reacting with A1203.
Because of the enormous technological advances which would result from
successful deposition of high quality YBa2Cu307.x films on Si and sapphire, a great
deal of research has been dedicated towards overcoming the interfacial reaction
problem. The most common approach has been to grow an intermediate barrier
layer prior to YBa2Cu307_x. The principal requirements of the barrier layer are
that it must be chemically inert to both the substrate and the YBa2Cu307.x film,
and should lattice match YBa2Cu307.x and the substrate. The materials which
have been most successful as barrier layers have been the materials which are also
the best substrate materials, such as SrTi03, Y-Zr02, and MgO. Although

74
excellent quality YBa2Cu307_x films have been deposited on single crystal LaA103
substrates, LaA103 has not been reported to be a good barrier layer material for
growth of YBa2Cu307_x on either A1203 or Si substrates. Attempts to grow
LaA103 films on a variety of substrates at 760 C showed that epitaxial films
could be deposited on SrTi03 and LaA103 substrates, while attempts to grow
LaA103 films on Si, A1203, and MgO substrates resulted in amorphous films.108
This indicates that matching both the lattice parameters and crystal structures at
the film-substrate interface are important parameters which critically influence the
crystallinity and orientation of the YBa2Cu307.x films.
Experiments in which LaA103 and YBa2Cu307.x were simultaneously deposited
onto an MgO substrate showed that T0 gradually dropped as the fraction of
LaA103 in the film increased.101 YBa2Cu307.x films which contained no LaA103
had a T0 = 87 K, while the transition temperature dropped to 77.5 and 30 K
for YBa2Cu307.x films containing 9 and 13 mole percent LaA103, respectively.
When the transition from Tonset to T0 was determined by measuring the magnetic
response of the film to a magnetic field (inductive response), the transition
remained fairly sharp (< 5 K) for all the films. However, when measured by the
electrical resistance technique, the transitions became much broader as the
fraction of LaA103 in the films increased. Apparently, the formation of LaA103
at the grain boundaries decreased coupling of the superconducting wave function
between the grains. The discrepancies in the widths of the temperature ranges
over which the transitions occured was attributed to the higher sensitivity of the

75
inductive measuring technique to intragranular conductivity, whereas the electrical
resistance technique is more sensitive to the presence of intergranular defects.
To date, Y-Zr02 has been the most successful barrier layer for YBa2Cu307.x
films grown onto Si substrates. 1 /m thick YBa2Cu307.x films with T0 = 82 K
have been grown at substrate temperatures of 650 C, using 0.1 ¡urn Y-ZrOa
barrier layers.83 The orientation of the Y-Zr02 layer greatly influences the
orientation of YBa2Cu307_x, and the best superconducting films are grown onto
(100) oriented Y-ZrOz barrier layers.103 Fork et al. showed that the ratio of
(200)/(lll) x-ray diffraction peaks from Y-ZrOz is strongly influenced by the
depostion temperature, and to a lesser extent, the oxygen partial pressure during
growth. The conditions under which highly (100) oriented Y-Zr02 barrier layers
were grown on Si were at a substrate temperature of 780 C and PQ2 = 7X10-4
Torr. Growth of (100) Y-Zr02 was also promoted by degreasing and cleaning the
silicon substrates in a flowing N2 hood, then transferring the substrates to the
deposition chamber via a N2 purged glove box, thus insuring a hydrogen
terminated Si surface. Thin YBa2Cu307.x films (305 ) deposited at 750 C, and
grown on 500 Y-Zr02 layers have T0 = 86 88 K, and Jc values of 2.2X106
amps/cm2 at 77 K. These values are comparable to those obtained from
YBa2Cu307.x films deposited on SrTi03 substrates. However, increasing the
YBa2Cu307_x film thickness to 1300 reduced the Jc to 1.5 x10s A/cm2 at 77 K.
Table 2-4 lists some of the more successful YBa^jO^/barrier layer structures
deposited on Si substrates, along with the film thicknesses. These data show that

76
reductions in T0 and Jc as the YBa2Cu307_x film thickness increases result from
film cracking caused by tensile stresses in the YBajQ^O^ films. The cracking
seriously degrades the superconducting film properties as the YBa2Cu307.x film
thickness exceeds approximately 500 .
In addition to the stress created by the thermal expansion coefficient
differences between Si and YBa2Cu307.x, stresses in the YBa2Cu307_x films are
exacerbated because the thermal expansion coefficient of YBa2Cu307_x is highly
anisotropic.104 Table 2-5 tabulates the thermal expansion coefficients of
YBa2Cu307_x in three directions, and in different temperature regimes. The
differences in thermal expansion coefficents in the <100> and <010> directions
are caused by ordering of oxygen atoms along the <010> direction, and
formation of the orthorhombic phase. The largest thermal expansion coefficients
are observed in the <001> direction. For YBa2Cu307_x films in which the <001>
direction is normal to the substrate, nucleation of a grain with the <010>
direction along a given direction parallel to the substrate will also result in
nucleation of another grain with the <010> direction orthogonal to the first
grain. This is the method by which the system minimizes the stresses caused by
differing thermal expansion coefficients in the <100> and <010> directions. If
an (001) oriented YBa2Cu307.x film is deposited on a substrate (such as SrTi03 or
LaA103) in which the thermal expansion coefficient is close to the average
thermal expansion of YBa2Cu307_x in the <100> and <010> directions, film
cracking does not appear to be a problem. However, for substrates with lower

77
thermal expansion coefficients (Si and A1203), YBa2Cu307.x film cracking caused
by thermally induced stresses is a significant problem.
The cracking problems caused by the low thermal expansion coefficients of Si
have been circumvented by using a silicon-on-sapphire structure.105 YBa2Cu307.x
films grown onto a Y-Zr02 barrier layer, which in turn was deposited on a Si film
grown on a sapphire substrate, had Jc values of 4.6 xlO6 A/cm2 at 77 K for
YBa2Cu307.x film thicknesses up to 4000 A.
Y-ZrOz has also been successfully used as a barrier layer for YBajQijO^
films deposited on the (1102) plane of A1203. Highly (lOO)-oriented Y-Zr02 films
were deposited on (1102) A1203 at substrate temperatures greater than 780 C,
whereas growing the Y-ZrOz layer at lower temperatures increased the ratio of
(111)/(100) Y-Zr02 x-ray diffraction peak intensities.36 YBa2Cu307.x films
deposited on highly (100) oriented Y-Zr02 barrier layers had T0 = 90 K, and Jc
values of 1.2X106 amps/cm2 at 77 K.
SrTi03 has emerged as a very good barrier layer material for growth of
YBa2Cu307_x on (1102) A1203 substrates. 1.2 pm thick YBa2Cu307_x films grown
on a 4000 SrTi03 barrier layer, had T0 values of 86.5 K and Jc values of lxlO6
amps/cm2 at 77 K.106 X-ray diffraction indicated that the SrTi03 layers
preferentially grew with a (110) orientation, although the (200) peak was
significant. The YBa2Cu307.x films were (001) oriented. Secondary ion mass
spectroscopy showed a drastic reduction of A1 concentration in films grown on

Table 2-4. Superconducting transition temperatures and critical current densities for YBa2Cu307.x/barrier layer films
on silicon and LaA103 substrates.
Substrate
Barrier layer and
Thickness ()
YBa2Cu30,x
thickness (A)
Transition
temperature
(K)
Critical current
density
(amps/cm2)
Reference
Si
Y-Zr02 (1000)
10,000
82

83
Si
BaT i03/MgAl204
(3500/5000)

70

107
Si
BaTi03/MgAl204
(3500/750)
1000
86-87
6xl04 at 77 K
108
Si
Yp/^-ZrOj
(100/900)
600
82-84
lxlO6 at 77 K
109
Si
Y-Zr02
(500)
130
86-88
2xl06 at 77 K
103
Si
Y-Zr02
(500)
1350

1x10s at 77 K
103
LaA103

1300
88-90
5xl06 at 77K
73
o
00

79
Table 2-5. Thermal expansion of YBajCi^O**. Reference 104.
Thermal Expansion (xl06/C).
Dila-
tometer
average
<100>
<010>
<001>
Average
Orthorhombic
14.3
5.8
25.5
15.2
12.9
25-400 C
400-610 C
37.5
0.0
39.5
25.7
25
25-610 C
22.6
3.5
30.3
18.8
16.6
Tetragonal
11.5
17.0
13.3
10.9
25-800 C
SrTiOs barrier layers relative to superconducting films grown directly on A1203,
confirming that SrTi03 is an excellent barrier to A1 diffusion. Char et al.110 found
that growth of the YBa2Cu307.x/SrTi03 film on (1102) A1203 at 750 C produced
films with T0 = 86.5 K and Jc values of 2xl06 amps/cm2 at 74 K. Higher
YBa2Cu307.x deposition temperatures and hence better in-plane epitaxy, which
were made possible by the SrTi03 barrier layers, were credited as the cause for
improved electrical properties, relative to YBa2Cu307_x films deposited directly
onto (1102) A1203.

CHAPTER 3
EXPERIMENTAL TECHNIQUES
Film Growth by Laser Deposition
Barrier layer and YBajCi^O^ films were sequentially deposited at 730 750
C on silicon (Si), aluminum oxide (A1203), yittria-stabilized zirconia (Y-Zr02),
lanthanum alumnate (LaA103), or strontium titanate (SrTi03) substrates using a
pulsed laser deposition system. A Questek model 2560 pulsed excimer laser,
using KrF gas and operating at 248 nanometers, 30 nanosecond pulses, and 5
pulses/second, was focused to 2.5 3.0 Joules/cm2 with a 50 centimeter focal
length lens onto a one-inch diameter YBa2Cu307.x or barrier layer target. A
schematic diagram of the deposition system is shown in figure 3-1. The
stoichiometric YT^CujOy.* target was obtained from Ceracon, Inc. and was 96%
dense. The barrier layer targets were fabricated by mixing stoichiometric ratios of
the powders, calcining at 950 C, then pressing the powders into disks and re
firing at 950 C for 12 hours. The barrier-layer targets were approximately 65%
dense. Up to three targets could be mounted on a stainless steel holder. By
rotating the target holder, sequential films were deposited without breaking
80

PULSED LASER DEPOSITION
Vacuum Chamber
Excimer Laser
248 nm, 30 ns pulses
Focusing Lens
50 cm Focal length
Figure 3-1. Schematic diagram of the pulsed laser deposition system.

82
vacuum or reducing the substrate temperature. This holder did not allow
continuous rotation of the target during deposition, but the target was moved
slightly every 800 pulses to expose a new surface to the laser radiation. The
deposition temperatures were measured by a thermocouple spot-welded to the
heater block. The barrier layer films were approximately 1000 thick, and unless
specified otherwise, were grown in 40 mTorr 02. The YBa2Cu307_x films were
2000 3000 thick, and were always deposited at PD2 = 200 mTorr. After the
YBa2Cu307.x films were deposited, the chamber was filled with oxygen, and the
temperature was maintained at 730 750 C for 20 minutes in order to facilitate
the tetragonal-to-orthorhombic transition. The films were cooled to 450 C over
60 minutes, held at 450 C for 45 minutes at approximately 300 Torr of oxygen to
ensure complete oxygenation, then cooled to room temperature.
All of the substrates used in this study were single crystal. The SrTi03 and Si
substrates were cut parallel to the (100) planes, LaA103 was cut parallel to the
(1102) planes, and A1203 was cut parallel to the (1102), (1210), or (0001) planes.
The Y-Zr02 substrates were cut 5 12 degrees from the (100) planes, as
determined by Laue back-diffraction patterns. By depositing the YBa2Cu307.x
films on off-axis Y-Zr02 substrates, we increased the tendency for high-angle
grain boundaries to form in the barrier layer and superconducting films. This
feature enabled us to determine whether the barrier layers would passivate the
YBa2Cu307.x films from the defects introduced by these substrates, and allow
growth of superconducting films with high Jc values.

83
A variety of techniques were used to characterize the superconducting and
barrier-layer microstructures. Film orientation and interdiffusion at the barrier
layer/substrate and YBa2Cu307.Jt/barrier layer interfaces critically affected the
electrical properties of the YBa2Cu307_x films, and evaluation of the interfacial
reactions and diffusion phenomena which promoted various types of
microstructures were required in order to correlate film microstructure with
electrical performance. Several analytical techniques, including x-ray diffraction
(XRD), scanning Auger electron spectroscopy (AES), scanning electron
microscopy (SEM), and Raman spectroscopy were used to evaluate the film
microstructures. Because the beam/sample interactions and detection techniques
were different for each of the measurement techniques, various microstructural
features could be examined. By understanding the mechanisms by which the data
was generated, the sample volume which was probed by each technique, and the
factors which were likely to reduce the validity of the data (such as electron
charging in AES), a complementary set of data were obtained which uncovered
many of the microstructural features which influenced the superconducting
properties of the films. Similarity, electrical and magnetic measurements provided
essential information about the film microstructures, and the suitability of the
YBa2Cu307_x/barrier layer/substrate combinations for various devices. A basic
understanding of how the techniques work, and the potential sources of error are
essential in order to assess the data. In this section, a description of the various

84
experimental techniques, and the regimes in which they were used for this study,
is presented.
X-ray Diffraction
X-ray diffraction was initially used to verify that YBa2Cu307.x and barrier layer
films were being grown. Interfacial phases were also detected using x-ray
diffraction (XRD). A Phillips model APD 3720 x-ray diffractometer operating at
40 kilvolts and 20 milliamps was used to generate Cu Ka radiation of A = 1.54060
and 1.54439 . A graphite monochromater filtered out most of the Cu k/2
radiation. Peaks within the 20 range of 5 65 were detected, and the x-ray
detector was rotated at 3 per minute. Interplanar spacings were calculated using
the Bragg equation:111
nX = 2dsin0 (3-1)
where n is an integer, A is the photon wavelength, d is the interplanar spacing,
and is the angle (relative to the sample surface) at which the x-rays enter and
leave the sample. By matching the experimentally observed interplanar spacings
with those predicted by the Joint Committee of Powder Diffraction Standards
index, the phases and orientations of the films were determined. Although the
graphite monochromater eliminated over 99% of the Kfi radiation, samples which
produced extremely large Ka diffraction peaks, such as the single crystal
substrates, also produced measurable Kfi x-ray diffraction peaks.
The orientations of the YBa2Cu307_x films were highly dependent on the
orientations of the barrier layer films. To determine whether phase information

85
about the barrier layer or interfacial phases buried beneath the superconducting
film could be obtained, the x-ray attenuation depths for the various films were
calculated. Assuming the x-ray intensity decreases exponentially as it enters the
sample, the attenuation for each compound was calculated using the expression:
-f =exp[-(-£)p] (3-2)
h P
where I/I0 is the fraction of the incident x-ray intensity which penetrates to a
depth = t, (p/p) is the mass absorption coefficient for each phase, and p is the
density of the phase, (p/p) was calculated from the weighted fractions of the
mass absorption coefficients for the individual elements:
(tW = wi(-) +w2(t) + w3() +-+w(t)
(3-3)
P Pi P 2 P 3 P
where wn is the weight fraction of element n in the phase, and (p/p)n is the mass
absorption coefficient for element n. For the calculation, the penetration depth at
which I/Iq was equal to 0.368 was taken to be the absorption depth (= t). Of the
films examined in this study, YI^CujO** had the smallest absorption depth (=9
pm). Since the YBa2Cu307.x films were typically 3000 thick, and the barrier
layer films were 1000 thick, we concluded that diffraction data was obtained
from all of the films in the multilayered structure. The calculation of the x-ray
absorption depth for YBajQ^O^ is presented in appendix A.

86
Scanning Electron Microscopy
A JEOL JSM 35C scanning electron microscope (SEM) operated at 15 20
kilovolts accelerating voltage and 100 microamps beam current was used to
visually determine the surface morphology and microstructures of the films.
Microstructural features are readily imaged in the SEM because secondary
electron detection is highly sensitive to the angle between the ejected electrons
and detector, so the number of secondary electrons detected varies as the
topography of the film changes.112 YI^Q^O^ films deposited on highly reactive
substrates or barrier layers tended to be cracked or have rough surfaces, whereas
superconducting films grown on inert substrates were smooth and featureless.
Correlations between SEM micrographs and electrical data were instrumental for
clarifying the types of microstructures which resulted in YBa2Cu307.x film
degradation.
Scanning Auger Electron Spectroscopy
Film uniformity and interdiffusion between the various layers and Si substrates
were analyzed using Auger electron spectroscopy (AES). A Phi model 660
scanning Auger microprobe, controlled by an Apollo domain series 3500
computer, was used to determine the film composition as a function of depth.
The operating parameters for the Auger were 5 kilovolts accelerating voltage, 25 -

87
35 nanoamps beam current, 132 volts emission voltage, and 40 60 microamps
emission current. Interdiffusion at the YBajCuaO^/barrier layer and barrier
layer/substrate interfaces was observed by ion sputtering a crater in the films, so
as to expose the barrier layer and substrate, then making a line scan across the
edge of the crater. With this method, errors induced by electrical charging and
Auger peak shifting in the insulating barrier layers were minimized.
Auger is often used to qualitatively measure the relative concentrations of
components within the top 10 of the surface. The type of element is
determined by the energy of electrons emitted as a result of the Auger process.
The Auger process is started by an incident electron beam with sufficient energy
to remove an inner shell electron, which creates a core hole. The ion energy is
reduced by filling the core hole with an electron from a more shallow energy
level, and emitting another electron from a shallow energy level. The energy of
the emitted electron is given by:113
Kinetic Energy = EA EB Ec (3-4)
where EA is the energy of the core level electron, EB is the energy of the shallow
level electron which fills the core hole, and Ec is the energy of the shallow level
electron which is emitted. Although Ea, Eb and Ec are all sensitive to the
chemical state of the atom, the time constant for Auger emission is short, so the
peaks are broad. Therefore the energies of electrons emitted via the Auger
process are less sensitive to the chemical state of the element than are electrons
emitted by other techniques, such as x-ray photoelectron spectroscopy. Hence

88
AES is widely used to determined which elements are present at the surface, with
limited determination of chemical state.
Quantification of the surface concentration is a difficult process because there
are many factors which influence the Auger yield. For an Auger transition from
species i at a site (x,y,z), where N¡ is the background Auger count and dN¡ is the
number of Auger electrons resulting from the transition:114
dN¡ = (incident electron flux of energy Eprimary at x,y,z)
x (ionization cross-section of EA for species i at Ep)
x (backscattering factor for Eprimaiy at the incident direction)
x (probability of decay of EA for species i to give the Auger
transition)
x (probability of no loss escape of electrons from region (x,y,z))
x (acceptance angle of analyzer)
x (instrumental detection efficiency).
To as much of an extent as possible, the Auger operating parameters were
kept constant in these experiments so that comparisons between the atomic
concentrations of different samples could be made. The incident electron flux was
dependent on the beam current, and was maintained at 30 40 nanoamps. The
ionization cross section is heavily influenced by the incident beam energy; low
beam energies are not adequate to produce core holes, and high beam energies
reduce the Auger yield from the shallow core levels. Generally, the optimal beam
energy is 3 5 times the binding energy of the deepest core level of interest. In

89
this set of experiments the beam accelerating voltage was always 5 kilovolts
because the sensitivity of the YLMM transition is highest at this accelerating
voltage.
Atoms can be ionized by backscattered and secondary electrons as well as
primary electrons, and the backscattering and secondary-electron yields are
sensitive to the chemical environment and electrical properties of the sample.119
Discrepancies in ionization cross section resulting from different backscattering
yields of the same element in different compounds is the primary phenomena
which limits quantitative Auger analysis. In this set of experiments, interdiffusion
betwen YBa2Cu307.x, the barrier layers, and the Si substrates were of interest.
Since the chemical environment and backscattering yields of each of the structures
were similar, semi-quantitative comparisons between the chemical compositions of
these structures could be made. Two final parameters which significantly effect
the Auger yield are the angle between the sample surface and the incident
electron beam, and the angle between the surface and detector. As the angle
between the beam and surface normal increases, the incident beam path length in
the surface region increases by a factor of sec 0,114 and the Auger yield increases.
Experimentally, this factor was kept constant by always keeping the angle between
the incident beam and surface normal at 60 .

90
Electrical Resistance Measurements
Electrical resistance versus temperature data were taken in order to correlate
normal state resistances, the onset of superconductivity (Tonsct), and the
temperatures at which the DC resistance dropped to zero (T0), with YBa2Cu307.x
film microstructures. Resistance measurements were obtained using a four-point
probe apparatus in which current was transported through the film by the outer
two terminals, and the voltage drop was measured across the inner two terminals.
The four probes were mechanically pressed against the sample, and either 10 or
100 microamps were transported through the sample. T0 was determined when
the voltage drop was less than 1 microvolt (R < 0.1 or 0.01 ohms. Resistance vs.
temperature data were also obtained on many of the samples at NASA Lewis
Research Center. The system at NASA was more sensitive than the one at the
University of Florida; gold contacts leads were wire bonded directly to the
superconducting films, which increased the sensitivity in the low resistance regime
near the transition temperature. The criteria for superconductivity in this system
was R < 0.001 ohms. Despite the differences, both measurement apparatuses
produced very similar normal state resistance and transition temperature data
when the same samples were tested on both systems.
For thin films, resistivity rather than film resistance is usually plotted because
resistivity is a material property, and the calculations used to obtain resistivity

91
values take into account film thickness, probe spacings, and sample geometry. In
this set of experiments, resistance was documented because the variations in
resistance caused by microstructural features overshadowed the relatively minor
changes in resistivity caused by varying YBa2Cu307_x film thicknesses and sample
geometries. An explanation of how resistivities are calculated, and how they are
affected by sample geometries is presented in order to support the hypothesis that
electrical resistance was the more appropriate parameter to monitor. Film
resistivity is given by:
p=FRt (3-5)
where F is a correction factor, R is the measured resistance, and t is the film
thickness. Assuming the probes are equally spaced, there are two correction
factors which will improve the accuracy of the resistivity measurements.115 The
first correction factor is given by the thickness correction factor, F(t/a). As the
sample length, a, becomes significantly greater than the film thickness, F(t/a)
approaches 1. Since the films were less than 3000 thick, and the length of the
substrates were usually greater than 1 cm, F(t/a) had a negligible influence on the
measurement. The second correction factor, F2, is a geometric correction factor
which takes into account the increased current densities in the sample caused by
narrow samples with relatively large distances between the probes. Films
deposited on narrow substrates have large a/d ratios, where d is the sample width.
If we assume the YBa2Cu307_x films were deposited on rectangular substrates in
which the length was twice the width (a/d = 2), the appropriate correction factors

92
for the film resistivity as the ratio of sample width to probe spacing (d/s) are
presented in table 3-1.
There were several practical difficulties which made calculating the
resistivities difficult. First, there was a trade-off between growing the highest
quality films and accurately assessing the film thickness, because masking off an
area of the substrate to create a step for profilometer analysis created thermal
gradients in the sample, which damaged the superconducting film quality. Second,
the substrates on which the films were deposited had a variety of shapes, so the
precise geometrical correction factor was difficult to establish. Although the
geometric factors were diverse, it is unlikely that they altered the resistivities by
more than + /- 50%, since the sample width/probe spacing ratios were
constrained by the design of the R vs. T measurement apparatus to be between 2
and 3. On the other hand, measured resistances often varied by an order of
magnitude or more, depending on the film microstructure and type of substrate
used. The additional information which could have been obtained by determining
the film resistivities would have been minor, and the fundamental
microstructural/electrical property correlations which controlled the electrical
properties were more fully uncovered by correlating data obtained by electrical
resistance measurements with other types of microstructural data.

93
Table 3-1. Resistivity correction factors as the sample width to probe spacing
increases. The sample length to width is kept constant (a/d = 2).
Ratio of sample width
to probe spacing (= d/s)
Resistivity correction
factor (= F2), assuming
a/d = 2.
1.50
1.4788
1.75
1.7196
2.00
1.9454
2.50
2.3532
3.00
2.7000
4.00
3.2246
5.00
3.5746
7.50
4.0361
10.00
4.2357
15.00
4.3947
20.00
4.4553
40.00
4.5129
00
4.5324

94
-Critical .CuirentPensity
For commercial applications such as interconnect lines, the critical current
density (Jc) is the most important figure of merit of a superconducting film. In
this set of experiments, Jc values were determined using the Bean model,116,117,118
in which circulating currents are induced in the superconducting films by an
applied magnetic field (Ha). In this technique, a DC magnetic field is applied
perpendicular to the film surface. Circular eddy currents are generated in the
film, and these currents propagate near the edges of the film in order to shield
the interior of the film from Ha. Because of the shielding currents, the magnitude
of H is less near the center of the film than at the edges, and the relationship
between H and Jc is given by:
(3-6)
10
To measure Jc, the intensity of Ha is increased beyond the point at which the
film can shield its interior from the magnetic field via circulating currents, so the
magnetic field completely penetrates the sample, and the film becomes a normal
conductor. Ha is then reduced, and the circulating currents reverse direction in
response to the change in Ha. When Ha = 0, the remanant magnetization (Br)
created by the circulating currents is given by:

95
fi=£ (3-7)
r 3
when the plane of the film is circular. Hp is the strength of the applied magnetic
field which completely penetrates the film. The polarity of Ha is then reversed,
and the magnitude is increased beyond -Hp. By cycling the sample through the
applied magnetic fields, a B vs Ha hysteresis loop is generated. The Bean model
asserts that Jc is constant throughout the sample, so
j = E (3-8)
e dr
where r is the radius of the film. Combining equations 3-7 and 3-8, then
integrating:
J = (3-9)
c 2 r
where AM is the remnant magnetization per unit volume (emu/cm3), and is the
magnetization difference for increasing and decreasing fields. The radius is given
in centimeters. After correcting for the units (1 emu/cm3 = 10 amps/cm), the
final expression is obtained:
j = iM (3-10)
c r
where Jc is given in amps/cm2. In practice, rectangular samples were cut, and the
films were oriented with the (001) planes perpendicular to Ha. r(cm) was the
radius of the largest circle which could be inscribed in the rectangular sample.

96
Baman_.SpggtrQ£CQpy
Raman spectroscopy was used to probe the intragranular microstructures of
the YBa2Cu307.x films. The intensities and positions of two of the Raman peaks
are sensitive to the oxygen content of YBa2Cu3O7.jp and were used to qualitatively
determine whether the films were oxygen deficient.119,120 Also, since each of the
Raman peaks are highly sensitive to the direction of the incident electric field, a
rough measure of the ratio of (001) vs. non-(OOl) oriented grains could be
determined. This feature was especially useful for films in which the grain sizes
were too small to produce x-ray diffraction peaks.
Raman scattering occurs when a photon is inelastically scattered by a
molecule, and the scattering is accompanied by either the absorption or emission
of a phonon with an energy equal to the difference in energies between the
incident and scattered photons:19
scattered photon ~ incident photon ^ phonon H)
When the incident light is in the visible region, the Raman spectra is dominated
by the change in polarizability as a molecule goes from its ground state to an
excited state. These polarization changes are caused by changes in the vibrational
energies of the molecule when it is excited by the electric field of the incident
light. The energies of the absorbed (or emitted) phonons are quantized, and
correspond to differences between the vibrational energy levels of the molecule.

97
The phonon energies are characteristic of particular vibrational motions, and
therefore can be used to identify the molecular environment in which the
vibrations occur.
In our experiments, a 3 milliwatt Ar+ laser was used to generate unpolarized
541.5 nanometer radiation, which struck the sample perpendicular to the sample
surface. The intensity and wavelength of the reflected light was measured with a
Jobin-Yvon photomultiplier, operating in the photon counting mode. In
YT^CuaQy.jj, the most prominant Raman spectra are at 335, 440, and 500
cm'1. The line at 335 cm'1 results from out-of-phase bending of the 0(2) and
0(3) atoms (see figure 2-16), and the scattering intensity is greatest when the
incident electric field is parallel to the <100> or <010> directions. The 440 cm'
1 line is caused by in-plane bending of the 0(2) and 0(3) atoms, and the line at
500 cm'1 is caused by the Cu(l) 0(4) vibration parallel to <001>. The 440 cm'1
and 500 cm*1 scattering intensities are greatest when the incident electric field is
in the <001> direction.121
A great deal of microstructural information regarding the oxygen content of
the film, as well as disorder and substitution on the Cu(l) sites, and the relative
fractions of (001) oriented vs (100) and (010) oriented grains was obtained by
comparing the relative peak heights and widths of of these films. Oxygen
deficiency in YBa2Cu307.x manifests as a decrease in the frequency of the 500 cm'1
peak, with a shift from 505 cm'1 for x = 0, to approximately 480 cm'1 for x = 1.
The 500 cm'1 peak is strongly influenced by the 0(1) atoms, and oxygen vacancies

98
in the Cu(l)-0(1) chains reduce the frequency of the Cu(l)-0(4) vibration. In
addition, substitution of Cu(l) by impurity atoms (such as Si or Al) broaden the
500 cm'1 peak. Qualitative information about the extent of oxygenation is also
obtained from the height of the 335 cm'1 peak. In fully oxygenated YBajQ^O?.*,
repulsive interactions between the 0(2) and 0(4) atoms reduce the peak height,
and vacancies on the 0(4) sites slightly increase the peak intensity. The shape of
the 335 cm'1 peak also provides qualitative information about the oxygen content
of the sample. Several studies have shown that the peak is highly asymmetric for
superconducting films and bulk samples, even at room temperature.120,121 This
asymmetry is caused by Fano-type interference between the phonon mode and the
electronic continuum in the Cu-O planes. The Fano-type lineshape122 is caused by
reduced electronic scattering at energies slightly above 335 cm'1, hence the
peak/background ratio is higher at energies greater than 335 cm'1. In oxygen
depleted, non-superconducting YBa2Cu307.x the 335 cm'1 peak is symmetric, which
has been attributed to the reduction in the free carrier concentration in the Cu-0
planes.
For completely (OOl)-oriented films, the peak at 500 cm'1 is not observed,
since the electric vector of the incident radiation is perpendicular to the <001>
direction; thus the ratio of the 335 cm4/500 cm'1 peaks can be used to
qualitatively determine the ratios of (001) to (100) or (OlO)-oriented grains. The
Raman sensitivity for the 440 cm'1 peak is highest in the <001> direction, so the
appearence of this peak indicates that a significant fraction of the grains have

99
non-(001) orientations. Correlations between the height of the 440 cm"1 peak and
oxygen content have not been established, so this peak was only used to identify
the presence of non-(OOl) oriented grains.
Millimeter-Wave Transmission Measurements
The amplitude and phase angle of millimeter-wave signals transmitted through
a YBa2Cu307.x/SrTi03 film deposited on (1102) A1203 were measured using the
millimeter-wave power transmission technique. From this data, surface resistance
and complex conductivity for the YI^Q^O^ film were calculated.
A 36 gigahertz signal, operating in the transverse electric (01) mode, was
propagated in a rectangular waveguide, at a power level of 16 milliwatts. The
substrate was clamped to the end of the waveguide with two waveguide flanges,
with the film side facing the incoming signal. The magnitude and phase of the
transmitted power were measured as a function of temperature with a Hewlett-
Packard 8510 Network Analyzer. Changes in the fraction of power transmitted
through the film/substrate combination, and the phase of the transmitted power,
were dominated by changes in the resistance of the YBa2Cu307_x film as it was
cooled from the normal to the superconductive states. A complete description of
the formulas used to calculate the transmitted power, phase angle of the signal,
surface resistance, and complex conductivity are presented in reference 12.

CHAPTER 4
RESULTS
The superconducting properties of YBajQ^O^ films must be nearly optimal
in order to be used in commercial devices. When deposited on substrates which
are chemically inert and have a similar lattice structures, lattice spacings, and
thermal expansion coefficients to YBa2Cu307.x, films with transition temperatures
near 90 K and Jc values greater than 4xl07 A/cm2 at 4.2 K can be grown.
However, when YBajCujO-^ is deposited on Si or A1203 substrates, there are a
number of microstructural defects which degrade the superconductivity, such as
high angle grain boundaries, grain boundary phases, and poor lattice matching
between YBa2Cu307_x and the substrate or barrier layers. Microcracking is a
particularly troublesome defect which results from the tensile stresses caused by
different thermal expansion coefficients between YBa2Cu307_x and the substrate.
The types of defects which are mainly responsible for YBa2Cu307.x film
degradation are often related to the substrate properties, and in this chapter we
will compare and contrast the effectiveness of various barrier-layer materials on
different substrates.
100

101
YB^Ci^Ot x on (1102) LaAlQ3
Before examining the properties of the various YBa2Cu307.x/barrier layer
structures, the electrical properties of a nearly ideal YBa2Cu307.x film are
reviewed in order to establish a benchmark to which the YBajQ^O^/barrier
layer films deposited on Si and A1203 substrates can be compared. Figure 4-1
shows resistance versus temperature data for two YBa2Cu307_x films deposited on
(1102) LaA103 substrates. The first film was cooled to 550 C before filling the
chamber with oxygen. The resistance versus temperature profile of this film was
non-linear, with a T0 of 79 K. For the second film, the chamber was filled with
oxygen while the temperature was still 730 C, and this film had nearly ideal
electrical properties. The normal state resistance was metallic, and extrapolated
to 0 ohms at 0 K. There was a sharp transition to the superconducting state at
88 K, and the critical current density was 5xl07 amps/cm2 at 4.2 K (figure 4-2).
X-ray diffraction data shows the film was predominately (001) oriented, with
minor (100) and (010) peaks also observed (figure 4-3). SEM micrographs (figure
4-4) show the film is smooth, and the most prominent features are debris from the
target.
The Raman spectra from YBa2Cu307_x films deposited on (1102) LaA103
substrates was sensitive to the deposition temperature and the film thickness.
Raman spectra from the film deposited at 730 C shows a pronounced peak at
500 cm*1 and an asymmetric peak at 335 cm*1 (figure 4-5). Figure 4-6 shows

8
CQ
3

o
w
u
s
CO
II
CQ
w

6
4
2
0
i i
O YBaaCus07_x ON (10Z) LaA10a. CHAMBER
FILLED WITH 0a AFTER DEPOSITION.
YBagCua07_x ON (ll02) LaA10#. FILM
COOLED TO 660 C BEFORE FILLING . '
CHAMBER WITH 0-. . '
* /
0 50 100 150 200 250 300
TEMPERATURE (K)
Figure 4-1. Resistance versus temperature data for YBajQ^O^ films deposited
on (102) LaA103. (O) Oxygen chamber filled with oxygen
immediately after the deposition, while the substrate temperature was
730 C. () Film cooled to 550 C before filling the chamber with
oxygen, and the lower T0 (79 K) was attributed to an incomplete
tetragonal-to-orthorhombic phase transition.

YBa2Cu307.x on La AI03 (1102)
Figure 4-2. Magnetization (M) versus magnetic field intensity for a YBa2Cu307_x film on (1102) LaA103. AM is the
width of the magnetization curve at a given magnetic field intensity, and Jc is derived from Jc = 15 AM/r,
where is the radius of the film. The Jc of this film was 5xl07 amps/cm2 at 4.5 K. _*

INTENSITY (ARBITRARY UNITS)
H*
8

20KU XI1000 0192 ^0U UFMSE
Figure 4-4. Scanning electron micrograph of YI^CujO^ film on (1102)
LaA103.

INTENSITY (arbitrary units)
Figure 4-5. Raman spectra for YBa2Cu307.x film deposited on (1102) LaA103.

INTENSITY (ARBITRARY UNITS)
250 300 350 400 450 500 550 600
RAMAN SHIFT (cm'1)
Figure 4-6. Raman spectra from 500 and 1000 YBa2Cu307_x films on (1102) LaA103, and from bare LaAlO
substrates.

108
Raman spectra from a LaA103 substrate, as well as 500 and 1000 thick
YBa2Cu307.x films deposited at 750 C. Increasing the deposition temperature
significantly reduced the height of the 500 cm'1 peak. For thicknesses greater than
1000 , Raman transitions from the substrate were masked by the film.
YBa3Qi3Q:, It/Y-ZrO, Films on Si. Y-ZrCX,. and LaAlQ3 Substrates
The most notable microstructural differences between YBa2Cu307_x films
deposited on Si versus Y-ZrOz substrates was the presence of cracks in
superconducting films on Si (figure 4-7). In all of the YBa2Cu307.x/barrier layer
films deposited on Si substrates, cracking was observed to some extent. SEM
micrographs of Y-Zr02 films deposited at 730 C on Si were featureless (figure
4-8), indicating that the cracking occurred during the YBa2Cu307.x film growth
and cooling process. We will show that strain induced by the difference in the
thermal expansion coefficients (table 4-1) was the principal cause of YBa2Cu307.x
film cracking on Si substrates, and reduction of the thermally induced stresses
significantly reduced YBa2Cu307.x cracking.
A primary objective of this study was to identify the mechanisms by which the
thermally induced stresses were relieved. Figure 4-9 shows the resistance vs.
temperature data for YBa2Cu307.x grown directly on an off-axis single crystal Y-
Zr02 substrate, and on Si with a Y-Zr02 buffer layer. After multiplication of the
normal state resistance values of the YBa2Cu307_x/Y-Zr02 film on Si by 0.061, the
resistance versus temperature plot is similar to that of YBa2Cu307.x films grown

109
directly onto Y-Zr02 substrates, although there was a wider AT between Tonset
and T0 in the YBa2Cu307.x/Y-Zr02 film on Si. AES data (figure 4-10) shows that
Y-Zr02 was an effective barrier layer for preventing interdiffusion between
YBa2Cu307.x and the Si substrate. The x-ray diffraction patterns obtained from
YBa2Cu307.x/Y-Zr02 on Si and YBa2Cu307.x on Y-Zr02 are similar, with textured
(001) YBa2Cu307.x predominant on both substrates (figure 4-11). There were,
however, small (100) diffraction peaks observed in the YBa2Cu307.x film on
Y-Zr02. Critical current density measuerements taken at 4.2 K show that the Jc
for this film was 6.8x10 3 amps/cm2.
p
v* -
x I
V ,
, *
% ~
. > <
V. ; *
w )
20Kt> XI2000 $060
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, 1 ID"UFMSE
Figure 4-7. Scanning electron micrograph of a YBa2Cu307.x/Y-Zr02 film on (100)
Si.

110
20KU X4000 3080 1.0U UFMSE
Figure 4-8. Scanning electron micrograph of a Y-Zr02 film on (100) Si.
Raman spectra of the YBa2Cu307.x/Y-Zr02 film deposited on Si (figure 4-12)
show an asymmetric peak at 335 cm'1, and a smaller peak at 500 cm'1. The
spectra from a YBa2Cu307_x/YA103 film on Si is also shown in this graph, and the
similarities between these two spectra and the spectra obtained from YBa2Cu307_x
films deposited on (1102) LaA103 indicates that the intragranular regions of the
YBa2Cu307.x/(Y-Zr02 or YA103) films deposited on Si were undamaged.
Deposition of Y-Zr02 barriers on single-crystal Y-ZrOz substrates prior to the
YBa2Cu307_x films significantly improved the normal state electrical behavior and
Jc values, as shown by the resistance versus temperature behavior in figure 4-13.

Ill
Both films show a sharp drop to zero resistance at 86-87 K, but the normal state
resistance of the YBajQ^O^* film deposited on a Y-Zr02 barrier layer was more
Table 4-1. Thermal expansion coefficients for YBa2Cu307.x, substrates, and
barrier layer films. Lattice parameter mismatches at the YBa2Cu307_x/
barrier layer and barrier layer/substrate interfaces are also calculated.
The lattice parameter for SflBa2Cu307_x is a = 3.82 .
MATERIAL
THERMAL
EXPANSION
COEFFICIENT
(xlO'6/C)
LATTICE PARAMETER
MISMATCH (%) WITH
RESPECT TO
YBa2Cu307.x
Si
YBa2Cu307.x
13.2

-0.5
Si
3.8
-0.5

ai2o3
6.7(||[0001])
5.0 (J-100011)
Y-ZrOz
13.5
+5.1
-5.3
SrTi03
9.4
-2.1
+ 1.5
LaA103
10
+ 0.8
-1.3
yaio3
?
+3.0
-3.4
y2o3
9.2
+ 1.9
-2.3
Y-Al-Si-O
3.1-7.0


i3
5.1
-0.7
+ 0.3

RESISTANCE (OHMS)
112
Figure 4-9. Resistance versus temperature data for a YBa2Cu307.x film deposited
on a Y-Zr02 substrate, and a YBa2Cu307_x/Y-Zr02 film on (100) Si.
Resistance values for the YBa2Cu307.x/Y-Zr02 film on Si are
multiplied by 0.061.

AUGER SIGNAL
(peak-background) / (background)
en

INTENSITY (ARBITRARY UNITS)
114
Figure 4-11. X-ray diffraction patterns from a YBajO^O^ film on Y-ZrO,, and a
YBa2Cu307.x/Y-Zr02 film on (100) Si.
(200)

INTENSITY (arbitrary units)
Ln
Figure 4-12. Raman spectra from YBa2Cu307.)/Y-Zr02 and YBa2Cu307./YA103 films deposited on (100) Si.

RESISTANCE (OHMS)
116
Figure 4-13. Resistance versus temperature data for a YBa2Cu307.x/Y-Zr02 film
deposited on Y-Zr02.

117
metallic, and extrapolated closer to zero ohms at 0 K than did the YBa2Cu307.x
film grown directly on the Y-ZrOz substrate. The more metallic behavior with a
barrier layer indicates lower grain boundary resistances than for films deposited
directly on Y-Zr02 substrates.
The Jc values of YBa2Cu307..x/Y-Zr02 films deposited on Y-Zr02 were close
to the best values reported for YBa2Cu307.x films deposited on (100) oriented Y-
Zr02.104 Figure 4-14 shows Jc versus H data taken at temperatures 77 to 4.5 K.
At zero field (0 tesla), the Jc ranges from 9 x10s A/cm2 at 77 K to 1.5 xlO7
A/cm2 at 4.5 K.
No x-ray diffraction peaks were detected from the Y-ZrOz barrier layer film
deposited on a randomly oriented Y-Zr02 substrate (figure 4-15). The diffraction
pattern from the Y-ZrOz film on the Y-Zr02 substrate was identical to that
obtained from a Y-ZrOz substrate. X-ray diffraction from the YBa2Cu307.x/Y-
Zr02 film on Y-ZrOz shows that the YBa2Cu307.x film was highly (001) oriented,
since only peaks from this orientation were detected. SEM indicates the film is
comprised of 0.3 -1.0 /m grains (figure 4-16), and the grain boundaries are much
clearer in this picture than in the YBajCu^^ film on LaA103 (figure 4-4).
High Jc values were also observed for YBa2Cu307.x/Y-Zr02 films deposited on
(1102) LaA103. The Jc of this film (3XlO7 amps/cm2) at 4.5 K was close to the
best values we have obtained in YBa2Cu307_x films deposited directly on (1102)
LaA103 substrates (5 x 107 A/cm2). Resistance versus temperature data are shown

CRITICAL CURRENT DENSITY (A/cm2)
MAGNETIC FIELD (T)
Figure 4-14. Critical current density versus magnetic field intensity for a YBa2Cu307.x/Y-Zr02 film on Y-Zr02. Data
taken at temperatures from 4.5 to 77 K.
00

INTENSITY (ARBITRARY UNITS)
119
Figure 4-15. X-ray diffraction patterns from Y-ZrO, and YBa,Cu,07 /Y-ZrO,
films on Y-Zr02.

120
4
41
20KU X5400 *0068 \.0U UFMSE
Figure 4-16. Scanning electron micrograph of YBa2Cu307x/Y-Zr02 film on
Y-Zr02.

121
in figure 4-17, and Jc versus H is given in figure 4-18. X-ray diffraction shows that
the Y-ZrOz film was predominately (100) oriented, but small (220) peaks were
also apparent (figure 4-19). The only diffraction peaks observed from the
YBa2Cu307.x film were from the (001) orientation. SEM indicates pinholes were
prevalent in the YBa2Cu307_x film. Through the pinholes, a fine-grained Y-Zr02
barrier layer is visible (figure 4-20).
In order to correlate yittria content in the barrier layers with Jc, Zr02 films
without yittria were deposited prior to the YBa2Cu307.x films. Growth of a Zr02
barrier layer on a Y-Zr02 substrate prior to the YBajQ^O^ deposition resulted
in a superconducting film with highly metallic normal state electrical properties,
and a zero resistance temperature of 88.9 K (figure 4-21). The x-ray pattern
from this film was identical to the pattern from the YBa2Cu307_x/Y-Zr02 film
deposited on Y-Zr02 (figure 4-22). However, the Jc was only 9.4 X104 amps/cm2
at 4.5 K. SEM shows the film is comprised of 0.5 -1.0 fim grains (figure 4-23).
The normal state resistance values and x-ray pattern of a YBa2Cu307.x/Zr02
film deposited on (100) SrTi03 were very similar to the YBa2Cu307_x/Zr02 film
deposited on Y-ZrOz. The resistance versus temperature curve was metallic, and
the T0 was 88.1 K (figure 4-24). Only x-ray peaks from the (001) YBa2Cu307.x
phase were detected; no ZrOa peaks were observed (figure 4-25). However, the Jc
of this film (3.8 X106 A/cm2) at 4.5 K was over a factor of ten higher than a

122
8
T
YBa2Cu307_x/Y-Zr02
on (1102) LaAlOg
m
S
W
o
m
o
53
<
H
m
M
CQ
W
ftt
2
o
/
/
0 L
0
o
_i i_i 1 1 ; 1
50 100 150 200 250
TEMPERATURE (K)
Figure 4-17. Resistance versus temperature data for a YBa^QuO,_/Y
on (1102) LaA103.
300
Zr02 film

CRITICAL CURRENT DENSITY (A/cm2)
Figure 4-18. Critical current density versus magnetic field intensity for a YBa2Cu307.x/Y-Zr02 film deposited on
(102) LaA103.
to
OJ

INTENSITY (ARBITRARY UNITS)
124
Figure 4-19. X-ray diffraction from Y-Zr02 and YBa^O? J Y-ZrO, films
deposited on (1102) LaA103.

125
. 7 j|jJ ft
"'ft
%
#
20KU XI 1000 0058 ~ UFMSE
Figure 4-20. Scanning electron micrograph of YBa2Cu307 J Y-Zr02 film on
(1102) LaA103.

RESISTANCE (OHMS)
126
100 150 200
TEMPERATURE (K)
250 300
Figure 4-21. Resistance versus temperature for a YBa2Cu307.x/Zr02 film
on Y-Zr02.

INTENSITY (ARBITRARY UNITS)
127
Figure 4-22. X-ray diffraction pattern from a YBa2Cu307.x/Zr02 film deposited
on Y-Zr02.

128
20KU XI 10 0 0 0077
Figure 4-23. Scanning electron micrograph of YBa2Cu307.x/Zr02 film deposited
on Y-Zr02.

RESISTANCE (OHMS)
129
20
16
12
8
0
i r
YBagCu307_x/Zr02 ON
(100) SrTiOa
$
/
-6-
J L
50 100 150 200 250 300
TEMPERATURE (K)
Figure 4-24. Resistance versus temperature data for a YBajCh^O^/ZrOj film on
(100 SrTi03.

INTENSITY (ARBITRARY UNITS)
130
YBa2Cu307.x/Zr02
ON (100) SrTi03
A Y Ba2Cu3C>7.x
s
CO
o
o
CM
O
O
Mm
I
*+*++**
P
O
O

X
X
X
10
20 30 40 50
DIFFRACTION ANGLE (20)
60
Figure 4-25. X-ray diffraction pattern from a YBa2Cu307_x/Zr02 film on (100)
SrTi03.

131
similar film deposited on Y-Zr02. SEM micrographs of the YBa2Cu307.x/Zr02
film on (100) SrTi03 (figure 4-26) indicate the film is smooth, and the grain
boundaries are not well defined.
Comparisons between the electrical properties and microstructural features of
YBa2Cu307.x/Y-Zr02 films deposited on different orientations of A1203 helped to
clarify the types of defects which degraded superconductivity. Figure 4-27 shows
the resistance vs. temperature behavior for YBa2Cu307.x/Y-Zr02 films deposited
on the (1102), (1210), and (0001) faces of A1203. The YBa2Cu307.x/Y-Zr02 film
deposited on (1102) A1203 had the most metallic normal state resistivity and the
highest transition temperature (T0 = 83 K), while the YBa2Cu307_x/Y-Zr02 film
grown on (1210) A1203 was slightly less metallic and the T0 was 80 K.
The poorest electrical performance for YBa2Cu307.x/Y-Zr02 films deposited on
A1203 was observed on the (0001) face. The room temperature resistance values
for films on this plane were ~ 5 times as high as for YBa2Cu307.x/Y-Zr02 films
deposited on the (1102) and (1210) faces of A1203, the resistance vs. temperature
curve was less metallic, and T0 was depressed to 77 K.
Despite the varience in the orientations of the Y-Zr02 barrier layers, x-ray
diffraction data indicates the YBa^jO^ layers were highly (001) oriented for
each of the films (figures 4-28, 4-29, and 4-30). The Y-Zr02 layer deposited on
(1102) A1203 was primarily (200) oriented, with a significant (111) Y-Zr02 peak.
Although the diffraction peak from (111) is larger than the peak from (200) Y-
Zr02, the film is mostly (200) oriented because the relative x-ray intensity factor

Figure 4-26. Scanning electron micrograph of a YBa2Cu307.x/Zr02 film on (100)
SrTi03.

RESISTANCE (OHMS)
133
YBa2Cu307_x/Y-Zr02 FILMS DEPOSITED
ON DIFFERENT ORIENTATIONS OF A10,
A 3
Figure 4-27. Resistance versus temperature data for YBa2Cu307.x/Y_-Zr02 films
on different orientations of A1203: (O) (1102); () (1210); (A)
(0001).

INTENSITY (ARBITRARY UNITS)
134
Figure 4-28. X-ray diffraction patterns from Y-Zr02, and YBa2Cu307.x/Y-Zr02
films on (1102) A1203.

INTENSITY (arbitrary units)
135
YBa2Cu307.x/Y-Zr02
on A!203 (0001)
o
o
o
in
o
o
<0
o
o
co
s
CM
O
O

JL
J
o
o
r
JL
o
8
UU
&
A
A
Y-Zr02 on
Al203 (0001)
o
o
CM
YBa2Cu307.x
Y-Zr02
10
20 30 40 50
DIFFRACTION ANGLE (20)
60
Figure 4-29. X-ray diffraction patterns from a Y-ZrO, and YBaoCuX^.,./Y-ZrO,
film on (0001) A1203.

INTENSITY (ARBITRARY UNITS)
136
Figure 4-30. X-ray diffraction patterns from Y-Zr02, and YBa2Cu307.x/Y-Zr02
films on (1210) A1203.

137
for (111) is greater than that of (200) Y-Zr02 (100 vs. 25). The highest
(200)/(lll) ratio was observed for the Y-Zr02 film deposited on (0001) A1203.
Although highly (200) oriented Y-Zr02 barrier layers have been reported to
improve the in-plane epitaxy of YBa2Cu307.x films grown on (102) A1203, the
electrical performance of YBa2Cu307.x/Y-Zr02 deposited on (0001) A1203 was
poor. Raman spectroscopy (figure 4-31) of the YBa2Cu307_x/Y-Zr02 films
deposited on the (1102) and (0001) planes of A1203 show that both films had
minor concentrations of (100) and (010) oriented grains, and there was no
significant difference between the two spectra. Scanning electron micrographs of
a YBa2Cu307.x/Y-Zr02 film on (0001) A1203 revealed small cracks (figure 4-32).
These cracks were not observed in films deposited on (1102) or (1210) oriented
ai2o3.
YBa3Cu3Q7 ./SrTiC^ Films on A13Q3 Substrates
YBa2Cu307.x films deposited directly onto both (100) SrTi03 substrates had
optimal electrical properties (figure 4-33). The normal state resistance was
metallic, and there was a sharp drop to zero resistance at 89 K.
SrTi03 was also very effective as a barrier layer for the growth of YBa2Cu307_x
on (1102) A1203. The orientation adopted by the SrTi03 barrier layer was highly
sensitive to the oxygen pressure during growth. At PG2 = 200 mTorr, SrTi03 grew
with the (100) orientation (figure 4-34), whereas dropping the PG2 to 40 mTorr

INTENSITY (ARBITRARY UNITS)
Figure 4-31. Raman spectra of YBa2Cu307.x/Y-Zr02 films on (102) and (0001) orientations of A1203.
w
00

139
Figure 4-32. Scanning electron micrograph of a YBa2Cu307.x/Y-Zr02 film on
(0001) ai2o3.

RESISTANCE (OHMS)
140
100 150 200 250
TEMPERATURE (K)
300
Figure 4-33. Resistance versus temperature for a YBa2Cu307.x film on (100)
SrTi03.

INTENSITY (arbitrary units)
141
<0
8
YBa2Cu307_x/SrT103
on AI2O3 (1102)
CO
o
CO
8
m
8
52 I
CM
8

CO
o



SITIOg

AI2P3
I

SrTlC>3 on Al203 (1102).
Deposited at 200 mTorr 02>
-Q
a
CM
O
1
o
o
1
1
I
Jl
11
10
20 30 40 50
DIFFRACTION ANGLE (20)
60
Figure 4-34. X-ray diffraction patterns for SrTi03 and YBa2Cu307.x/SrTi03 films
deposited on (1102) A1203. SrTi03 films deposited at an oxygen
pressure of 200 mTorr.

142
during the deposition resulted in a highly (110) oriented SrTi03 film (figure 4-
35). YBa2Cu307.x films subsequently deposited on these barrier layers adopted
orientations which minimized lattice misfit at the YBajQ^O^/barrier layer
interface. YBa2Cu307_x films on (100) SrTi03 barrier layers were (001) oriented,
while (103) oriented YBa2Cu307.x films grew on (110) SrTi03 barrier layers. The
electrical properties of these films were sensitive to the YBajQ^O^ orientations
(figure 4-36). The normal state properties of an (001) YBa2Cu307_x /(100) SrTi03
film deposited on (1102) A1203 were similar to YBa2Cu307.x deposited directly on
(100) SrTi03, with a metallic normal state resistance, but a lower zero resistance
temperature (83 K). The Jc value for this film was 2.5xl06 amps/cm2 at 4.5 K.
SEM micrographs of the (001) YBa2Cu307_x/(100) SrTi03 film on (1102) A1203
show a large grained microstructure with clearly visible grain boundaries (figure 4-
37). By contrast, the surface of a YBa2Cu307.x film deposited on (100) SrTiOs
was featureless.
Millimeter-wave transmission data taken at 36 gigahertz shows that the
surface resistance of the YBa2Cu307.x/SrTi03 film deposited on (1102) A1203 was
approximately 10 milliohms at 4.2 K (figure 4-38), which is lower than the value
reported for a YBa2Cu307.x film grown on single-crystal LaA103 (50 milliohms).
Figure 4-39 shows a sharp drop in the phase angle as the film went from the
normal to the superconducting state. The imaginary part of the conductivity,

INTENSITY (arbitrary units)
143
YBa2Cu307_x/SrTi03
on AI2O3(lT02)
CO
T-
o
A
A
V.
SrTi03 on AI2O3
(1T02). Deposited
at 40 mTorr 02.
A.
J V
A

smo3

AI2O3
10
CM
.O
O
O
CM
1
O'
.O
I CM
CM
4-A-
20
30
40
50
DIFFRACTION ANGLE (20)
60
Figure 4-35. X-ray diffraction patterns for SrTi03 and YBa2Cu307_x/SrTi03 films
deposited on (1102) A1203. SrTi03 films deposited at an oxygen
pressure of 40 mTorr.

RESISTANCE (OHMS)
144
10
8
6
4
2
0
0
Figure 4-36.
YBa2Cua07_x/SrTi03 ON (1102) Al20g.
SrTi03 BARRIER LAYERS DEPOSITED AT
50 100 150 200 250 300
TEMPERATURE (K)
Resistance versus temperature data for YBa2Cu307.x/SrTi03 films
deposited on (1102) A1203. The orientations of the YBajCujO^
films were determined by the SrTi03, which, in turn, were
determined by the oxygen pressure during the deposition.
Resistance values for the (103) oriented YBa2Cu307.x film are
multiplied by 0.50.

145
Figure 4-37.
0 0 %
4
IR *
20KU X8600 |000 1.0U UFMSE
Scanning electron micrograph of a (001) YBa2Cu307x/(100) SrTi03
film on (1102) A1203.

Rs (OHMS)
146
Figure 4-38. Surface resistance of a YBa2Cu307.x/SrTi03 film deposited on
(1102) A1203. Data taken at 36 gigahertz.

147
Figure 4-39. Phase angle versus temperature for a YBa2Cu307_K/SrTi03 film on
(1102) A1203. Data taken at 36 gigahertz.

148
which was due to the superconducting charge carriers, was greater than the real
part of conductivity at temperatures below 80 K (figure 4-40).
Lowering the oxygen pressure during growth of the SrTi03 barrier layer
resulted in a damaged YBajCujO^ film. The (103) oriented YBa2Cu3O7.x/(110)
SrTi03 film on (1102) A1203 was only slightly metallic, and did not reach T0 until
77 K. The critical current density for this film was 5x10s amps/cm2 at 4.2 K.
The electrical properties of YBa2Cu307.x/SrTi03 films deposited on the
(1210) and (0001) planes of A1203 were severely degraded. Figure 4-41 shows the
YBa2Cu307.x/SrTi03 film on (0001) A1203 had high normal state resistance values
which were almost independent of temperature, and there was a broad
temperature range (88 51 K) over which the resistance dropped from 36 to 0
ohms, with a long tail near 0 ohms. Similarily, the normal state resistances of
YBa2Cu307.x/SrTi03 on (1210) A1203 were much higher than the resistance values
of an ideal YBa2Cu307.x film, and the temperature dependence was
semiconducting, with a broad transition to T0 = 52 K. X-ray diffraction data
shows that SrTi03 grew with a strong (111) orientation on (0001) A1203, and the
YBa2Cu307.x film deposited on this structure was (001) oriented (figure 4-42).
SrTi03 grew with a weak (110) orientation on (1210) A1203, and the YBa2Cu307.x
film was weakly (103) oriented (figure 4-43).
Raman spectra from YBa2Cu307.x/SrTi03 films show the 335 cm'1 peak is
higher than the 500 cm'1 peak for the film on (1102) A1203, but the 500 cm'1 peak
is larger than the 335 cm'1 peak for the film on (1210) A1203 (figure 4-44).

COMPLEX CONDUCTIVITY (S/m)
149
Figure 4-40. Real (ctj) and imaginary (a2) parts of the electrical conductivity for a
YBa2Cu307.x/SrTi03 film on (1102) A1203. Data taken at 36
gigahertz.

RESISTANCE (OHMS)
150
50
40
30
20
10
0
i i i i i
o YBa2Cua07_x/SrTi0g on (1210) Al20a
YBa2Cu307_x/SrTi0g on (0001) A1203
0 50 100 150 200 250 300
TEMPERATURE (K)
Figure 4-41. Resistance versus temperature data for YBa2Cu307.x/SrTi03 films
on: () (0001) A1203; (O) (1210) A1203.

INTENSITY (ARBITRARY UNITS)
151
Figure 4-42. X-ray diffraction patterns from SrTi03 and YBa2Cu307.x/SrTi03
films on (0001) A1203.

INTENSITY (ARBITRARY UNITS)
Figure 4-43. X-ray diffraction patterns from SrTi03 and YBa2Cu307_x/SrTi0
films on (1210) A1203.

INTENSITY (arbitrary units)
Figure 4-44. Raman spectra of YBa2Cu307_x/SrTi03 films on (1210) and (102) A1203 substrates.
LA
LO

154
YBa3Cu307j,/LaAlQ3 Films on Si and A13Q3 Substrates
The optimal normal state and superconducting properties of YBajQ^O^*
films deposited on (1102) LaA103 substrates have been described earlier. LaA103
was considerably less effective as a barrier layer material than as a substrate
material, and YBa2Cu307.x/LaA103 films deposited Si were severely degraded.
The normal state resistance for YBa2Cu307_x/LaA103 on Si was high and
semiconducting, and there was a broad transition to the superconducting state,
from Tonset = 90 K to T0 less than 46 K (figure 4-45). AES data do not show
La-Si interdiffusion or formation of an interfacial phase at the substrate/barrier
layer interface (figure 4-46), and there are no well defined impurity phases at the
YBa2Cu307.x/barrier layer interface. Diffraction peaks from the rare-earth silicide
or any other impurity phases were not observed in the x-ray pattern (figure 4-47).
The YBa2Cu307_x film was weakly (001) oriented, and a smaller (103) peak was
also observed. The film surface was very rough (figure 4-48), and an
interconnected crack network was observed throughout the film.
LaA103 was also a poor barrier layer material for YBajQ^O^ films deposited
on the (102) or (1210) faces of A1203. The normal state resistance values of the
YBa2Cu307.x/LaA103 film on (1102) A1203 were moderately metallic, with a T0 of
58 K (figure 4-49). The resistance values of the YBa2Cu307.x/LaA103 film

RESISTANCE (OHMS)
155
Figure 4-45. Resistance versus temperature for a YBa2Cu307.x/LaA103 film on
(100) Si.

Figure 4-46. Auger line scan across crater edge of a YBa2Cu307.x/LaA103 film on (100) Si.
CT\

Figure 4-47. X-ray diffraction pattern from a YBa^O^/LaAlOa film on (100)
INTENSITY (ARBITRARY UNITS)
ro
o
O
T1
n
33
>
O
H
O u
O
r~
m
13
3 *
o
Ol

(002)
(003)
(005)
-<
CD
O
F
r

>
cP
O
3
W
8
(006)
-j

158
. +' * V** 'T
*

-f/ki* > J*;
# * .
< > *, 0 m
*'%.* * * :m
V- *Z ^ ^

V* ^#V.V t
#
v

/**
15KV X6000
8335 1.0T UFMSE
Figure 4-48. Scanning electron micrograph of a YBa2Cu307.x/LaA103 film on
(100) Si.

RESISTANCE (OHMS)
159
Figure 4-49. Resistance versus temperature data for YBa2Cu307.x/LaA103 films
deposited on: () (1102) A1203; (O) (1210) A1203.

160
on (1210) A1203 was semiconducting, and there was no transition to the
superconducting state. X-ray diffraction peaks from LaA103 barrier layers
deposited on both A1203 substrates were amorphous (figures 4-50 and 4-51), and
the only diffraction peaks observed were the Ka and Kfi peaks from the substrate.
An explanation is presented in the discussion section chapter as to why LaA103
films grown on (1102) A1203 appeared to be amorphous, whereas films such as
SrTi03, which has a crystal structure similar to LaA103, is highly oriented.
Despite the lack of x-ray diffraction from the LaA103 layers, the orientation of
the YBa2Cu307.x films was dependent on the orientation of the A1203 substrate.
In all cases, the intensities of x-ray diffraction signals from YI^CujO^ films
deposited on LaA103 barrier layers were much smaller than those obtained from
YBa2Cu307.x films deposited on single crystal LaA103 substrates. Raman spectra
of the YBa2Cu307.x/LaA103 films deposited on (102) A1203 and Si substrates
(figure 4-52) show large peaks at 500 cm'1, and shoulder peaks at 440 cm'1, which
indicate a large fraction of non-(001) oriented grains.
YBa^CugO-^/YAlOj Films on Si and A13Q3
YA103 barrier layers effectively limited interdiffusion between YBa2Cu307.x
films and Si or (1102) A1203 substrate. Figure 4-53 compares the resistance vs.
temperature performance of YBa2Cu307.x/YA103 films on Si and (1102) A1203
substrates. The YBa2Cu307.x/YA103 film on (1102) A1203 showed metallic
normal state behavior, with a sharp drop to T0 = 82 K. To compare the

INTENSITY (ARBITRARY UNITS)
161
Figure 4-50. X-ray diffraction patterns from LaA103 and YBa2Cu307.x/LaA103
films on (1102) A1203.

INTENSITY (ARBITRARY UNITS)
162
YBa2Cu307.x/LaAI03 _
on AI2O3(1210) i
LaAI03 on
AI2O3(1210)
O
ic3
/
YBa2Cu307x
LaAI03
ai2o3
+ Al203kp
/
20 30 40 50
DIFFRACTION ANGLE (20)
60
Figure 4-51. X-ray diffraction patterns from LaA103 and YBa,Cu,07_v/LaA10,
films on (1210) A1203.

INTENSITY (arbitrary units)
Figure 4-52. Raman spectra from YBa2Cu307.x/LaAI03 films on (100) Si and (102) A1203.

RESISTANCE (OHMS)
164
Figure 4-53. Resistance versus temperature data for YBa2Cu307.x/YA103 on: (O)
(100) Si; () (1102) A1203.

165
YBajCi^Oy.x/YAlC^ films deposited on Si and (1102) A1203, the resistance was
normalized by multiplying the resistance values of the YBa2Cu307.x/YA103 film
on Si by 0.125. The resistance versus temperature curve of this film was slightly
less metallic than for YBa2Cu307.x /YA103 on (1102) A1203, and had a lower
transition temperature (T0 = 78 K). In this study, the YBa2Cu307.x/YA103 film
on Si had the highest T0 of all the films deposited on Si substrates. SEM data
(figure 4-54) show that the YBa2Cu307_x/YA103 film on Si is composed of
interconnected submicron grains, with much fewer cracks as compared to samples
of YBa2Cu307.x/Y-Zr02 on Si. A cross-sectional view (figure 4-55) shows the
film/substrate interface is sharp, but there is not a well defined interface between
the YBa2Cu307.x and YA103 layers. AES data of YA103 deposited onto Si (figure
4-56) show segregation of A1 towards Si and formation of an interfacial Al-Si-O
phase. X-ray diffraction (figure 4-57), however, does not indicate any Al-Si-O
phase. AES analysis of the YBa2Cu307.x/YA103 film on Si (figure 4-58) show
very little interdiffusion between YBa2Cu307.x, YA103, or Si, indicating YA103 was
an effective barrier to chemical diffusion. SEM of YA103 films deposited on Si at
730 C (figure 4-59) show the films are featureless, hence the morphology
observed in the YBa2Cu307.x/YA103 film must have emerged during or after the
YBa2Cu307.x deposition.
YA103 was also an effective diffusion barrier between YBa2Cu307.x and
(1102) A1203. The only x-ray diffraction peaks from YA103 deposited on (1102)
A1203 are from the substrate (figure 4-60), indicating the film was amorphous.

166
Figure 4-54. Scanning electron micrograph of a YBa2Cu307_x/YA103 film on
(100) Si.

167
Figure 4-55. Cross-sectional view of YBa2Cu307.x/YA103 film on (100) Si, taken
with SEM.

Figure 4-56. Auger line scan across crater edge of a YA103 film on (100) Si.

Figure 4-57. X-ray diffraction pattern of a YBa2Cu307_x/YA103 film on (100) Si.
INTENSITY (ARBITRARY UNITS)

Figure 4-58. Auger line scan across crater edge of a YBa2Cu307.i[/YA103 film on (100) Si.
o

Figure 4-59. Scanning electron micrograph of a YA103 film on (100) Si.

INTENSITY (ARBITRARY UNITS)
172
Figure 4-60. X-ray diffraction pattern of a YBa2Cu307.x/YA103 film on (1102)
ai2o3.

YBajQ^CVx films deposited on this structure were highly (001) oriented. The Jc
was lxlO4 amps/cm2 at 4.2 K.
YBa3Cu3Q7l[/Y3Q3 Films on Si. Y-ZrQ3. and SrTiQ3 Substrates
Y203 barrier layers dramatically improved the Jc values of YBajQ^O^ films
deposited on Y-Zr02 substrates. The normal state resistance of the YBa2Cu307_x/
Y203 film on Y-Zr02 was metallic, with a T0 of 89.6 K (figure 4-61). The Jc
versus H data was similar to that of the YBa2Cu307.x/Y-Zr02 film on Y-Zr02,
with zero field Jc values of 7x10s A/cm2 at 77 K, and lxlO7 A/cm2 at 4.5 K.
X-ray diffraction of the YBa2Cu307_x/Y203 film on randomly oriented Y-ZrOz
show small peaks from the (001) YBa2Cu307_x phase (figure 4-62). The peak
intensities from this film were much smaller than the peak intensities from the
YBa2Cu307.x/Y-Zr02 film deposited on Y-Zr02. SEM shows this film is fine
grained, with numerous pinholes (figure 4-63).
The normal state resistance and Jc data for a YBa2Cu307.x/Y203 film
deposited on (100) SrTi03 (figures 4-64) were very similar to those observed on
Y-Zr02 substrates. The Jc values of the YBa2Cu307.x/Y203 film on (100) SrTi03
were 7x10s A/cm2 at 77 K, and lxlO7 A/cm2 at 4.5 K. X-ray diffraction
patterns from this film show the Y203 barrier layer was highly (100) oriented, and
only YBa2Cu307.x peaks from the (100) orientation were detected (figure 4-65).
SEM micrographs show the YBa2Cu307_x/Y203 film is smooth (figure 4-66).

RESISTANCE (OHMS)
174
Figure 4-61. Resistance versus temperature for a YBajQ^O^/YjC^ film
on Y-Zr02.

INTENSITY (ARBITRARY UNITS)
175
Figure 4-62. X-ray diffraction pattern from a YBa2Cu307.x/Y203 film on Y-Zr02.

176
m > *
A
*.4|
.\ ^
4 *
20KU &600 5300 1.0UUFMSE
Figure 4-63. Scanning electron micrograph of a YBajQ^O? JY203 film on
Y-Zr02.

RESISTANCE (OHMS)
177
100 150 200
TEMPERATURE (K)
250 300
Figure 4-64. Resistance versus temperature data for a YBa2Cu307.x/Y203 film on
(100) SrTi03.

Figure 4-65. X-ray diffraction pattern from a YBa2Cu307.I/Y303 film on (100)
INTENSITY (ARBITRARY UNITS)

179

%

%
20KU XI 1000 0070
1.0U UFMSE
Figure 4-66. Scanning electron micrograph of a YI^Q^O^/YjOj film on (100)
SrTi03.

180
Comparing Jc data from YBa2Cu307x films deposited on Y-Zr02 substrates using
Zr02, Y-Zr02, or Y203 barrier layers, we conclude that Y203 concentrations
greater than 13% in the barrier layers are crucial to the formation of YB^CujO^
films with high Jc values.
YBa2Cu307.x/Y203 films on Si had very poor electrical properties, with a
semiconducting normal state resistance and a long transition from Tonset = 90 K
to T0 less than 20 K (figure 4-67). SEM micrographs (figure 4-68) show an
interconnected crack network through the YBa2Cu307.x/Y203 film on Si, and the
film surface was very rough. AES data (figure 4-69) show that Ba has diffused
from the film to the substrate. Y203 films deposited on (100) Si at 730 C (figure
4-70) are smooth and featureless, which indicates the microstructural changes in
figure 4-68 occur during the YBajQ^O^ deposition and oxygenation.
YBa7Cu?Q7 x/(YA10?1LaAlQ3. or Y7Q,)/AbSi7On_Films_Qn i
-an3 Substrates
A series of experiments were performed in which 800 thick Al6Si2013 barrier
layers were deposited on Si or (1102) A1203 substrates prior to growth of the
second barrier layer (LaA103, YA103, or Y203) and the YBa2Cu307_x film. For Si
substrates, the role of the Al6Si2013 was to create a layer with a slightly larger
thermal expansion coefficient than Si. On (1102) A1203, the lower thermal
expansion coefficient of Al6Si2013 (relative to A1203) altered the stress in the
second barrier layer and YBa2Cu307.x films. Because the electrical resistance of
the YBa2Cu307.x films were very sensitive to the extent of cracking within the

RESISTANCE (OHMS)
181
Figure 4-67. Resistance versus temperature data for a YBa2Cu307_x/Y203 film on
(100) Si.

182

Figure 4-69. Auger line scan across crater edge of a YBa2Cu307.x/Y203 film on (100) Si.

Figure 4-70. Scanning electron micrograph of a Y203 film on (100) Si.

185
films, a more complete description of the mechanisms by which thermally induced
stresses were transferred through the barrier layers was obtained by comparing
the resistance values of YBa2Cu307.x/barrier layer films grown on (1102) A1203
with or without an Al6Si2013 layer. A complicating factor was that Al6Si2013 was a
source of Si, which favors formation of the tetragonal phase when it diffuses into
YBa2Cu3O7.jp and thereby reduces the transition temperature. Previous data
showed the presence of tetragonal YBa2Cu307_x will lower T0 (see figure 4-1), but
only increase the normal state resistance values slightly. On the other hand, the
resistance values of cracked YBa2Cu307.x films are about a factor of 10 higher
than those of uncracked films. This study was useful for distinguishing whether
YBa2Cu307.x degradation was caused by tensile cracking, or Si diffusion into the
YBa2Cu307.x film.
Addition of Al6Si2013 into the barrier layer structures significantly altered the
electrical performance and microstructures of YBa2Cu307.x films deposited on Si
substrates. Resistance versus temperature data (figure 4-71) show the normal
state resistances for both the YBa2Cu307_x/YA103/Al6Si2013 and YBa2Cu307.x/
LaA103/Al6Si2013 films on Si are metallic, but the resistance values for the film
with a LaA103 layer are less than one-half of those observed in the film with a
YA103 layer. The transition temperatures were very similar, with T0 = 61.5 and
59 K for the samples with a YA103 or LaA103 barrier layer, respectively. SEM
data (figures 4-72 and 4-73) show the morphologies of both films are virtually
identical, with unconnected crack networks running through the films. The x-ray

RESISTANCE (OHMS)
186
120
100
80
60
40
20
0
0 50 100 150 200 250 300
TEMPERATURE (K)
T
o YBa2Cu307_x/YA10g/Al6Si2013 ON Si
YBa2CUg07_3t/LaA10g/Al6Si2013 ON Si
Figure 4-71. Resistance versus temperature data for: (O) YBa2Cu307.x/
YA103/Al6Si2013 film on (100) Si; () YBa^CW
LaA103/Al6Si2013 film on (100) Si.

187
2@tu X8600 0004 1.0U UFMSE
Figure 4-72. Scanning electron micrograph of a YE^Q^O^/YA103/Al6Si2013
film on (100) Si.

188


*
>

fe



1%

%
1
it %
m
w
20KU XI1000 0004
*
1.0U
0
A
"UFMSE
Figure 4-73. Scanning electron micrograph of a YBa2Cu307_x/LaA103/Al6Si2013
film on (100) Si.

189
diffraction patterns for both of these films show strongly (001) oriented
YBa2Cu307.x peaks (figure 4-74). The T0 for the YBa2Cu307_x/YA103/Al6Si2013
film on Si film is significantly lower than the YBa2Cu307_x/YA103 film on Si (61.5
versus 78 K), whereas T0 is higher for the YBa2Cu307.x/LaA103/Al6Si2013 film
on Si than for the YBa2Cu307_x/LaA103 film on Si (59 versus 46 K). In addition,
the normal state resistance values in the film containing the LaA103/Al6Si2013
barrier layer were a factor of 10 lower than in the film with only a LaA103 barrier
layer. Although the electrical and x-ray diffraction data indicate similar
microstructures for the YBa2Cu307_x/(LaA103 or YA103)/Al6Si2013 films
deposited on Si, the Raman spectra for these films are quite different (figure
4-75). The 500 cm"1 peak of the film with a LaA103 layer is smaller than the 335
cm"1 peak, which indicates a high degree of (001) orientation. By contrast, the
500 cm"1 peak in the film with a YA103 barrier layer is more pronounced.
Comparing these data to the Raman spectra obtained from the YBa2Cu307.x
/YA103 film on Si, we see that addition of the Al6Si2013 layer resulted in more
(100) and (010) oriented YBajC^O^ grains.
The effectiveness of YA103/Al6Si2013 relative to Y203/Al6Si2013 as a barrier
layer structure to (1102) A1203 presents a unique comparison, since both are
comprised of the same elements, but form different phases. The resistance vs.
temperature data (figure 4-76) show the normal state behavior of theYBa2Cu307.x
/YAlcyAlgSiAa film on (1102) A1203 was metallic, with a T0 = 75 K. This
was the highest transition temperature observed for YBa20u3O7.x films grown on

190
barrier layer structures which contained an Al6Si2013 layer. Conversely, the
normal state behavior of the YBa2Cu307.x/Y203/Al6Si2013 film on (1102) A1203
was semiconducting, and T0 was less than 39 K. X-ray diffraction data (figure 4-
77) show that the YBa2Cu307.x layer was primarily (001) oriented for both
structures, with minor (100) contributions. The predominant barrier layer peaks
in the film with a Y203 barrier layer were from the (222) Y203 phase; there were
no YA103 peaks in the YBa2Cu307.x/YA103/Al6Si2013 film on (1102) A1203. The
Raman spectra for the two films are nearly identical, with major peaks at 335
cm"1, and smaller peaks at 500 cm"1 (figure 4-78).

INTENSITY (ARBITRARY UNITS)
191
Figure 4-74. X-ray diffraction patterns from YBa2Qi307_x/YA103/Al6Si2013 and
YBa2Cu307.x/LaA103/Al6Si2013 films on (100) Si.

INTENSITY (arbitrary units)
Figure 4-75. Raman spectra of YBa2Cu307.x/YA103/A]6Si2013 and YBa2Cu307.x/LaA103/Al6S2013 films on (100) Si.

RESISTANCE (OHMS)
193
60
50
40
30
20
10
0
Figure 4-76.
Resistance versus temperature data for: (O) YB2^Ca307.J
YAlOj/A^SiA, on (1102) A1203; () YBa^O^/YA/AljSijO*
on (1102) A1203.

INTENSITY (ARBITRARY UNITS)
YBa2Cu307.x/Y203/AI6Si2013 on Al203 (1102)
Figure 4-77. X-ray diffraction patterns for YBajO^O-^/YAlCVAl^O^ and
YBa2Cu307.x/Y203/Al6S2013 films on (1102) A1203.

INTENSITY (arbitrary units)
YBa2Cu307.x /Y203/AlgSl20-j3
on (1102) Al203
250 300 350 400 450 500 550 600
RAMAN SHIFT (cm*1)
Figure 4-78. Raman spectra for YBa2Cu307_;x/YA103/Al6Si2C)13 and YBa2Cu307.J/Y203/Al6Si2013 on (1102) Al;

CHAPTER 5
DISCUSSION
A number of barrier layer/substrate combinations were examined in order to
what the most critical parameters were for optimal quality YBa2Cu307_x films on
the various substrates. These barrier layer/substrate combinations, and the zero
resistance temperatures and critical current densities observed in the
superconducting films are summarized in table 5-1.
Intergranular Versus Intragranular Effects
YBa2Cu307.x films with nearly ideal normal state and superconducting
properties were deposited on (1102) LaA103 substrates. Figure 4-1 shows that
the normal state resistance behavior for the fully oxygenated film was metallic,
with a sharp drop to 0 ohms at 88 K. Using the Halbritter equation (2-34) to
separate the resistance values into inter and intragranular contributions, we see
that the intergranular resistance was negligible compared to the intragranular
contributions, since the R vs. T curve extrapolates to 0 ohms at 0 K. The film
which was cooled to 550 C before filling the chamber with oxygen does not show
ideal R vs. T behavior, and the temperature at which the resistance became zero
was depressed to 79 K. This degradation was attributed to an incomplete
transition from the tetragonal to the orthorhombic crystal structure during cooling,
196

197
Table 5-1. Superconducting transition temperatures (zero resistance) and critical
current densities for YBa2Cu307_x/barrier layer films on different
substrates. Dashed lines (---) indicate no data were taken.
Substrate
Barrier Layer
T0 (K)
Jc (amps/cm2)
(1102) LaA103
none
88
5X107
off-axis Y-Zr02
none
86
6.8X103
(100) SrTi03
none
89

(100) Si
Yo.i3Zr0870194
76

off-axis Y-ZrOz
Yo.i3Zr0i8701>94
87
1.5 X107
(1102) LaA103
Yo.i3Zr0g70194
88
3X107
(1102) A1203
Yo.i3Zr0870194
83
1.9 xlO4
(1210) A1203
^o.i3^r0870194
80

(0001) ai2o3
^0.13^r0.87^1.94
77

off-axis Y-Zr02
Zr02
89
9X104
(100) SrTi03
Zr02
88
3.6 xlO6
(1102) A1203
(100) SrTi03
83
2.5 xlO6
(1102) A1203
(110) SrTi03
77
5X103
(1210) A1203
SrTi03
52

(0001) A1203
SrTi03
51


Table 5-1. Continued.
Substrate
Barrier Layer
T0 (K)
Jc (amps/cm2)
(100) Si
LaA103
< 46

(1102) A1203
LaA103
58

(1210) A1203
LaA103
< 30

(100) Si
yaio3
78

(1102) A1203
yaio3
82
lxlO4
off-axis Y-Zr02
y2o3
90
lxlO7
(100) SrTi03
y2o3
88
1.8 X107
(100) Si
y2o3
< 20

(100) Si
Y A103/Al6Si2013
61.5

(100) Si
LaA103/ Al6Si2013
59

(1102) A1203
YA103/Al6Si2013
75

(1102) A1203
Y203/AlfjSi2013
< 39


199
which resulted from an inadequate supply of oxygen as the film was cooled below
the tetragonal-to-orthorhombic transition temperature (700 C). The tetragonal
regions would be normal junctions through which the superconducting charge
carriers had to pass, hence the amplitude of the superconducting wave function
was diminished and T0 lowered. The Halbritter equation was used to determine
the microstructural features most responsible for the normal state electrical
performance of the films. Since the resistance versus temperature curve does not
extrapolate to 0 ohms at 0 K for the 550 C film, this equation suggests
intergranular defects, such as grain boundaries were impeding the superconducting
charge carriers. Because the slope of the resistance versus temperature curve of
this film is larger than the 730 C film, the Halbritter equation suggests that
either intragranular defects were present, or the percolation pathway was
increased. Since the magnitude of the normal state resistance values for the 550
C film are similar to those of the 730 C film (figure 4-1), it is unlikely that the
percolation path lengths are significantly different, and we conclude that the film
was degraded by both intragranular and intergranular defects; intragranular
defects because the slope of the curve is higher than the ideal film, and
intergranular because the curve extrapolated to higher than 0 ohms at 0 K.
Scanning electron micrographs of the fully oxygenated films are virtually
featureless, except for the particulates which resulted from the laser deposition
process. This reinforces the conclusion that increased percolation distance was
not the cause of increased resistance in the oxygen depleted film. X-ray

200
diffraction showed that the fully oxygenated YBa2Cu307.x film was primarily (001)
oriented, with small (100) peaks also present. Although the extent of in-plane
epitaxy cannot be determined from the 0-20 x-ray diffraction scans used in this
study, the high critical current density for this film (5xl07 amps/cm2 at 4.2 K) is
strong evidence that the YBa2Cu307.x grains were strongly coupled.
Comparisons between YBa2Cu307.x films deposited directly onto single-crystal
Y-Zr02 substrates with YBa2Cu3O7_x/Y-Zr02 films on Si show that the films
deposited on Si are cracked, whereas the YBa2Cu307.x films on the Y-Zr02
substrate were not cracked. SEM micrographs of Y-Zr02 films deposited on Si
do not show cracks, thus cracking occurs during YBa2Cu307.x deposition or the
cooling cycle. One approach to understand why cracking occurs in YBa2Cu307_x/
Y-Zr02 films deposited on Si is to look for microstructural differences between
these films and YBa2Cu307.x films deposited directly onto Y-ZrOz substrates.
Auger analysis shows the Y-Zr02 layer effectively prevented interdiffusion
between YBa2Cu307.x and Si, so degradation via Si interdiffusion into the
YBa2Cu307.x film was probably not the cause of cracking. X-ray diffraction
patterns from YBa2Cu307.x/Y-Zr02 films on Si versus YBa2Cu307.x films on Y-
Zr02 substrates show both are predominately (001) oriented, but the YBa2Cu307.x
film on Y-Zr02 also had minor (100) YBa2Cu307.x peaks. Based on the Auger
and x-ray data, the microstructures of the two films seem to be very similar.
Interpretation of the normal state resistance values of the two films in terms of
the Halbritter equation shows that after multiplying the resistance values of the

201
YBa2Cu307.x/Y-Zr02 film on Si by 0.061, the two curves are very similar. This
indicates that both films have about the same intragranular and grain boundary
resistance, but the percolation parameter (p) was much larger for the film
deposited on Si. This is consistent with the observation of cracks in this film,
which increase the length of current pathways and thus increase the normal state
resistances. The degradation of T0 in this film was attributed to poor coupling of
the superconducting charge carriers across the cracked regions. Even though
there appear to be pathways by which the superconducting holes could travel,
microcracking and dislocations not imaged by the SEM will reduce the
superconducting order parameter in these regions. The resistance versus
temperature curve for YI^CujO-^ deposited on single crystal Y-Zr02 suggests
the grain boundary resistance was at least as high as the intragranular resistance,
and the low Jc (6.8 xlO3 A/cm2 at 4.5 K) was probably due to high angle grain
boundaries, which resulted in poor intergranular coupling of the superconducting
holes at the grain boundaries.
Effect of Y-ZrQ3.Y3Q3. and ZrO-, Barrier Layers on Jc Values
Figure 5-1 shows Jc versus magnetic field intensity for YBa2Cu307.x/
Y0.i3Zr087O194 films deposited on Y-Zr02 and LaA103 substrates. The Jc values
are roughly proportional to If0-5, indicating the Jc was limited by flux creep within
the grains.34 Similar plots for YBa2Cu307.x/Y203 films deposited on Y-Zr02 and

202
SrTi03 substrates shows the same H-0-5 dependence at 4.5 and 77 K (figures 5-2
and 5-3).
The Jc versus H behavior of YBa2Cu307_x/Zr02 films deposited on off-axis Y-
Zr02 and SrTi03 substrates are shown in figures 5-4 and 5-5. The shape of the
curves are similar to those reported by others for bulk and thin film samples with
Josephson weak links.123,124 The shape of Jc versus Ha curves for the YBa2Cu307_x
/Zr02 films, and lower Jc values (relative to the films with Y203 or Y-Zr02
barrier layers) indicates the Jc was limited by decoupling of Josephson weak links
at the grain boundaries. Since Zr02 is monoclinic, there could not be epitaxial
lattice matching at the YBa2Cu307.x/Zr02 interface. The Jc values of the
YBa2Cu307.x/Zr02 films deposited on (100) SrTi03 are close to the highest values
predicted for films with aligned (but not epitaxial) <100> and <010>
YBa2Cu307.x directions (2xl06 amps/cm2 at 4.5 K),125 and the Jc of the
YBa2Cu307.x/Zr02 film on Y-ZrOz is close to the maximum value predicted for
(001) films with random in-plane orientations (3 x10s amps/cm2 at 4.5 K).92 The
Jc and resistance versus temperature data indicate high-angle grain boundaries
were the primary reason for lower Jc values.
Presumably, the high Jc values observed in the YBa2Cu307.x/Y203 film on
SrTi03 could be attributed to epitaxial growth of both the Y203 and YBa2Cu307.x
films, which would limit the fraction of high-angle grain boundaries. Whether
highly (100) oriented Y203 was also deposited on Y-Zr02 substrates is not clear,
since the substrates were cut off-axis from the (100) planes, and diffraction peaks

203
from the Y203 barrier layer or Y-Zr02 substrate were not observed. The Y-Zr02
film deposited on (100) LaA103 was primarily (100) oriented, and a smaller (220)
peak was also observed. The presence of both peaks usually creates a
YBa2Cu307.x films with mixed in-plane orientations, and high-angle grain
boundaries. The high Jc values of this film indicate good coupling of the
superconducting holes across the grain boundaries.
MAGNETIC FIELD (T)
Figure 5-1. Jc versus magnetic field intensity for YBa2Cu307_x/Y-Zr02 films
deposited on Y-Zr02 and (1102) LaA103 substrates. Data taken at
. 4.5 K. The solid and dashed lines are fits of the experimental data
with the flux creep model

204
MAGNETIC FIELD (T)
Figure 5-2. Jc versus magnetic field intensity for YBa2Cu307.x/Y203 films on Y-
Zr02 and SrTi03 substrates. Data taken at 77 K. The solid line is
a fit of the experimental data to the flux creep model.

CRITICAL CURRENT DENSITY (10 A/CM )
205
MAGNETIC FIELD (T)
Figure 5-3. Jc versus magnetic field intensity for YB2l2Oi301.JY203 films on Y-
Zr02 and (100) SrTi03 substrates. Data taken at 4.5 K. The solid
and dashed lines are fit of the experimental data to the flux creep
model.

CRITICAL CURRENT DENSITY (10 A/CM )
Figure 5-4. Jc versus magnetic field intensity for YBa2Cu307.x/Zr02 film on
Y-Zr02 substrate. Data taken at 4.5 K.

>H
Eh
<
m
S3
w
Q
Eh
S3
m

K
S3
u
l-J
o

Eh
P
o
0.0
YBa2Cua07_x/Zr02 ON
(100) SrTiOg
4.5 K
0.4 0.8 1.2 1.6
MAGNETIC FIELD (T)
2.0
Figure 5-5. Jc versus magnetic field intensity for a YBa2Cu307.x/Zr02 film
(100) SrTi03. Data taken at 4.5 K.

208
Despite similar resistance versus temperature behavior for all the films, there
were large variations in Jc. The T0 for each of the films was 87-90 K, and the
resistance versus temperature curves extrapolated close to 0 ohms at 0 K. This
shows that the normal state intergranular resistances were neglible compared to
the intragranular resistances. However, superconducting films deposited on Y203
or Y-Zr02 barrier layers had high Jc values, but the Jc values of films on Zr02
barrier layers were lower, and were limited by poor intergranular coupling of the
superconducting hole pairs. SEM micrographs show that films deposited with
Y203, Y-Zr02, or Zr02 barrier layers had different grain sizes and surface
roughnesses, but there was no correlation between the features observed by SEM
and the Jc values. We believe the YBa2Cu307.x films with Y203 or Y013Zr0 87O194
barrier layers grew with yittria-rich grain boundaries, which were flux pinning sites
at the grain boundaries, similar to the suggestion of Catana et al.126
Effective flux-pinning sites result from steep gradients in the superconducting
order parameter between superconducting and non-superconducting regions of the
film. The ability of the YBa2Cu307_x/Y203 interface to be a flux-pinning site can
be traced to two properties. First, Catana et al.126 showed that the YBa2Cu307.x
/Y203 interface was coherent, and the intersection of a growing YBa2Cu307.x
grain with a Y203 inclusion resulted in epitaxial growth of YBa2Cu307.x around
the inclusions, with formation of edge dislocations and other defects near the
Y203 inclusions. Second, Ying et al.127 reported that the mobility of oxygen
through Y203 is very high. This means that as the YBajQ^O^ films are cooled,

209
oxygen readily diffuses through the yittria-rich grain boundaries, into the
YBa2Cu307.x grains, thus insuring the entire YBa2Cu307.x grain is fully oxygenated.
The combination of coherent YBa2Cu307_x/Y203 interfaces along with the high
oxygen mobility through Y203 minimizes degradation and oxygen depletion of
YBa2Cu307.x near the grain boundaries as long as the oxygen pressure is
sufficiently high during cool-down.
Dependence of T0 and Jc on A13Q3 Substrate Orientation
Variations in the electrical properties of YBa2Cu307.x/Y-Zr02 films deposited
on different orientations of A1203 substrates were also observed. YBa2Cu307.x/
Y-Zr02 films deposited on (1102) and (1210) A1203 substrates displayed metallic
resistances, but the transition to the superconducting state was sharper (83 K)
for the film deposited on (1102) A1203. The transition for the film deposited on
(1210) A1203 was broader, with a T0 of 80 K. The normal state electrical
properties for the YBa2Cu307.x/Y-Zr02 film on (0001) A1203 were much worse
than those of the films deposited on the other A1203 orientations, with resistance
values ~ 5 times higher than the YBa2Cu307.x/Y-Zr02 films deposited on (1102)
and (1210) A1203. There was also a very broad (89 77 K) transition to the
superconducting state. X-ray diffraction showed that the Y-Zr02 film deposited
on (0001) A1203 had a higher ratio of (200)/(lll) orientation than Y-Zr02 films
deposited on (1102) or (1210) A1203, and the YBa2Cu307.x/Y-Zr02 films with the
highest Jc values are typically observed when the Y-ZrOz layer is highly (100)

210
oriented.93 In addition, Raman spectra from YBa2Cu307.x/Y-Zr02 films deposited
on (1102) and (0001) A1203 show the peak at 500 cm'1 is small, which suggests
both YBa2Cu307.x films were highly (001) oriented; the non-symmetric lineshape
of the 335 cm'1 peaks indicate the films were highly oxygenated. The explanation
for this apparent anomaly is given by the SEM micrograph of YBa2Cu307_x/Y-
Zr02 on (0001) A1203 (figure 4-32). This picture shows a faint series of cracks in
the film, and the degradation of the electrical and superconducting properties are
attributed to these cracks. Like the YBajQ^O^/Y-Zr02 films on Si, these cracks
form because of the difference in thermal expansion coefficients between
YBa2Cu307_x and the substrate. The thermal expansion coefficient of A1203 is
anisotropic (table 4-1), with the lowest values observed in directions parallel to
the (0001) plane, and the largest values being perpendicular to (0001). Thus we
conclude that for YBa2Cu307.x/Y-Zr02 films on A1203, the thermally induced
stresses are largest in YBajCujO^ films deposited on the (0001) plane. This
accounts for the cracking and degradation of electrical properties in these films.
In general, film cracking dramatically increased the normal state electrical
resistances of YBa2Cu307.x films. Since the electrical resistance across a crack is
not very sensitive to changes in the temperature, resistance versus temperature
profiles of YBa2Cu307_x films were less metallic than those of the uncracked films,
and were sometimes semiconducting. For cases in which cracking was the
dominant defect, any excess resistance caused by intragranular damage was
masked by the cracking.

211
Estimate of Stress and Cracking Due to Differential Thermal Expansion
A rough analysis of the magnitude of the thermally induced stresses can be
obtained from the continuum mechanics model. The difference in thermal
expansion coefficients necessary to cause film cracking can be obtained by
inserting values128 of a = 50 MPa, E = 86 GPa, v = 0.290, and AT = 705 C for
the fracture strength, Young's modulus, Poisson's ratio, and temperature
difference into equation 4-57:
50jc106 pa (86x 109 Pa) (705) (Aa)
0.71
This analysis suggests that if Aa exceeds 5.9xlO7/C, thermally induced stresses
will cause cracking in the YBajQ^O,.* film upon cooling down from 750 C.
However, there are several factors which affect the validity of this arguement.
First, the fracture strength used in the calculation is for bulk YBa2Cu307.x, and the
film fracture strength is probably much higher. Second, the assumption that all of
the stress is relieved by either elastic deformation or cracking must be incorrect,
since uncracked YBa2Cu307_x films can be deposited on substrates such as SrTi03
and LaA103, for which the difference in thermal expansion coefficients exceeds
the calculated critical values. Since the melting temperature of YBa2Cu307.x is
1010 C, it is probable that atomic motion is relatively easy near the deposition
temperature, so plastic deformation is expected to be the primary stress reduction

212
mechanism. At lower temperatures, plastic deformation becomes more difficult,
so the thermal stresses become more elastic. From this analysis, we conclude that
estimation of the thermal stresses via the continuum mechanics model is too high
because it neglects plastic deformation in the high temperature regime. The
temperature below which elastic deformation becomes appreciable is a function of
the bonding strength at the YBajQ^O^/substrate (or barrier layer) interface.
While the model cannot be taken too literally, it predicts that if thermally induced
elastic stresses exceed the fracture strength of the film, cracking will occur. This
was experimentally observed in YBa2Cu3Q7_x films deposited on Si, and to a lesser
extent on (0001) A1203.
The slightly broadened transition from Tonset (88 K) to T0 (80 K) in the
YBa2Cu307.x/Y-Zr02 film deposited on (1210) A1203 resulted from the
polycrystalline Y-Zr02 barrier layer. The lack of an x-ray diffraction pattern from
Y-ZrOz suggests the film was fine grained and randomly oriented. We conclude
that the YBa^Q^O^ film must have had random in-plane orientation, and T0 was
lowered by poor coupling across high-angle grain boundaries.
Effects of Texture and In-Plane Alignment
The best electrical properties for YBa2Cu307.x films deposited on A1203
substrates were obtained using SrTi03 barrier layers. YBa2Cu307.x/SrTi03 films
on (1102) A1203, in which both films were deposited at 200 mTorr oxygen,
showed a metallic resistance vs. temperature behavior which extrapolated to 0

213
ohms at 0 K, indicating that grain boundary resistance was negligible compared
to the intrinsic intragranular resistance of YT^Q^O^. The critical current
density for this film was 2.5xl06 amps/cm2 at 4.2 K, which was lower than the
value observed for YBa2Cu307.x films deposited on (1102) LaA103 substrates, but
over 100 times greater than the Jc value of the YBajQ^O^/Y-Zr02 film on
(1102) A1203. The high Jc value for this film indicates there was a large degree of
in-plane alignment between the YI^CujO^ and SrTi03 films, hence the number
of high-angle YBa2Cu307.x grain boundaries were minimal. Raman spectra shows
the 335 cm'1 peak was larger than the 500 cm'1 peak, which indicates the film was
highly (001) oriented, and the asymmetry of the 335 cm'1 peak shows the film was
highly oxygenated. Raman spectra shows the intragranular microstructure of the
film was nearly optimal. The reason the YBa2Cu307_x film grew with a high
degree of (001) texture was attributed to the predominately (100) growth of the
SrTi03 barrier layer. Since SrTi03 has the same perovskite crystal structure and
similar lattice parameters as YBa2Cu307.x (a<100> for SrTi03 = 3.90 , a<100> <010>
= 3.82 and 3.89 for YBa2Cu307.x), YBa2Cu307_x grows epitaxially on single
crystal SrTi03. The Jc for this film was not as high as the Jc observed in
YBa2Cu307.x films deposited on LaA103 substrates, possibly because of the small
amount of (110) oriented SrTi03 grains which could have led to high-angle grain
boundaries in the YBa2Cu307.x film and reduced Jc across these grains. SEM
micrographs show the surface to be composed of clearly defined grains, whereas

214
YBa2Cu307.x films deposited on LaA103 or SrTi03 substrates are featureless at
similar magnifications. This morphology may have caused the reduction in T0 and
Jc values for the YBa2Cu307.x/SrTi03 film on (1102) A1203.
Millimeter-Wave Properties
The millimeter-wave properties of YBa2Cu307.x/SrTi03 films were very good
(figures 4-38 and 4-39). The resistance at 36 GHz was 10'2 ohms, which was lower
than the value observed in a YBa2Cu307_x films deposited on LaA103 (5 X 10~2
ohms). There was a sharp change in the phase angle of the film as it went from
the normal to superconducting state. In the normal state, the phase angle was
close to 0, which is typical for normal conductors. At the transition temperature,
the phase angle of the current sharply dropped to 90 out of phase with the
voltage, which is expected for films in which the superconducting charge carriers
dominate the conductivity. There was no evidence that the SrTi03 barrier layer
affected the dielectric properties of the film or substrate, despite the fact that bulk
SrTi03 is lossy and has a dielectric constant of over 1500 at 77 K. Apparently,
film stresses inhibit the tetragonal-to-orthorhombic phase transition in SrTi03
films. Since the high dielectric constant is caused by the structural phase
transformation, inhibiting this transition prevents the dielectric constant from
increasing. Recent experiments in which the dielectric constant of a 4000
SrTi03 film was measured as a function of temperature show that the dielectric
constant is 300 from 4.2 to 300 K.129

215
Effects of Surface Energy
The good superconducting properties of YBa2Cu307.x/SrTi03 films on (1102)
A1203 were largely due to the ease with which (100) oriented SrTi03 films could
be deposited on (1102) A1203 substrates. Comparisons between the orientations
of SrTi03 films deposited on (1102), (0001), and (1210) A1203 indicates that
lattice mismatch, excess AHf for SrTi03, and phenomena intrinsic to the pulsed
laser deposition process synergistically affected the microstructures of the SrTiOa
films.
Reduction of interfacial free energy by epitaxial growth is a strong driving
force for highly textured film growth. Table 5-1 lists the smallest in-plane lattice
mismatches which can occur between the low index directions of SrTi03, LaA103,
(100) MgO, and (1102), (1210), or (0001) orientations of A1203. The minimum
lattice mismatch between SrTi03 and any cut of A1203 is 7.4%. By comparison,
the minimum lattice mismatch between LaA103 and A1203 is only 4.5%.
Experimentally, we observe that highly oriented SrTi03 can be grown on (1102)
and (0001) A1203, whereas LaA103 films deposited on (1102) A1203 are
amorphous. The cause of the behavior was attributed to the more negative excess
AHf for SrTi03. For the reaction

216
SrO + Ti02 = SrTi03
the system lowers its AHf ^ .c by 47.4 kilojoules/(gram atom) of oxygen by
adopting the SrTi03 structure, compared to the extrapolated AHf m .c which a
mixture of SrO + TiOz phases would have. Although the AHf for LaA103 has not
been measured, it is likely that the excess AHf is much smaller than that of
SrTi03, since the crystal structures of La^ and A1203 are both hexagonal, and
the structure of LaA103 is basically an alternating series of LaOx and A10x layers.
Hence we conclude that the excess AHf for SrTi03 was a prime driving force for
crystallization of the film, whereas the lower excess AHf for LaA103 caused the
film to remain amorphous.
Table 5-2. Lattice mismatch between SrTi03 or LaA103 films, and A1203 or MgO
substrates. Mismatches were calculated by using the value of ardm for
the direction indicated in the vertical column, and calculating the
minimum mismatch which could be achieved by matching afllm with
one of the low index directions of the substrate surface, and using the
lattice parameter of this direction for a,.ubstrate.
(1102) A1203
(1210) A1203
(0001) ai2o3
(100) MgO
<100> SrTi03
+ 11.5
-5.3
-18.1
-7.1
<110> SrTi03
+ 7.4
+15.1
+ 15.8
+ 31.2
<100> LaA103
+ 8.3
-8.0
-20.4
-9.8
<110> LaA103
+ 4.5
-17.5
+ 12.6
+ 27.6

217
Effects of Lattice Matching
Minimization of the lattice mismatch and matching of the surface meshes
strongly influenced the orientation of SrTi03 films deposited on A1203 substrates.
SrTi03 films deposited at 200 mTorr onto (1102) A1203 were predominately (100)
oriented, while films grown on (0001) A1203 under identical conditions were
strongly (111) oriented. Because the two dimensional mesh of the (1102) A1203
surface has an almost tetragonal mesh (the angle between the <100> and
<012> directions is 86), it is not suprising that SrTi03 adopts a cubic overlayer
mesh. Similarity, the oxygen atoms of the (0001) A1203 surface are hexagonalty
close-packed, so the SrTi03 overlayer adopts the (111) orientation, in which the
oxygen atoms are also hexagonalty close-packed. Note that the (111) SrTi03
orientation is predominant despite a lattice mismatch of 15.8% with the A1203
(0001) substrate. Matching the type of surface mesh appears to be as important
as the lattice mismatch in the growth of highly oriented SrTi03 films. Oriented
SrTi03 films were not observed on (1210) A1203, and this was attributed to the
lack of a close-packed surface mesh on the A1203 substrate. The smallest surface
mesh for the (1210) surface is a tetragonal mesh with lattice parameters of 12.99
x 8.24 (see figure 5-6).

218
SURFACE MESHES OF (0001), (1102) AND (1Z10)
ORIENTATIONS OF Al203
O = Oxygen
0 = Aluminum
(1210) A
1.299 nm
>
0.824 nm
*0.0 0.0-0 o
0.0 0.0.0
*0 o O- o om o
o o o o' o
Figure 5-6. Surface meshes of (0001), (1102), and (1210) orientations of A1203.

Effects of Oxygen Pressure on SrTi03 Growth
The oxygen pressure during growth of SrTi03 films on (1102) A1203 substrates
dramatically affected the orientation of the SrTi03 films. Films deposited at 200
mTorr 02 were highly (100) oriented, while films grown at 40 mTorr 02 were
(110) oriented. This behavior can be explained by correlating the surface
structures and lattice mismatches of the (100) and (110) SrTi03 films with the
deposition conditions which caused these orientations to flourish. Table 5-1 shows
that the minimum lattice mismatch occurs between the SrTiOa <110> and A1203
<012> directions. Starting with the premise that this relationship is the basis for
epitaxy between the SrTi03 film and (1102) A1203 substrate, it is equally
favorable for the film to adopt either the (100) or (110) orientations. The
principal difference between growth at the two different oxygen pressures was that
at 40 mTorr 02, the laser plume expanded from the target and encompassed the
substrate. Within the plume, the gas density is higher than the ambient, so the
flux of species from the target which reaches the substrate is greater than the
oxygen flux from the ambient.130,131,132 When the oxygen pressure was raised to
200 mTorr, the plume only extended approximately 3 cm, so it did not reach the
substrate, indicating the 02 gas density at the substrate was higher than the gas
density in the plasma. The (100) SrTi03 surface is comprised of alternating SrOx
and TiOx layers (figure 5-7). The O/Ti ratio in the first layer is 2, and the O/Sr

220
SURFACE STRUCTURE OF (100) SrTiOg
Q = Oxygen
0 = Titanium
^ = Strontium
Ti-q* LAYER
0.390 nm ; ^
>
0.390 nm
SrOxLAYER
0.390 nm
0.390 nm
o o o o o o
000*000
;o:o:o:o:o:o
o o o o o o
k-
Figure 5-7. Surface structure of (100) SrTi03.

221
Sr-Ti-0
LAYER
0.39 nm
OXYGEN
LAYER
0.39 nm
SURFACE STRUCTURE OF (110) SrTi03
O = Oxygen
= Titanium
= Strontium
ooooo
8*8888

> iO oo o o
0-550 nnn If
*o oooo
k >\
oo oo oo oo
oooooooo
^ lOO oo oo oo
0.550 nm
*0 0 oo oo oo
K H
Figure 5-8. Surface structure of (110) SrTi03.

222
ratio of the second layer is 1. We postulate that because the Sr and Ti species
reached the substrate surface as SrOx and TiOx species, incorporation of Sr and Ti
in the same plane was discouraged, so layer-by-layer growth was favored.
Evaluation of the (110) SrTi03 surface yields a different scenario. Both Sr
and Ti atoms are present in the same layer, and the fraction of oxygen atoms in
this layer is smaller than in the case of the (100) layer. In the (110) layer, both
the O/Ti and O/Sr ratios are 1, and the oxygen deficiency is alleviated by
interleaving layers composed entirely of oxygen. The structure of the (110)
surface is shown in figure 5-8.
Growth of (110) SrTi03 was favored at low oxygen pressures because the
oxygen/cation ratios were sufficiently low that a significant fraction of the Sr and
Ti species were not completely oxidized when they reached the substrate surface.
This enabled Sr and Ti to occupy the same plane, which favored growth of the
(110) orientation. Since SrTiOa is a compound with a limited range of non
stoichiometry, small deviations from the ideal SrTi03 can drastically alter the
chemical potential of the compound. The most prevalent type of defect in SrTi03
is oxygen vacancies.133 A plausible set of mechanisms which could have lead to
growth of the (110) phase were that the initial Sr-Ti-O oxygen deficient layer grew
because of the low oxygen/cation flux. Because the compound was oxygen
deficient at that point, the chemical potential for oxygen was greatly increased.
The oxygen chemical potential and oxygen deficiency was corrected for by growth

223
of the interleaving oxygen layers, which occured between laser pulses, when the
cation flux was zero.
SrTi03 films were deposited on (100) MgO substrates at 40 mTorr and 200
mTorr oxygen. In both cases, the SrTi03 films were (100) oriented. Table 5-1
shows that the lattice mismatch is a minimum in the <100> SrTi03 // MgO
directions, hence we expect the <100> SrTi03 direction to be parallel to the
substrate. However, this system differs from the SrTiO3/(1102) A1203 case
because the interfacial energy can be further minimized by matching the <010>
SrTi03 // MgO directions, so growth of (100) oriented SrTi03 is favored.
In the SrTi03 //(1102) A1203 system, only the <110> SrTi03 direction provides a
reasonably close lattice match, so minor changes in the deposition parameters can
induce either the (100) or (110) orientations to form.
A YBa2Cu307.x film subsequently deposited on the (110) SrTi03 barrier layer
at 200 mTorr was (103) oriented, which minimized lattice mismatch at the
YBa2Cu307_x/barrier layer interface. The resistance versus temperature behavior
of this film was slightly metallic, which was partially attributed to the lack of an
easy direction parallel to the substrate in which the charge carriers could travel.
In terms of the Halbritter equation, both the grain boundary and non-intrinsic
intragranular resistance were significant compared to the intrinsic intragranular
resistances. YBa2Cu307_x films deposited onto (110) SrTi03 substrates show
similar normal state resistances, but exhibit a sharp transition to the
superconducting state at 86 K. On the other hand, YBa2Cu307.x films deposited

224
on (110) SrTi03 barrier layers did not become superconducting until 77 K. This
discrepancy was attributed to microcracking in the YBa2Cu307_x film deposited on
(1102) A1203. Since the thermal expansion coefficient of YBa2Cu307.x is largest in
the <001> direction, a (103) oriented YBa2Cu307_x film will be subjected to
larger thermal stresses than an (001) oriented YBa2Cu307_x film, so microcracking
is more probable.
SrTi03 films deposited on (0001) A1203 at 200 mTorr were strongly (111)
oriented, and YBa2Cu307.x films deposited on this structure were primarily (001)
oriented, with minor fractions of (100), (013), and (113) phases also present. The
non-metallic normal state resistance was attributed to high-angle grain boundaries
resulting from the polycrystalline nature of the film, and the reduced T0 was
caused by thermally induced microcracks, similar to the case in which the
YBa2Cu307.x/Y-Zr02 film was deposited on (0001) A1203.
(1210) A1203 was the only plane of A1203 on which highly oriented SrTi03
barrier layers were not grown. X-ray diffraction shows only weak (110) SrTi03
peaks from films deposited at 200 mTorr onto (1210) A1203, and the principal
peak from a YBa2Cu307.x film deposited on this structure is the (103) orientation.
The resistance vs. temperature behavior of this film was semiconducting, with a T0
= 52 K. The poor electrical performance of this film resulted from several
factors, including high angle grain boundaries, microcracking, and judging from
the weak intensities of the x-ray peaks, a large fraction of amorphous Y-Ba-Cu-O
constituents.

LaAlOj Barrier Layers
In contrast to SrTi03 films, LaA103 barrier layers deposited on (1102) A1203
did not provide a suitable template on which to grow YBa2Cu307_x films. Despite
having a smaller lattice mismatch with (102) A1203 than SrTi03, and having a
similar crystal structure (LaA103 is a distorted pervoskite), x-ray diffraction of a
LaA103 film deposited on (1102) A1203 did not show any LaA103 peaks. The
YBa2Cu307.x/LaA103 film on (1102) A1203 was (001) oriented, but the x-ray
diffraction peaks were less intense than the (001) YBa2Cu307_x/(100) SrTi03 film
grown on (1102) A1203. The normal state resistance was slightly metallic, and the
film did not become superconducting until 58 K. The non-ideal normal state
resistance and depressed T0 were attributed to grain boundaries and lack of
completely crystallized YBa2Cu307.x phase in the film. Interpretation of the
resistance vs. temperature behavior in terms of the Halbritter equation showed
the intragranular resistance was high, since the curve did not extrapolate close to
0 ohms at 0 K; this was partially due to the presence of non-(001) oriented
grains. Raman spectra showed a large peak at 500 cm'1, which indicates that a
significant portion of the film is oriented with the <001> direction in the plane of
the film, or at some angle with the surface other than 90. This is consistent with
x-ray diffraction data. The increased thermal stresses in non-(001) oriented
YBa2Cu307_x grains may have induced microcracking.

226
Another factor which contributed to increased intragranular resistance was
incorporation of La into the YBa2Cu307.x film. For comparison, a YBa2Cu307.x
film deposited on (1102) A1203 using a YA103 barrier layer had lower resistances
and a higher transition temperature (T0 = 82 K) than did the YBa2Cu307.x films
grown on LaA103 barrier layers. Both barrier layers were amorphous, and this
indicates that La was more detrimental to the superconducting properties that
excess Y. It is difficult to determine whether the La is most damaging because it
segregates to the grain boundaries and forms insulating junctions through which
the superconducting charge carriers cannot penetrate, or if La diffuses into the
grains, replaces Y, and inhibits formation of the orthorhombic phase. Since the
Tonset for YBa2Cu307_x films deposited on (1102) A1203 using both LaA103 and
YA103 is ~ 89 K, it seems as though the temperature at which localized
superconductivity occurs within the grains is similar in both films. However, the
wide temperature region over which the resistance drops to zero in the
YBa2Cu307.x/LaA103 film on (1102) A1203 is indicative of poor intergranular
coupling. Thus we conclude that excess La is damaging to the YBa2Cu307.x film
because it segregates to the grain boundaries and inhibits superconducting charge
transport between grains. However, only (001) YBa2Cu307.x x-ray peaks were
observed in the film grown on the YA103 barrier layer, so the absence of
microcracking may also explain the difference in electrical properties.
YBa2Cu307_x/LaA103 films deposited on (1210) A1203 were even more degraded.
Resistance vs. temperature measurements show the film was semiconducting.

227
Although Tonset was 88 K, the transition was very broad and 0 ohm resistance
was not observed. X-ray diffraction showed only a weak LaA103 peak, and the
YBa2Cu307.x film grew with (103) and (001) orientations. All of the YBa2Cu307_x
peaks were weak, which indicates non-crystallized Y-Ba-Cu-O constituents. The
weak x-ray peaks may have also resulted from a crystalline YBa2Cu307_x film in
which the grain size was too small to yield x-ray diffraction peaks. Although the
LaA103 barrier layer films did not yield well defined x-ray peaks, the orientation
of the YBa2Cu307.x films was dependent on the orientation of the A1203
substrates. This dependence supports the hypothesis that the LaA103 barrier
layer and YBa2Cu307_x films were more crystalline than the x-ray diffraction
suggests, but the very small grain size broadened the x-ray diffraction peaks so
they were undetected. Like the YBa2Cu307_x/LaA103 film deposited on (1102)
A1203, the wide transition temperature range was attributed to intergranular
damage caused by La diffusion to the grain boundaries.
YBa2Cu307_x/LaA103 films deposited on (100) Si had very poor normal state
properties. SEM micrographs show the film was rough, and was separated into
isolated regions by an interconnected crack network. There are two primary
reasons why this film was heavily cracked. First, the large difference in thermal
expansion coefficients between YBa2Cu307_x and Si induced tensile stresses which
exceeded the fracture stress of YBa2Cu307_x. Second, cracking was more extensive
in this film than in the YBa2Cu307.x/Y-Zr02 film on (100) Si because the
YBa2Cu307.x film contained mixed orientations, and the non-(001) oriented grains

228
have a higher thermal expansion coefficient than the (001) oriented grains. This
difference in expansion leads to increased tensile stresses and cracking.
Y3O3 Barrier Layers
YBa2Cu307.x/Y203 films deposited on (100) Si had a morphology very similar
to the YBa2Cu307_x/LaA103 films deposited on Si. These films had a very rough
surface and showed an interconnected crack network, which drastically limited the
ability of normal state or superconducting charge carriers to tunnel across the
cracks. In addition, the rough surface increased the surface area and thus
percolation distance for the hole pairs. Auger data show that Ba diffused through
the Y203 film and accumulated at the Y203/Si interface, which probably
contributed to the degradation. This behavior may have been influenced by the
high interfacial free energy of the Y-Si interface, which could have provided a
large driving force for Ba migration to and reaction with the substrate. This
implies that transport paths existed through the Y203 films. Prior studies in which
thin (< 40 ) metallic yittrium films were deposited onto Si (111) substrates by
electron beam evaporation showed that Si reacts with Y to form YSi2134. The
YSi2 grew as three dimensional islands, and large pinholes were observed after
post-annealing at 800 C. In our study, Y203 films deposited on (100) Si at 730
C then cooled to room temperature appeared to be smooth. However, we
postulate that during the subsequent superconductor deposition and extended
anneal necessary to form orthorhombic YBa2Cu307.x, the Y203 barrier layer

229
reacted with Si and formed pinholes as the film coalesced, thus providing open
channels for Ba to diffuse towards the substrate. Although there is no direct
proof of the existence of a Ba-Si-O phase in the superconducting film, it seem
likely that the same pathways which permitted diffusion of Ba to the substrate
(postulated to be pinholes) would allow Si diffusion into YBa2Cu307_x.
YAlQj Barrier Layers
YA103 films deposited onto (1102) A1203 were similar to LaA103 films in that
no crystalline peaks were observed. However, the electrical properties of
YBa2Cu307_x/YA103 films were far superior to YBa2Cu307_x/LaA103 films on
(1102) A1203. The normal state resistance values were more metallic, although
they were not ideally metallic, indicating grain boundary resistance was still
significant. The higher transition temperature and sharp drop to T0 at 82 K was
attributed to improved coupling between the YBa2Cu307_x grains, which shows
that excess Y in the YBa2Cu307.x film was less damaging than La. The extent to
which A1 was responsible for the slight drop in T0 is difficult to quantify. At 750
C, A1 diffuses into YBa2Cu307.x and degrades the superconductivity. Initial
experiments in which YBa2Cu307.x films were deposited directly onto A1203
substrates with no barrier layer produced white films which were almost
transparent, and there was no evidence of a superconducting phase. Clearly, A1
incorporation into YBa2Cu307.x films degrades the superconductivity. The slight
drop in T0 for YBa2Cu307.x/YA103 films deposited on (1102) A1203 probably

230
resulted from Al diffusion to the YBajCojO^, grain boundaries and reduced the
intergranular coupling between grains, which would account for the broadened
transition between Tonset and T0. Despite the adverse effects of A1 incorporation
into the YBa2Cu307_x film, the combination of Y and A1 oxides in the barrier layer
caused only a slight degradation in T0 of the YBajQ^O^ film.
Much of the phenomena responsible for the microstructure observed in the
YBa2Cu307.x/YA103 films can be inferred from phase equilibria. The flow chart
diagram for the Y-Al-Si-Cu-O system is shown in appendix B. This diagram
shows that incorporation of A1 into the Y-Cu-Si-O system lowers the liquidus
temperature. The flow chart predicts that the lowest temperature for a liquid
phase occurs at the l^Si02-Al203-CuA102-Cu0 five-phase equilibria, and SEM
micrographs of YBa2Cu307.x/YA103 films on Si indicate a liquid phase may have
formed at the grain boundaries.
The sharpest transition for YBajQijO^ films deposited on Si substrates were
observed when YA103 barrier layers were used to prevent interdiffusion. The
sharp transition was attributed primarily to the greatly reduced cracking (relative
to the YBa2Cu307.x/Y-Zr02 film on Si). The most probable reason for reduced
cracking in these films was that the thermally-induced tensile stresses were
relieved by viscoelastic relaxation of the YBa2Cu307.x/YA103 film. We believe
A10x diffused to the grain boundaries of the YBa2Cu307.x film during the = 1 hour
anneal at 730 C (required to form the orthorhombic YBa2Cu307.x phase),
forming liquid and CuA102 phases. It is probable that incorporation of BaO

231
lowered the liquidus temperature below the temperature suggested by the flow
charts. The glass transition temperature for Y-Al-Si-O is 930 C,135 and SEM
micrographs of a YA103 film deposited on Si at 730 C were featureless.
However, after the deposition and in-situ annealing of the YBajQ^O^ films, the
growth of circular, overlapping grains is apparent. The circular shape of the
grains is evidence that the viscosity of the film was sufficiently low during the
YBa2Cu307_x growth process to adopt a configuration which minimized the surface
tension. The overlap of the YBa2Cu307_x grains resulted from YBajCojO-^ film
contraction, which was the mechanism by which the tensile stresses were relieved.
We conclude that incorporation of A10x into the YBajCujO,.* film was
necessary for viscoelastic relaxation of the thermal stresses. However, the A10x
also contributed to degradation of the superconducting properties of the film
because it degraded intergranular coupling of the superconducting holes at the
grain boundaries.
Al6Si3Q13 Barrier Lavers
Deposition of an Al6Si2013 film onto (100) Si substrates prior to growth of
YA103 or LaA103 barrier layers and the YBa2Cu307_x film dramatically altered
the microstructures of the YBa2Cu307_x films, compared to the cases in which an
Al6Si2013 was not deposited. Figures 4-72 and 4-73 show that the surface
morphologies for the YBa2Cu307.x/(YA103 or LaA103)/Al6Si2013 films on Si are

232
very similar. This indicates that the Al6Si2013 layer was equally successful at
preventing the diffusion of Y or La to the substrate. Comparison of these
micrographs with the morphologies of YBa2Cu307.x/(YA103 or LaA103) films
deposited on Si also showed that the microstructures of YBa2Cu307_x films grown
on Si using YA103 or LaA103 barrier layers (without the Al6Si2013 layer) resulted
largely from a rare earth silicide type of reaction. X-ray diffraction shows that
both the YBa2Cu307.x/(YA103 or LaA103)/Al6Si2013 films on Si were (001)
oriented, and the smaller degree of cracking relative to the YBa2Cu307_x/LaA103
film on Si was attributed to the reduced strain associated with a highly (001)
oriented film. YBa2Cu307_x/LaA103 films deposited on Si contained both (001)
and (103) oriented grains, and the cracking was much more extensive.
The normal state electrical resistances for the YBa2Cu307_x/(YA103 or
LaA103)/Al6Si2013 films on Si were both slightly metallic, with T0 values of 61.5
and 59 K, respectively. Analysis of this behavior in terms of the Halbritter
equation suggests that the predominant contribution to the normal state electrical
resistances were intergranular defects (i.e. cracks), and the tail in the resistance
versus temperature profile could be attributed to weakened coupling of
superconducting charge carriers across the grains. However, examination of the
Raman spectra show that the intragranular microstructures of these two films are
markedly different. Comparison of Raman data from the YBa2Cu307_x/(YA103 or
LaA103)/Al6Si2013 films on Si shows that the film containing YA103 has a
pronounced peak at 500 cm"1, indicative of a large fraction of non-(001) oriented

233
grains. The lack of x-ray diffraction from these non-(OOl) peaks indicates that the
grain size was too small to yield diffraction peaks. The small size of the 500 cm'1
Raman peak from the YBa2Cu307.x/LaA103/Al6Si2013 film shows that the fraction
of non-(OOl) oriented grains is much smaller, but the wide breadth of this peak
indicates disorder along the oxygen chains or substitution of Cu(l) by impurity
atoms (such as A1 or Si). Both types of phenomena could have changed the
bonding environment of the 0(4) atoms, and cause the spread in the 500 cm'1
peak. According to this data, intragranular defects should have been a significant
cause of degradation in the YBa2Cu307_x/LaA103/Al6Si2013 film on Si. Except for
the * 2x difference in the normal state resistance values for the YBa2Cu307.x/
(YA103 vs LaA103)/Al6Si2013 films on Si, the resistance versus temperature
patterns were very similar. Using the Halbritter equation to uncover the
predominant type of defect responsible for the electrical behavior, the normal
state resistances were dominated by cracking, and the cracking was less
pronounced in the YBa2Cu307.x/LaA103/Al6Si2013 film. However, the Raman
spectra shows the intragranular microstructures were quite different for the two
films.
Comparisons between YBa2Cu307.x/(YA103 or Y203)/Al6Si20i3 films
deposited on (1102) A1203 provided insights into the role which the mechanical
properties of the barrier layers play in the overall degradation of the YBa2Cu307..x
films. Since the thermal expansion coefficient of Al6Si2013 is smaller than that of
(1102) or any of the other films used in this study, prediction of how the Al6Si2013

234
film influences the stress in the subsequently deposited barrier layer and
YBa2Cu307.x films presents an interesting dichotomy. According to the continuum
mechanics model for thin films, one would expect the Al6Si2013 film to have a
negligible impact on the stress in the other films. The small expected effect of
the Al6Si2013 film would be to reduce the total tensile stress in the multilayered
film structure.
There were many similarities between YBa2Cu307_x/(YA103 or
Y203)/Al6Si2013 films deposited on (1102) A1203. Both barrier layer structures
were composed of the same components, hence differences in degradation
resulting from incorporation of different types of cations in the YBa2Cu307.x film
were minimized. Both YBa2Cu307.x films were predominately (001) oriented, yet
the resistance vs. temperature data for these films were vastly different. The
YBa2Cu307.x/YA103/Al6Si2013 film deposited on (1102) A1203 was metallic, with
a T0 = 75 K. The normal state resistance values and T0 for this film indicated
that incorporation of Al6Si2013 into the barrier layer structure slightly degraded
the performance of the YBa2Cu307.x film.
Degradation of the YBa2Cu307_x/Y203/Al6Si2013 film on (1102) A1203 was
much more extensive. The normal state resistance behavior of this film was
semiconducting, with a long transition to the superconducting state at T0 less than
39 K. An initial explanation for the degradation of the YBa2Cu307.x film was
that the growth of the Y203 was three dimensional, so pathways existed by which
Si could diffuse to the YBa2Cu307.x film. However, Raman spectra showed that

235
the intragranular regions for both the YBa2Cu307.x/(YA103 or Y203)/Al6Si2013
films on (1102) A1203 were nearly ideal, with the asymmetry of the 335 cm*1 peak
and the small, narrow peak at 500 cm'1 indicating highly oxygenated, (001)
oriented YBa2Cu307_x grains in both films.
We believe chemical interdiffusion and incorporation of Si into the
YBa2Cu307.x film was not the mechanism responsible for degradation of
YBa2Cu307_x films when an Al6Si2013 layer was added to the barrier layer
structure. The degradation was caused by increased cracking in the YBa2Cu307.x
films. This data shows that the Al6Si2013 layer increased the magnitude of the
tensile stresses in the YBa2Cu307.x films, which indicates that the "continuum
mechanics model"48 for stress transfer between the substrate and various film
layers is not completely valid. However, this model does the best job of
explaining the experimentally observed behavior. The data also suggests that
cracking is not necessarily initiated in the YBa2Cu307.x layer, but the poor
electrical properties of the YBa2Cu307.x/Y203/Al6Si2013 film may result from the
poor fracture strength of the Y203 layer, which would allow cracks to nucleate
and propagate into the YBa2Cu307.x layer. The tensile stresses which accompany
crystallization may have also increased the stress in the Y203 layer, and induced
cracking. (222) Y203 diffraction peaks were observed, whereas the Y-A1-0 layer
did not crystallize. The thermally induced stresses in the amorphous Y-Al-O layer
were less than in the Y203 layer, so cracking was suppressed.

CHAPTER 6
SUMMARY AND CONCLUSIONS
YBa2Cu307_x is a promising material for low loss, low dispersion applications
such as electrical interconnects and microstrip lines on Si and A1203 substrates.
However, YBa2Cu307_x deposited directly onto Si or A1203 forms interfacial
phases which degrade both the YBa2Cu307.x film and substrate, the critical
current densities (Jc) of YBa2Cu307_x films deposited directly on (100) Y-Zr02
substrates are sensitive to the deposition conditions, and it is difficult to
reproducibly deposit films with high Jc values. To suppress interfacial
degradation, barrier layer films are deposited prior to growth of the YBa2Cu307.x
films. In this study, YBa2Cu307.x and barrier layer films were deposited in-situ
using a pulsed laser deposition process, with an excimer laser operating at a
wavelength of 248 nanometers. The barrier layers and YBa2Cu307_x films were
sequentially deposited at 730 C, without lowering the substrate temperature.
The orthorhombic, superconducting YBa2Cu307_x phase was formed during the
cooling process, and no post-annealing was required. Chemical reactions at the
substrate/barrier layer as well as at the barrier layer/YBa2Cu307_x interface
significantly influenced the normal state resistivities, the superconducting
transition temperatures, and Jc values of YBa2Cu307_x films. Optimal electrical
236

237
properties were observed when the YB^Ci^O^ films were predominately (001)
oriented, and free of microcracks.
The orientation and lattice parameters of the barrier layers dictated the
orientation of the YBa2Cu307.x films. For SrTi03 barrier layers deposited on
(1102) A1203 substrates, the SrTi03 texture was heavily influenced by the oxygen
pressure during deposition. SrTi03 films deposited at 200 mTorr 02 were (100)
oriented, and the subsequently deposited YBajCujO^ film grew with an (001)
orientation. When the oxygen pressure during the SrTi03 deposition was dropped
to 40 mTorr, the SrTi03 layer adopted the (110) orientation, and the YBa2Cu307.x
film was (103) oriented. In both cases, the YBa2Cu307_x film grew with the
orientation which minimized lattice mismatch at the YBa2Cu307.x/SrTi03
interface.
The normal state and superconducting properties of the (001) and (103)
YBa2Cu307.x films were dramatically different. The normal state resistance of the
(001) YBa2Cu307.x/(100) SrTi03 film deposited on (1102) A1203 was metallic,
with a sharp transition to the superconducting state at 83 K. The Jc was 2.5 X106
amps/cm2 at 4.5 K, and the microwave resistance was 10 milliohms at 36 GHz
and 65 K. These values indicate the YBa2Cu307.x film had good in-plane
epitaxy. By contrast, the normal state electrical resistance of the (103)
YBa2Cu3O7.x/(110) SrTi03 film on (1102) A1203 was only slightly metallic, with a
transition to the superconducting state at 77 K. At 4.5 K, the Jc was 6.8 X103
amps/cm2. The primary source of degradation in this film was microcracking,

238
which occured because the thermally induced stresses exceeded the fracture
strength of YBa2Cu307_x. The stresses in the (103) oriented YBajQ^O^ film
were greater than in the (001) oriented film because the thermal expansion
coefficient is highest in the <001> direction, and the contribution of the <001>
direction to the in-plane thermal expansion of the film is greater in the (103)
oriented films than in the (001) YBajQ^O^ films.
YA103 and LaA103 films deposited on (1102) A1203 were both amorphous,
but the electrical properties of the YBa2Cu307x films deposited on the barrier
layers were quite different. The normal state resistance of the YBa2Cu307_
x/YA103 film deposited on (1102) A1203 was metallic, with a transition to the
superconducting state at 82 K. However, the YBa2Cu307_x/LaA103 film
deposited on (1102) A1203 was less metallic, and the transition to the
superconducting state was depressed to 57 K. The difference in electrical
properties was attributed to La segregation to the grain boundaries and decreased
intergranular coupling of the superconducting hole pairs.
Because of the large difference in thermal expansion coefficients between
YBa2Cu307.x and Si (lSxlO'VC vs 3.8xl0'6/oC, respectively) thermally induced
cracking is prevalent in the YBa2Cu307.x films. Elastic continuum mechanics
predicts that the thermally induced stress in a film will result almost entirely from
the difference in thermal expansion coefficients between the film and substrate,
and will be virtually independent of the presence of any other layers.

239
YBa2Cu307_x films (2500 ) were deposited onto Si substrates using either Y-Zr02
or YA103 barrier layers. The YBa2Cu307.x/Y-Zr02 film became superconducting
at 76 K, and was heavily cracked. However, the superconducting transition
temperature of the YBa2Cu307.x/YA103 film was 78 K, and the cracking was not
as extensive. Scanning Auger analysis showed diffusion of A1 to the interface to
form an Al-Si-O layer, followed by Y-Al-Si-O and Y-Al-O layers. The reduced
cracking in the YBa2Cu307.x/YA103 film deposited on Si was attributed to
relaxation of the thermally induced stresses by viscoelastic strain relaxation at the
YA103/Si interface.
The Jc values of YBa2Cu307.x films deposited on single crystal Y009Zr091O195
substrates cut 5 12 degrees from the (100) planes were significantly enhanced by
depositing a Y-Zr02 barrier layer film prior to the YBa2Cu307.x film. The normal
state properties of YBa2Cu307.x films deposited with or without Y-Zr02 barrier
layers were similar, with superconducting transition temperatures ranging from 86
to 89 K in both cases. However, the Jc for the YBajQ^O^ film deposited
directly onto the Y-Zr02 substrate was 6.8 xlO3 amps/cm2 at 4.5 K, while the Jc
of the YBa2Cu307_x/Y-Zr02 film was lxlO7 amps/cm2 at 4.5 K. The
composition of the Y-Zr02 target was Y013Zr0 87O193, so the increased critical
current density was attributed to pinning of the magnetic flux by excess Y203 in
the grain boundaries. To test this hypothesis, Y203 or Zr02 barrier layer films
were deposited prior to the YBajQrjO^ films. The Jc of the YBa2Cu307.x/Y203
film was lxlO7 amps/cm2 at 4.5 K, while the critical current density of the

240
YBa2Cu307.x/Zr02 film on the Y-Zr02 substrate was 9X104 amps/cm2 at 4.5 K.
These data confirm that excess Y203 in the grain boundaries significantly
improves intergranular coupling of the superconducting hole pairs, and Jc values
similar to those obtained from epitaxial YBa2Cu307.x films grown onto single
crystal substrates.
Selected barrier layers were used to dramatically improve the superconducting
properties of YBajQ^O^ films deposited on various substrates. The type of
barrier layer which worked best was different for each of the substrates, and was a
function of the lattice matching, and chemical reactivity at both the barrier
layer/substrate and YBa2Cu307_x/barrier layer interfaces. In this study, we
determined that SrTi03, YA103, and Y203 were the best barrier layer materials
for YBa2Cu307_x films deposited on (1102) A1203, Si, and Y-Zr02 substrates,
respectively.

APPENDIX A
CALCULATION OF X-RAY ABSORPTION DEPTH FOR YBa2Cu306J
Atomic weight of: yittrium = 88.9 g/mole
barium = 137.3 g/mole
copper = 63.5 g/mole
oxygen = 16.0 g/mole
Total weight = 88.9 + 2(137.3) + 3(63.5) + 6.5(16) = 658.2 g/mole.
Weight fraction of: yittrium = 88.9/658.2 = 0.135
barium = 274.6/658.2 = 0.417
copper = 127.0/658.2 = 0.290
oxygen = 104.0/658.2 = 0.158.
Mass absorption coefficient (u/p) of:
yittrium = 127.1
barium = 336.1
copper = 51.5
oxygen = 11.0
(u/p) for YBa2Cu307.x = 2wi(p/p)i
= 0.135(127.1) + 0.417(336.1) + 0.290(51.5) + 0.158(11.0) =
= 174.0
241

242
I/Io = exp -(fi/p)pt, where I/Iq is the ratio of x-ray intensity at a given depth (= t)
relative to the incident x-ray intensity, and p is the density of the phase under
consideration (= 6.38 g/cm3 for YBa2Cu307_x). Assuming the x-ray penetration
depth is equal to the depth at which I/Iq drops to 0.368, we obtain:
0.368 = exp -(174.0)(6.38)t.
Solving for t yields an x-ray penetration depth of 9 pm.

APPENDIX B
FLOW CHART OF THE Y AI Si Cu O SYSTEM
A flow chart of the Y Al Si Cu O system was constructed in order to
determine the chemical reactions responsible for the microstructure observed in
the YBa2Cu307.x/YA103 films on Si and A1203 substrates.
Construction of the flow chart requires that all phases appear or disappear at
equilibrium in a manner consistent with the Gibb's phase rule:
P+F=C+1 (A1)
where P = the number of phases in equilibrium, F = degrees of freedom, C =
the number of components, and the pressure is arbitrarily held at 1 atmosphere.
In this construction, only invarient points where F = 0 are considered. Three-
phase fields must begin or terminate at binary three-phase equilibria, ternary four-
phase equilibria, or quaternary five-phase equilibria (i.e. conditions where F = 0,
so the temperature at which the phases are in equilibrium, and the compositions
of each phase, are fixed). Since the temperatures of the three-phase equilibria
are known for the binary systems, the four-phase equilibria and temperature range
in which they must occur can be determined using a self-consistent analysis. In
this analysis, all three-phase regions emanating from or terminating at binary
phase diagrams are required to connect ternary four-phase equilibria. To be
243

244
consistent with Gibb's phase rule, there must be a total of 4 three-phase regions
connected to each four-phase equilibria. Once the four-phase equilibria are
determined, the five-phase equilibria can be determined using the same process.
In this study, the Y-Al-Si-Cu-O flow chart was constructed from the
existing binary oxide phase diagrams, and the Y203 A1203 Si02 liquidus phase
diagram. In our flow chart diagrams we treated the metallic oxides as
components, and did not take into account the effect of oxygen nonstoichiometiy
on reaction temperature; hence the most complex flow chart was the quaternary
diagram. For clarity, only ternary and quaternary reactions which involve a liquid
phase are listed in the flow charts.
This flow chart was constructed in order to uncover the chemical reactions
which caused the microstructures observed in figure 4-54 to emerge after
deposition of the superconductor on the YA103 barrier layer. Since CuO has the
lowest melting point of the metallic oxides which comprise YBa2Cu307.x (the
melting points of CuO, BaO, and Y203 are 1362, 1923, and 2410 C, respectively),
we hypothesized that CuO lowers the liquidus temperatures in the system, and
chemical interactions between CuO in the superconductor and the Y Al Si O
system were responsible for the microstructures in the YBa2Cu307_x/YA103 films
on Si.
The ternary flow charts are shown in figures Al A4. The binary phase
diagrams used to construct the flow charts are shown in figures A6 All. The
Y203 CuO phase diagram was not found, so the types of reactions which are

245
likely to occur in this system were inferred from the known YxQiyOz compounds.
The quaternary flow chart is shown in figure A5. This chart chart lists the four-
phase equilibria (determined in figures Al A4) in the columns marked "ternary",
and the five-phase equilibria in the "quaternary" column.

FLOW CHART FOR Y Al Si O
L <** Al6Si2013 + AI2
(1912 C)
)3 L ** Y203 + Y4A1209
(1940 C)
L Y4A1209 + YA103 -
(1875 C)
L YA103 + Y3A15012
(1850 C)
L * Y2Si05 + Y4Si3012
(1900 C)
-L ~ Y203 + Y2Si05
(1800 C)
L + Y203 Y4A1209 + Y2Si05
YA103 Y3A15012 Y4Si3012
L Y4A1209 Y2Si05
L + Y4A12Q9 > Y2SiOs + YA1Q3
Y4A1209 Y2Si05 YA103
L Y2Si05 YAJ03
Figure A-l. Flow chart of the Y Al Si O system.

L + YA103 Y3A15012 + Y2Si05
YA103 Y3A15012 Y2SiOs
L Y3A15012 Y2Si05
L + Y2Si05 Y3A15012 + Y4Si3012
Y2Si05 Y3A15012 Y4Si3012
L Y3A15012 Y4Si3012
L ** Y3A15012 + A1203
(1760 C)
L + Y3A15012 ** Y4Si3012 + A1203
Y3A15012 Y4Si3012 AI203
L Y4Si3012 A1203
Figure A-l. Continued

L ** Si02 + Al6Si2013
(1546 C)
L + AI203 ** Al6Si2013 4- Y4Si3012
A1203 Al6Si2013 Y4Si3012
L Al6Si2013 Y4Si3012
-L + Y4Si3012 ** Y2Si20
(1775 C)
L + Y4Si3012 ** Al6Si2013 + Y2Si207
Y4Si3012 Al6Si2013 Y2Si207
L Al6Si2013 Y2Si207
L ** Y2Si207 + Si02
(1660 C)
L ** AI6Si2013 + Y2Si207 + Si02
Al6^^2^13 Y2Si207 Si02
Figure A-l. Continued.

FLOW CHART FOR Al Cu Si O
A1,Q, SiO, CuO ALO, TERNARY CuO SiQ.
v
Figure A-2. Flow chart of the Al Cu Si O system.

L ** CuA102 + CuO
L ** CuO + SiO-,
(1060 C)
L ** Si02 + CuA102 + CuO
SiO-, CuAlO-, CuO
it it
Figure A-2. Continued.

FLOW CHART FOR Y Cu Si 0
Y,Q, CuO
Y^O-t SiO-7
TERNARY
CuO Si02
L + Y303 + Y2CU2O5
L * Y2Si05 + Y4Si3012
(1900 C)
L ~ Y203 + Y2Si05
(1800 C)
L ** Y4Si3012 + Y2SiOs
(1775 C)
L ** Y2S2O7 + Si02
(1660 C)
L + Y203 ** Y2Cu205 + Y2Si05
Y203 Y2Cu205 Y2Si05
L Y2Cu205 Y2Si05
Figure A-3. Flow chart of the Y Cu Si O system.

L + Y2Si05 Y4S3012 + Y2Cu2Os
Y2S05 Y4Si3012 Y2Cu2Os
L Y4Si3Or Y2Cu205
~i l
L + Y4Si3012 Y2Si207 + Y2Cu205
Y4S3012 Y2Si207 Y2Cu205
L Y2Si207 Y2Cu205
I i ,
L + Y2Si207 ** Si02 + Y2Cu2Os
Y2S207 Si02 Y2Cu2Os
L Si02 Y2Cu205
Ni
Ul
Figure A-3. Continued.

L <* Y-yCu7Or + CuO
(1060 C)
L ** Y2Cu205 + CuO 4- Si02
L ** CuO + Si02
(1050 C)
Y2Cu205 CuO Si02
N>
Ut
w
Figure A-3. Continued.

FLOW CHART FOR Y Cu A1 0
Y,Q, CuO Y,Oo ALO, TERNARY CuO ALO,
L ~ Y203 + Y4A1209
(1940 C)
L + Y4A1209 YA103
(1875 C)
L YA103 + Y3A15012
(1850 C)
L ~ Y3A15012 + A1203
(1760 C)
L + Y203 * Y2Cu2Os
L + Y203 ** Y2Cu205 + Y4A1209
Y203 Y2Cu205 Y4A1209
L Y2Cu205 Y4A1209
Figure A-4. Flow chart of the Y Cu Al O system.

L + Y4A]209 YA103 + Y2Cu205
Y4A1209 YA103 Y2Cu205
L Y4A1209 Y2Cu205
1 I
L + YA103 Y3A15012 + Y2Cu2Os
yaio3 Y3A15012 Y2Cu205
L Y3A15012 Y2Cu205
1 1
L + Y3A15012 ** Y2Cu205 + A1203
Y3A15012 Y2Cu205 A1203
L Y2Cu205 A1203
Figure A-4. Continued.

L ** Y2Cu205 + CuO
L + Y2Cu205 * CuO + A1203
Y2Cu205 CuO A1203
L CuO A1203
L A1203 + CuA102
(1180 C)
L ** CuA102 + CuO
(1040 C)
L ** AJ203 + CuO + CuAI02
AI203 CuO CuAI02
!s>
ON
Figure A-4. Continued.

FLOWCHART FOR Y Al Si Cu O
TERNARY
L + Y203 Y4A12
TERNARY
QUATERNARY
TERNARY
09 + Y2Si05
L + Y203 ** Y2Cu205 + Y2SOs
-L + Y203 ** Y2Cu205 + Y4A1209
L + Y203 ** Y4A1209 + Y2Cu205 + Y2Cu2Os
L Y4A1209 Y2Si05 Y2Cu205
L + Y4A1209 o Y2SOs + YAI03
L + Y4A1209 ~ YA103 + Y2Cu2Os
L + Y4A1209 ~ Y2Si05 + Y2Cu205 + YA103
L Y2Si05 Y2Cu2Os YAIO,
L + YA103 o Y3A15012 + y2so5
L + YA103 *> Y3A15012 + Y2Cu2Os
L + YA103 ~ Y3A15012 + Y2SiOs + Y2Cu205
L Y3A15012 Y2Si05 Y2Cu205
L + Y2SiOs Y3A15012 + Y4Si3012
1
L + Y2SiOs o Y4S3012 + Y2Cu2Os
L + Y2SiOs ** Y3A15012 + Y4S3012 + Y2Cu205
L Y3A15012 Y4Si3012 Y2Cu2Os
N
Figure A-5. Flow chart of the Y-Al-Si-Cu-O system.

L + Y3A15012 ** Y4Si3012 + A1203
L + Y3A15012 * Y2Cu205 + A1203
L + Y3A15012 ** Y4Si30j2 + Y2Cu205 + A1203
L Y4Si3012 Y2Cu205 A1203
L + Y4Si3012 ** Y2Si207 + Y2Cu205 j
L + Y4Si3012 + Y2Cu205 ** Y2Cu205 + A1203
L Y4Si3012 Y2Si20- A1203
L Y2Cu205 Y2Si207 A1203
L + A1203 ** Al6Si2013 + Y4Si3012
1
L + Y4Si3012 ** Al6Si2013 + Y2Si207
L + Y4Si3012 ** A1203 + Y2Si207 + Al^Si2013
L AJ203 Y2Si207 Al6Si2013
L + Y2Si207 *> Si02 + Y2Cu205 i
L + Y2Cu205 + Y2Si207 ** Si02 + A1203
L Y2Cu205 Si02 AI,0
2^3
L Y2Si207 Si02 A1203
Figure A-5. Continued.
S)
00

Si02 A1203 CuA102 CuO
is)
Lf
VO
Figure A-5. Continued.

260
A1203 Y203
Figure A-6. Phase diagram of the Y203 A1203 system. Reference 136.

Figure A-7. Phase diagram of the Y203 Si02 system. Reference 137.

262
Al203-Si02
Figure A-8. Phase diagram of the A1203 Si02 system. Reference 138.

263
Figure A-9. Liquidus projection of the Y203 A1203 Si02 system. Reference
139.

264
2000
1700
1400
1100
i 1 1 1 1 1 1 1 r
Li q u i d
Liq.
+
ai2o3
\ CuO ai2o3
\ +
CuO-Al203 -f CuO
i i i i
ALO, 20 40 60 80 CuO
Mol.% i/2cu20
Figure A-10. Phase diagram of the A1203 CuO system. Reference 140.

Cu20 Si02
2000
1600
1230
800
r
, |
Two Liquids
i '
J -1690
/
/

/
/
/
/
Si02 + Liquid
/
/
v /

\ /
v 1060
I
s*
CO
1
Cu20 + Si02
I 1
Cu20
25
50
75
1713
SiO-
Figure A-ll. Phase diagram of the CuO Si02 system. Reference 141.

266
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BIOGRAPHICAL SKETCH
The author was born November 7, 1959, in St. Louis, Missouri. He received
a Bachelor of Science degree in ceramic engineering from the University of
Missouri at Rolla in 1982, and a Master of Science in ceramic engineering from
the University of Illinois at Urbana-Champaign in 1985.
276

I certify that I have read this study and that in my opinion it conforms to
acceptable standards of scholarly presentation and is fully adequate, in scope and
quality, as a dissertation for the degree of Doctor of Philosophy.
Paul H. Holloway, Chair u
Professor of Materials Science
and Engineering
I certify that I have read this study and that in my opinion it conforms to
acceptable standards of scholarly presentation and is fully adequate, in scope and
quality, as a dissertation for the degree of Doctor of Philosophy.
4teza Abbaschian
Professor of Materials Science
and Engineering
I certify that I have read this study and that in my opinion it conforms to
acceptable standards of scholarly presentation and is fully adequate, in scope and
quality, as a dissertation for the degree of Doctor of Philosophy.
Timothy J. (Anderson
Professor of Chemical Engineering
I certify that I have read this study and that in my opinion it conforms to
acceptable standards of scholarly presentation and is fully adequate, in scope and
quality, as a dissertation for the degree of Doctor of Philosophy.
Richard Associate Professor of Materials Science
and Engineering

I certify that I have read this study and that in my opinion it conforms to
acceptable standards of scholarly presentation and is fully adequate, in scope and
quality, as a dissertation for the degree of Doctor of Philosophy.
Robert T. DeHoffy
Professor of Materials Science
and Engineering
This dissertation was submitted to the Graduate Faculty of the College of
Engineering and to the Graduate School and was accepted as partial fulfillment of
the requirements for the degree of Doctor of Philosophy.
December, 1992
Winfred M. Phillips
Dean, College of Engineering
Madelyn M. Lockhart
Dean, Graduate School

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TITLE:
Mueller, Carl
Microstructure/electrical property correlations for YBa2Cu307-
x/barrier layer films deposited on A1203 silicon and yittria-
stabilized zirconia substrates / (record number: 1904702)
PUBLICATION
DATE:
1992
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Ill
Both films show a sharp drop to zero resistance at 86-87 K, but the normal state
resistance of the YBajQ^O^* film deposited on a Y-Zr02 barrier layer was more
Table 4-1. Thermal expansion coefficients for YBa2Cu307.x, substrates, and
barrier layer films. Lattice parameter mismatches at the YBa2Cu307_x/
barrier layer and barrier layer/substrate interfaces are also calculated.
The lattice parameter for SflBa2Cu307_x is a = 3.82 .
MATERIAL
THERMAL
EXPANSION
COEFFICIENT
(xlO'6/C)
LATTICE PARAMETER
MISMATCH (%) WITH
RESPECT TO
YBa2Cu307.x
Si
YBa2Cu307.x
13.2

-0.5
Si
3.8
-0.5

ai2o3
6.7(||[0001])
5.0 (J-100011)
Y-ZrOz
13.5
+5.1
-5.3
SrTi03
9.4
-2.1
+ 1.5
LaA103
10
+ 0.8
-1.3
yaio3
?
+3.0
-3.4
y2o3
9.2
+ 1.9
-2.3
Y-Al-Si-O
3.1-7.0


i3
5.1
-0.7
+ 0.3


59
stoichiometry were grown in vacuum (typically < 10"5 Torr) onto unheated
substrates. The films were then post annealed in flowing 02 at 850 900 C,
thereby forming the tetragonal YBa2Cu307_x phase.56
During cooling from the high temperature tetragonal phase, YBa2Cu307_x
undergoes large structural changes. Understanding what these changes are, and
how they are affected by temperature and oxygen pressure are critically important
for optimizing growth of YBa2Cu307.x films. The temperature at which
YBa2Cu307_x transforms from the non-superconducting tetragonal phase to the
superconducting orthorhombic phase is dependent on oxygen pressure.71 In 100%
oxygen, the transition occurs near 700 C. The oxygen content and change in
oxygent content as a function of temeprature are shown in figure 2-13 for bulk
YBa2Cu307.x in 100 % oxygen. If the oxygen pressure is lowered to 20% oxygen,
the tetragonal s* orthorhombic transition temperature is lowered to 670 C, and
in 2% oxygen the transition is depressed to 620 C. Ordering of the oxygen
atoms into one-dimensional Cu-O chains along the <010> direction is the
primary mechanism responsible for the tetragonal it orthorhombic transition.
The dramatic increase in the <010>, and decrease in <100> lattice parameters
as the orthorhombic phase is formed are shown in figure 2-14.
The temperature at which the transition occurs affects the kinetics of the
phase change. The transition occurs via a nucleation and growth process, in which
the ordering of oxygen along <010> and lengthening of the <010> lattice
parameter begins at a grain boundary or free surface.72 If the transition


RESISTANCE (OHMS)
144
10
8
6
4
2
0
0
Figure 4-36.
YBa2Cua07_x/SrTi03 ON (1102) Al20g.
SrTi03 BARRIER LAYERS DEPOSITED AT
50 100 150 200 250 300
TEMPERATURE (K)
Resistance versus temperature data for YBa2Cu307.x/SrTi03 films
deposited on (1102) A1203. The orientations of the YBajCujO^
films were determined by the SrTi03, which, in turn, were
determined by the oxygen pressure during the deposition.
Resistance values for the (103) oriented YBa2Cu307.x film are
multiplied by 0.50.


166
Figure 4-54. Scanning electron micrograph of a YBa2Cu307_x/YA103 film on
(100) Si.


58
normal.66 Apparently the highly peaked, forward-directed component is primarily
responsible for stoichiometry in the growing film. One explanation for this
behavior is that the laser generated plasma results in rapid evaporation from the
target surface.65 The gas is initially at high pressures because the rate of
evaporation
is greater than the rate at which atoms and ions can leave the target surface. The
plasma expands into the vacuum, creating a supersonic molecular beam. Time-of-
flight measurements indicate that mean kinetic energies for Cu(l), Y(l), and
Ba(l) are 41.3, 43.4, and 47.9 eV, respectively for particles generated by a
193 nm laser.69 Time resolved optical spectroscopy measurements showed that at
oxygen partial pressures less than 1.3X10"4 Torr, the velocities of neutral and
ionized atomic species, as well as the diatomic species (such as YO, BaO, CuO)
were approximately 106 cm/sec. Increasing the oxygen pressure to 10'2 Ton-
reduced the velocities of the atomic and diatomic species to approximately
5x10s cm/sec.70
Although the ability to deposit a stoichiometric film is an extremely important
parameter for growing superconductor films with low surface resistivities and high
Je, there have been other advances in film growth techniques which have
significantly improved the film quality. Optimization of the growth temperature
and oxygen pressure during deposition, as well as the oxygen pressure and cooling
rate after the deposition, have generated YBa2Cu307.x films with nearly ideal
superconducting properties. Initially, amorphous films with the correct


56
Electronic mechanisms are also operative during laser deposition.66 Photons
with energies greater than the first ionization potential energies (7.726, 5.512, and
6.38 eV for Cu, Ba, and Y, respectively) will excite the target atoms, thereby
breaking bonds and causing ejection of ions. Experiments in which YBa2Cu307.x
film morphology and electrical properties of films grown at different wavelengths
showed conclusively that smoother films with fewer and smaller particulates, as
well as lower normal state resistivities and higher Jc values, were obtained in films
grown at shorter wavelengths (figure 2-12).57
The smaller particulate sizes which are observed when YBa2Cu307_x films are
deposited using shorter wavelength lasers is largely attributed to the higher
absorption coefficients in YBa2Cu307.x targets with decreasing A. The absorption
coefficients are 1.2 x 10s, 1.5 x 10s, and 1.7 x 10s cm'1 at 1064, 532, and 355 nm,
respectively.57 A higher absorption coefficient results in a thinner layer at the
surface into which the laser energy is coupled, thus creating a hotter
plume with finer fragments. However, the improved microstructures which result
from short wavelength radiation cannot be completely attributed to the slightly
larger YBa2Cu307.x target absorption coefficients. Strong absorption by
photofragments in short wavelength radiation, and subsequent fragmentation into
smaller particles, probably contributes to the smooth morphology of films grown
at shorter wavelengths.57
A great deal of effort has been made to understand the mechanisms by which
material is transferred from a YBa2Cu307_x target to the substrate via laser


37
(2'38)
where Ed is the energy for desorption of an adatom from the substrate. The
mean distance an adatom diffuses before being desorbed is given by:
<2-39)
where is the single jump distance. This result shows that the surface diffusion
length increases with decreasing substrate temperature because the mean stay
time is increased. However, epitaxy is not necessarily improved by lowering the
substrate temperature because the jump frequency is lowered, hence the rate at
which equilibrium is reached is lowered. When the number of adatoms diffusing
across the surface at a given time approaches the density of surface sites, energy is
lost by adatom-adatom collisions, and adatom-substrate interactions are
diminished, thereby decreasing epitaxy. For the remainder of this study, we
assume that the flux of impinging species was low enough, and the substrate
temperature sufficiently high that the number of adatoms diffusing along the
surface was negligible compared to the number of surface sites, hence the
adatoms were able to reach their lowest energy configuration and equilibrium
conditions prevailed.
Until the nucleus reaches a critical size, it is more likely that the atoms will
dissociate and join a different nucleus or desorb, thus sub-critical nuclei do not
participate in the film growth process. In a real deposition system, the flux of


68
Devices fabricated on Si substrates would benefit from the lower attenuation and
signal distortion which superconducting interconnects can provide. Because
silicon and sapphire are chemically reactive with YBajQ^O^ and attempts to
grow YBa2Cu307.x films directly on these substrates result in interfacial phases
which damage both the substrate and film,80 substantial efforts to grow
intermediate barrier layers on the substrate prior to YBa2Cu307_x film deposition
have been made. In order to understand why some barrier layers are successful
and thus obtain the expertise necessary to design better barrier layer structures, it
is helpful to characterize YBa2Cu307_x films grown on inert single crystal
substrates. The first class of substrates are materials which are chemically inert to
YBa2Cu3O7.jp lattice match reasonably well, and have the perovskite structure.
SrTi03 is cubic with 2^ = 3.905 at 25 C,81 and LaA103 is rhombohedral with 2lq
= 7.586 and and included angle of 901'. However, LaA103 undergoes a cubic
rhombohedral transformation at 430 C, and at 650 C LaA103 is cubic with 2^ =
3.818 ,82 so in-situ films are grown on a cubic perovskite LaA103 substrate.
YBa2Cu307_x films grown at 650 C on (100) SrTi03 substrates show T0 at 89 K
and Jc = 7.5xl06 amps/cm2 at 77 K. Cross-sectional transmission electron
microscopy indicates the interface is atomically flat and abrupt, and the film does
not contain any secondary phases.83,84,85 The films grew with the < 001 > direction
normal to the substrate, the films are heavily faulted, with staggered,
discontinuous (001) layers. Close to the substrate, Y or Ba combine with Cu and
O to form perovskite subshells which are epitaxial with the substrate.


>H
Eh
<
m
S3
w
Q
Eh
S3
m

K
S3
u
l-J
o

Eh
P
o
0.0
YBa2Cua07_x/Zr02 ON
(100) SrTiOg
4.5 K
0.4 0.8 1.2 1.6
MAGNETIC FIELD (T)
2.0
Figure 5-5. Jc versus magnetic field intensity for a YBa2Cu307.x/Zr02 film
(100) SrTi03. Data taken at 4.5 K.


229
reacted with Si and formed pinholes as the film coalesced, thus providing open
channels for Ba to diffuse towards the substrate. Although there is no direct
proof of the existence of a Ba-Si-O phase in the superconducting film, it seem
likely that the same pathways which permitted diffusion of Ba to the substrate
(postulated to be pinholes) would allow Si diffusion into YBa2Cu307_x.
YAlQj Barrier Layers
YA103 films deposited onto (1102) A1203 were similar to LaA103 films in that
no crystalline peaks were observed. However, the electrical properties of
YBa2Cu307_x/YA103 films were far superior to YBa2Cu307_x/LaA103 films on
(1102) A1203. The normal state resistance values were more metallic, although
they were not ideally metallic, indicating grain boundary resistance was still
significant. The higher transition temperature and sharp drop to T0 at 82 K was
attributed to improved coupling between the YBa2Cu307_x grains, which shows
that excess Y in the YBa2Cu307.x film was less damaging than La. The extent to
which A1 was responsible for the slight drop in T0 is difficult to quantify. At 750
C, A1 diffuses into YBa2Cu307.x and degrades the superconductivity. Initial
experiments in which YBa2Cu307.x films were deposited directly onto A1203
substrates with no barrier layer produced white films which were almost
transparent, and there was no evidence of a superconducting phase. Clearly, A1
incorporation into YBa2Cu307.x films degrades the superconductivity. The slight
drop in T0 for YBa2Cu307.x/YA103 films deposited on (1102) A1203 probably


Jc was 2.5 X106 amps/cm2 at 4.5 K. The surface resistance was 102 ohms at 36
gigahertz.
On silicon substrates, YBa2Cu307.x degradation is aggrevated by thermal stresses
created by the difference in thermal expansion coefficients between YE^Q^O^ and
Si (13.2 versus 3.8 xlO^C, respectively), which causes microcracking in the
YBa2Cu307.x films. Cracking and interdiffusion were minimized by depositing a
YA103 barrier layer prior to YBa2Cu307_x. The thermal stresses were relieved by
viscoelastic relaxation in the YBa2Cu307_x film, and the T0 was 78 K.
The Jc values of YBa2Cu307.x films on Y-Zr02 substrates were increased by
depositing Y-Zr02 or Y203 barrier layers. YBa2Cu307.x/Y203 films on Y-Zr02
substrates had Jc values of 9x10s and lxlO7 amps/cm2 at 77 and 4.5 K. The Jc of
YBa2Cu307.x films deposited on a Y-Zr02 substrate without a barrier layer was
6.8 x 103 amps/cm2 at 4.5 K. The higher Jc values were attributed to pinning of the
magnetic flux by excess Y203 at high-grain boundaries.
vii


211
Estimate of Stress and Cracking Due to Differential Thermal Expansion
A rough analysis of the magnitude of the thermally induced stresses can be
obtained from the continuum mechanics model. The difference in thermal
expansion coefficients necessary to cause film cracking can be obtained by
inserting values128 of a = 50 MPa, E = 86 GPa, v = 0.290, and AT = 705 C for
the fracture strength, Young's modulus, Poisson's ratio, and temperature
difference into equation 4-57:
50jc106 pa (86x 109 Pa) (705) (Aa)
0.71
This analysis suggests that if Aa exceeds 5.9xlO7/C, thermally induced stresses
will cause cracking in the YBajQ^O,.* film upon cooling down from 750 C.
However, there are several factors which affect the validity of this arguement.
First, the fracture strength used in the calculation is for bulk YBa2Cu307.x, and the
film fracture strength is probably much higher. Second, the assumption that all of
the stress is relieved by either elastic deformation or cracking must be incorrect,
since uncracked YBa2Cu307_x films can be deposited on substrates such as SrTi03
and LaA103, for which the difference in thermal expansion coefficients exceeds
the calculated critical values. Since the melting temperature of YBa2Cu307.x is
1010 C, it is probable that atomic motion is relatively easy near the deposition
temperature, so plastic deformation is expected to be the primary stress reduction


AUGER SIGNAL
(peak-background) / (background)
en


201
YBa2Cu307.x/Y-Zr02 film on Si by 0.061, the two curves are very similar. This
indicates that both films have about the same intragranular and grain boundary
resistance, but the percolation parameter (p) was much larger for the film
deposited on Si. This is consistent with the observation of cracks in this film,
which increase the length of current pathways and thus increase the normal state
resistances. The degradation of T0 in this film was attributed to poor coupling of
the superconducting charge carriers across the cracked regions. Even though
there appear to be pathways by which the superconducting holes could travel,
microcracking and dislocations not imaged by the SEM will reduce the
superconducting order parameter in these regions. The resistance versus
temperature curve for YI^CujO-^ deposited on single crystal Y-Zr02 suggests
the grain boundary resistance was at least as high as the intragranular resistance,
and the low Jc (6.8 xlO3 A/cm2 at 4.5 K) was probably due to high angle grain
boundaries, which resulted in poor intergranular coupling of the superconducting
holes at the grain boundaries.
Effect of Y-ZrQ3.Y3Q3. and ZrO-, Barrier Layers on Jc Values
Figure 5-1 shows Jc versus magnetic field intensity for YBa2Cu307.x/
Y0.i3Zr087O194 films deposited on Y-Zr02 and LaA103 substrates. The Jc values
are roughly proportional to If0-5, indicating the Jc was limited by flux creep within
the grains.34 Similar plots for YBa2Cu307.x/Y203 films deposited on Y-Zr02 and


98
in the Cu(l)-0(1) chains reduce the frequency of the Cu(l)-0(4) vibration. In
addition, substitution of Cu(l) by impurity atoms (such as Si or Al) broaden the
500 cm'1 peak. Qualitative information about the extent of oxygenation is also
obtained from the height of the 335 cm'1 peak. In fully oxygenated YBajQ^O?.*,
repulsive interactions between the 0(2) and 0(4) atoms reduce the peak height,
and vacancies on the 0(4) sites slightly increase the peak intensity. The shape of
the 335 cm'1 peak also provides qualitative information about the oxygen content
of the sample. Several studies have shown that the peak is highly asymmetric for
superconducting films and bulk samples, even at room temperature.120,121 This
asymmetry is caused by Fano-type interference between the phonon mode and the
electronic continuum in the Cu-O planes. The Fano-type lineshape122 is caused by
reduced electronic scattering at energies slightly above 335 cm'1, hence the
peak/background ratio is higher at energies greater than 335 cm'1. In oxygen
depleted, non-superconducting YBa2Cu307.x the 335 cm'1 peak is symmetric, which
has been attributed to the reduction in the free carrier concentration in the Cu-0
planes.
For completely (OOl)-oriented films, the peak at 500 cm'1 is not observed,
since the electric vector of the incident radiation is perpendicular to the <001>
direction; thus the ratio of the 335 cm4/500 cm'1 peaks can be used to
qualitatively determine the ratios of (001) to (100) or (OlO)-oriented grains. The
Raman sensitivity for the 440 cm'1 peak is highest in the <001> direction, so the
appearence of this peak indicates that a significant fraction of the grains have


INTENSITY (arbitrary units)
135
YBa2Cu307.x/Y-Zr02
on A!203 (0001)
o
o
o
in
o
o
<0
o
o
co
s
CM
O
O

JL
J
o
o
r
JL
o
8
UU
&
A
A
Y-Zr02 on
Al203 (0001)
o
o
CM
YBa2Cu307.x
Y-Zr02
10
20 30 40 50
DIFFRACTION ANGLE (20)
60
Figure 4-29. X-ray diffraction patterns from a Y-ZrO, and YBaoCuX^.,./Y-ZrO,
film on (0001) A1203.


30
YBa2Cu307.x films grown on (100) SrTi03 substrates, where two substrates were
sintered together so as to deliberately produce a misorientation angle in the
film yet insure the <001> directions in the YBajQ^O^ film on both sides of the
sintered SrTi03 junction were parallel, variations in Jc as a function of grain
misorientation showed a dramatic drop as the grain boundary angle was increased.
Inside the grains, Jc values of 4xl06 amps/cm2 at 5 K were measured. However,
across the grain boundaries Jc values dropped from 4x 106 amps/cm2 at 0
misorientation to O.lxlO6 amps/cm2 at 12.5 misorientation.
Further increases in resulted in similar or lower Jc values, and did not
follow a systematic trend. The authors observed that J^/Jc8 was approximately
proportional to 1/ for < 20, and proposed a model in which the dislocation
spacing was the predominant factor in determining Jc. The dislocation array acts
as a partial barrier to superconducting electrons, or perhaps as an easy path for
flux flow, since the superconducting order parameter is depressed at the
dislocation cores. Although various studies have observed significantly larger Jc
values (> 5xl06 amps/cm2 at 77 K) in epitaxial (OOl)-oriented films as opposed
to polycrystalline films, this study was the first to indicate that larger Jc would be
observed in (001) films with in-plane epitaxy and low grain boundary angles,
compared to to (001) oriented films with random in-plane orientation.
Later experiments by Mannhart et al. on epitaxial, (OOl)-oriented YBa2Cu307_x
films grown on (100) SrTi03 substrates determined that the values of
intergranular Jc as a function of temperature could be fitted by the
Ambegtaokatar-Baratoff equation for a SIS junction in which an energy gap A(0)


220
SURFACE STRUCTURE OF (100) SrTiOg
Q = Oxygen
0 = Titanium
^ = Strontium
Ti-q* LAYER
0.390 nm ; ^
>
0.390 nm
SrOxLAYER
0.390 nm
0.390 nm
o o o o o o
000*000
;o:o:o:o:o:o
o o o o o o
k-
Figure 5-7. Surface structure of (100) SrTi03.


69
Perpendicular to the substrate, each perovskite subshell contains both Y and Ba,
but Y and Ba segregate into different domains along the interface. Away from
the interface, the (001) layers become continuous, but very wavy. Increasing the
substrate temperature to 720 780 C reduced the number of defects. Films
grown at the higher temperatures were completely (001) oriented, with less wavy
(001) layers, and Jc values of 2.2 xlO6 amps/cm2 at 77 K.
YBa2Cu307.x films with T0 = 90 K and Jc = lxlO6 amps/cm2 have been
deposited on (1102) LaA103 at 750 C.86 These films were highly (001) oriented
during the first 4000 of their growth. At greater thicknesses however, the
growth mode abruptly changed, and the dominant growth mode was with the
(100) and (010) planes parallel to the substrate. Computer simulations based on
nucleation densities and growth rates of (001) and (lOO)-oriented grains help
explain the experimentally observed film microstructures.87 Assuming the
nucleation density of (OOl)-oriented grains is 109/cm2, and for (100) and (010)
oriented grains is 2xl08/cm2, and the growth rate in the <100> and <010>
directions is 10 times larger than in the <001> direction, a microstructure
emerges in which the film surface is initially dominated by (OOl)-oriented grains.
As the film thickens, the (100) and (OlO)-oriented grains, which have a high
growth rate normal to the substrate, coalesce and the film is covered by these
orientations.
YBa2Cu307.x films have also been deposited on several substrates which have
distorted perovskite structures. The difference between these substrates and the


158
. +' * V** 'T
*

-f/ki* > J*;
# * .
< > *, 0 m
*'%.* * * :m
V- *Z ^ ^

V* ^#V.V t
#
v

/**
15KV X6000
8335 1.0T UFMSE
Figure 4-48. Scanning electron micrograph of a YBa2Cu307.x/LaA103 film on
(100) Si.


CHAPTER 1
INTRODUCTION
There are several commercial and military application for thin films which are
superconducting at temperatures above 77 K. Since most semiconducting
devices perform optimally near 77 K, the potential for high-speed devices with
low-attenuation superconducting interconnects is promising1. In another
electronics area, passive microwave and millimeter-wave devices patterned into
superconducting thin films dramatically outperform the resonators and filters
presently being used.2 As the technology for depositing superconducting films on
A1203 substrates improves, this performance will be further enhanced. Potentially,
the largest applications for superconducting films are in power applications.3
Transmission cables, motors, and high field magnets which are too costly to
operate at 4.5 K would be economically feasible at 77 K. Progress in
fabricating superconducting wires has been difficult because the wires currently
being fabricated suffer from brittleness and low critical current density (Jc) values
(3X104 amps/cm2 at 77 K). A better technique for making superconducting
cables may be to deposit superconducting films on fibers such as yittria-stabilized
zirconia (Y-Zr02) which are mechanically strong and have a similar thermal
expansion coefficient as the superconducting film.
1


9
I
I
I
I
Figure 2-3. Schematic of a two-junction SQUID. Reference 4.
Because Josephson junctions can be changed from superconducting to
semiconducting and vice-versa by varying the current or magnetic field, they have
been used as switching devices in digital electronics.9 The intrinsic switching
speed of a Josephson junction is:


50
Several features regarding the stress distribution in a multilayered film are
implicit in the continuum mechanics model. Since the oftf force for each layer of
film is significantly less than adopt the surface dimensions of the relaxed substrate if there is no plastic
deformation at the film-substrate interface, or between any of the film layers. A
consequence of the small ot force of the film relative to that of the substrate is
that the stresses induced in each layer result almost entirely from interactions
between the film and substrate. According to this model, the stress in any given
layer of a multilayered film is independent of the sequence in which the films
were deposited, and the thermally induced stresses in each of the layers is given
by:46
J4oirat)T (2-57)
where Ef and vf are the Young's modulus and Poisson's ratio for the film, a{ and
as are the thermal expansion coefficients for the film and substrate, and AT is the
change in temperature to which the film-substrate combination is subjected. A
key assumption of this model is that the chemical composition of each layer is
uniform. Since the composition is uniform, the thermal expansion coefficient
within each layer is constant, so the dimensions of the top of each layer will be
the same as the bottom. This assumption, coupled with the assumption of no
plastic deformation at any of the interfaces, leads to the conclusion that it is not
possible to reduce the thermally induced stresses in the outermost film by


FLOW CHART FOR Y Al Si O
L <** Al6Si2013 + AI2
(1912 C)
)3 L ** Y203 + Y4A1209
(1940 C)
L Y4A1209 + YA103 -
(1875 C)
L YA103 + Y3A15012
(1850 C)
L * Y2Si05 + Y4Si3012
(1900 C)
-L ~ Y203 + Y2Si05
(1800 C)
L + Y203 Y4A1209 + Y2Si05
YA103 Y3A15012 Y4Si3012
L Y4A1209 Y2Si05
L + Y4A12Q9 > Y2SiOs + YA1Q3
Y4A1209 Y2Si05 YA103
L Y2Si05 YAJ03
Figure A-l. Flow chart of the Y Al Si O system.


19
(2-18)
e ** +1
where f(E) is the probability that a given electron state is occupied, E is the
energy, Ep is the Fermi energy, k is Boltzman's constant, and T is the
temperature. In normal metals, electrical conductivity is possible because there
are empty electronic states at energies greater than Ep, into which electrons can
hop and thus move through the lattice. Current is transported through a material
because the applied electric field raises the electron energy level of one terminal
relative to the other, thus increasing the probability that electrons will hop from
the high potential to the lower potential region of the sample. Electrical
resistance arises from processes which transfer kinetic energy to the crystal lattice
by electron-electron and electron-phonon collisions.
The mechanisms for current transport in superconductors are different from
those observed in normal metals. Superconductivity is a cooperative phenomena
involving many electron or hole pairs, and is possible because of electron (hole)-
phonon coupling, which creates an attractive force between electrons (holes). The
potential energy caused by an electronic transition from an initial state k, to
another state k', is given by:6


14
and
i =
1 + art
2_2
(2-11)
o<**
1 + (2-12)
ax and o2 are the conduction and displacement currents, and x is the mean time
between electron-phonon interactions. In normal metals, the skin depth (<5) is the
distance into the metal which an electric field can penetrate, and is a function of
the electrical conductivity:
v T7 l
(2uo0pw)-5
where c is the speed of light (3xl010 cm/sec) and fi is the magnetic permeability.
The dependence of the skin depth on electrical conductivity and hence
frequency means the dielectric constant of the metal changes with frequency. The
parameters pertinent to microwave signal transmission are contained within the
propagation constant, y,14 where
y = a +j p (2-14)
a is the attenuation constant, and is the wavenumber of the microstrip line. /?
is given by:


109
directly onto Y-Zr02 substrates, although there was a wider AT between Tonset
and T0 in the YBa2Cu307.x/Y-Zr02 film on Si. AES data (figure 4-10) shows that
Y-Zr02 was an effective barrier layer for preventing interdiffusion between
YBa2Cu307.x and the Si substrate. The x-ray diffraction patterns obtained from
YBa2Cu307.x/Y-Zr02 on Si and YBa2Cu307.x on Y-Zr02 are similar, with textured
(001) YBa2Cu307.x predominant on both substrates (figure 4-11). There were,
however, small (100) diffraction peaks observed in the YBa2Cu307.x film on
Y-Zr02. Critical current density measuerements taken at 4.2 K show that the Jc
for this film was 6.8x10 3 amps/cm2.
p
v* -
x I
V ,
, *
% ~
. > <
V. ; *
w )
20Kt> XI2000 $060
- t
, 1 ID"UFMSE
Figure 4-7. Scanning electron micrograph of a YBa2Cu307.x/Y-Zr02 film on (100)
Si.


LaAlOj Barrier Layers
In contrast to SrTi03 films, LaA103 barrier layers deposited on (1102) A1203
did not provide a suitable template on which to grow YBa2Cu307_x films. Despite
having a smaller lattice mismatch with (102) A1203 than SrTi03, and having a
similar crystal structure (LaA103 is a distorted pervoskite), x-ray diffraction of a
LaA103 film deposited on (1102) A1203 did not show any LaA103 peaks. The
YBa2Cu307.x/LaA103 film on (1102) A1203 was (001) oriented, but the x-ray
diffraction peaks were less intense than the (001) YBa2Cu307_x/(100) SrTi03 film
grown on (1102) A1203. The normal state resistance was slightly metallic, and the
film did not become superconducting until 58 K. The non-ideal normal state
resistance and depressed T0 were attributed to grain boundaries and lack of
completely crystallized YBa2Cu307.x phase in the film. Interpretation of the
resistance vs. temperature behavior in terms of the Halbritter equation showed
the intragranular resistance was high, since the curve did not extrapolate close to
0 ohms at 0 K; this was partially due to the presence of non-(001) oriented
grains. Raman spectra showed a large peak at 500 cm'1, which indicates that a
significant portion of the film is oriented with the <001> direction in the plane of
the film, or at some angle with the surface other than 90. This is consistent with
x-ray diffraction data. The increased thermal stresses in non-(001) oriented
YBa2Cu307_x grains may have induced microcracking.


CRITICAL CURRENT DENSITY (A/cm2)
MAGNETIC FIELD (T)
Figure 4-14. Critical current density versus magnetic field intensity for a YBa2Cu307.x/Y-Zr02 film on Y-Zr02. Data
taken at temperatures from 4.5 to 77 K.
00


L + Y3A15012 ** Y4Si3012 + A1203
L + Y3A15012 * Y2Cu205 + A1203
L + Y3A15012 ** Y4Si30j2 + Y2Cu205 + A1203
L Y4Si3012 Y2Cu205 A1203
L + Y4Si3012 ** Y2Si207 + Y2Cu205 j
L + Y4Si3012 + Y2Cu205 ** Y2Cu205 + A1203
L Y4Si3012 Y2Si20- A1203
L Y2Cu205 Y2Si207 A1203
L + A1203 ** Al6Si2013 + Y4Si3012
1
L + Y4Si3012 ** Al6Si2013 + Y2Si207
L + Y4Si3012 ** A1203 + Y2Si207 + Al^Si2013
L AJ203 Y2Si207 Al6Si2013
L + Y2Si207 *> Si02 + Y2Cu205 i
L + Y2Cu205 + Y2Si207 ** Si02 + A1203
L Y2Cu205 Si02 AI,0
2^3
L Y2Si207 Si02 A1203
Figure A-5. Continued.
S)
00


CHAPTER 5
DISCUSSION
A number of barrier layer/substrate combinations were examined in order to
what the most critical parameters were for optimal quality YBa2Cu307_x films on
the various substrates. These barrier layer/substrate combinations, and the zero
resistance temperatures and critical current densities observed in the
superconducting films are summarized in table 5-1.
Intergranular Versus Intragranular Effects
YBa2Cu307.x films with nearly ideal normal state and superconducting
properties were deposited on (1102) LaA103 substrates. Figure 4-1 shows that
the normal state resistance behavior for the fully oxygenated film was metallic,
with a sharp drop to 0 ohms at 88 K. Using the Halbritter equation (2-34) to
separate the resistance values into inter and intragranular contributions, we see
that the intergranular resistance was negligible compared to the intragranular
contributions, since the R vs. T curve extrapolates to 0 ohms at 0 K. The film
which was cooled to 550 C before filling the chamber with oxygen does not show
ideal R vs. T behavior, and the temperature at which the resistance became zero
was depressed to 79 K. This degradation was attributed to an incomplete
transition from the tetragonal to the orthorhombic crystal structure during cooling,
196


38
atoms striking the substrate consists of clusters of atoms as well as single atoms or
ions. Although it is difficult to determine what is the smallest stable nucleus size,
it is logical to expect that larger clusters striking the surface are more likely to
remain and form stable nuclei than are smaller clusters or atoms, since there are
more overlayer-substrate bonds formed. In addition, increasing the substrate
temperature is expected to increase the critical nucleus size, so the distribution of
clusters or atoms which eventually form stable clusters is further skewed towards
the higher cluster sizes.33
The relationship between critical cluster size and substrate temperature is
important because the orientation of the entire film can be heavily influenced by
the size of the original clusters from which stable nuclei are grown. Theoretical
and experimental data on face centered cubic metallic films deposited on NaCl
substrates indicates that for a critical cluster size of 3, the atoms will adopt a
triangular arrangement, and the film will grow with a (111) orientation.34
Similarily, when the smallest critical nucleus size is 4 atoms, the nucleus will
arrange itself in a square or rectangular mesh, and (100) oriented growth is
favored. In general, the first monolayer will grow with the surface mesh initiated
by the stable cluster, and subsequent layers will maintain this orientation and
grow with the closest packed planes parallel to the substrate. Based on this
analysis, face centered cubic films in which the critical cluster sizes are 3 or 4
atoms are expected to grow with (111) and (100) orientations, respectively. It has
been observed that several metallic films which grow with a (111) orientation at


89
this set of experiments the beam accelerating voltage was always 5 kilovolts
because the sensitivity of the YLMM transition is highest at this accelerating
voltage.
Atoms can be ionized by backscattered and secondary electrons as well as
primary electrons, and the backscattering and secondary-electron yields are
sensitive to the chemical environment and electrical properties of the sample.119
Discrepancies in ionization cross section resulting from different backscattering
yields of the same element in different compounds is the primary phenomena
which limits quantitative Auger analysis. In this set of experiments, interdiffusion
betwen YBa2Cu307.x, the barrier layers, and the Si substrates were of interest.
Since the chemical environment and backscattering yields of each of the structures
were similar, semi-quantitative comparisons between the chemical compositions of
these structures could be made. Two final parameters which significantly effect
the Auger yield are the angle between the sample surface and the incident
electron beam, and the angle between the surface and detector. As the angle
between the beam and surface normal increases, the incident beam path length in
the surface region increases by a factor of sec 0,114 and the Auger yield increases.
Experimentally, this factor was kept constant by always keeping the angle between
the incident beam and surface normal at 60 .


117
metallic, and extrapolated closer to zero ohms at 0 K than did the YBa2Cu307.x
film grown directly on the Y-ZrOz substrate. The more metallic behavior with a
barrier layer indicates lower grain boundary resistances than for films deposited
directly on Y-Zr02 substrates.
The Jc values of YBa2Cu307..x/Y-Zr02 films deposited on Y-Zr02 were close
to the best values reported for YBa2Cu307.x films deposited on (100) oriented Y-
Zr02.104 Figure 4-14 shows Jc versus H data taken at temperatures 77 to 4.5 K.
At zero field (0 tesla), the Jc ranges from 9 x10s A/cm2 at 77 K to 1.5 xlO7
A/cm2 at 4.5 K.
No x-ray diffraction peaks were detected from the Y-ZrOz barrier layer film
deposited on a randomly oriented Y-Zr02 substrate (figure 4-15). The diffraction
pattern from the Y-ZrOz film on the Y-Zr02 substrate was identical to that
obtained from a Y-ZrOz substrate. X-ray diffraction from the YBa2Cu307.x/Y-
Zr02 film on Y-ZrOz shows that the YBa2Cu307.x film was highly (001) oriented,
since only peaks from this orientation were detected. SEM indicates the film is
comprised of 0.3 -1.0 /m grains (figure 4-16), and the grain boundaries are much
clearer in this picture than in the YBajCu^^ film on LaA103 (figure 4-4).
High Jc values were also observed for YBa2Cu307.x/Y-Zr02 films deposited on
(1102) LaA103. The Jc of this film (3XlO7 amps/cm2) at 4.5 K was close to the
best values we have obtained in YBa2Cu307_x films deposited directly on (1102)
LaA103 substrates (5 x 107 A/cm2). Resistance versus temperature data are shown


51
depositing intermediate layers in which the thermal expansion coefficients
gradually bridge the gap between the substrate and outermost film.
Film stresses can significantly affect microstructural features such as grain size
and porosity in thin films. Using an energy minimization argument, Chaudhari50
showed that tensile stresses can significantly reduce the average grain size of a
film. Because the atomic packing density is lower in a region containing grain
boundaries than in a crystalline region, tensile stresses are reduced in films with
small grain sizes. A quantitative energy minima for the film is given by the
expression:
strain boundary
'boundary
(2-58)
where do is the initial grain size, d is the final grain size, e0 is the initial strain in
the film, and a is a normalized distance parameter used to compare the atomic
density of the grain boundary region with that of a grain. If the boundary region
has the same atomic density as the grain, then a is zero. If the boundary has a
monolayer of atoms missing, then a is 1. /? is a geometrical factor used to
characterize the shape of the grains. For grains with a square cross-section, ¡5 =
2. y is the grain boundary energy. For films subjected to tensile stresses, there is
some combination of strain and grain boundary energies for which the overall
energy of the film is a minimum. For films subjected to compressive stresses,
there is no overall energy minima, and the total film energy decreases as the grain
size increases.


Rs (OHMS)
146
Figure 4-38. Surface resistance of a YBa2Cu307.x/SrTi03 film deposited on
(1102) A1203. Data taken at 36 gigahertz.


RESISTANCE (OHMS)
181
Figure 4-67. Resistance versus temperature data for a YBa2Cu307_x/Y203 film on
(100) Si.


203
from the Y203 barrier layer or Y-Zr02 substrate were not observed. The Y-Zr02
film deposited on (100) LaA103 was primarily (100) oriented, and a smaller (220)
peak was also observed. The presence of both peaks usually creates a
YBa2Cu307.x films with mixed in-plane orientations, and high-angle grain
boundaries. The high Jc values of this film indicate good coupling of the
superconducting holes across the grain boundaries.
MAGNETIC FIELD (T)
Figure 5-1. Jc versus magnetic field intensity for YBa2Cu307_x/Y-Zr02 films
deposited on Y-Zr02 and (1102) LaA103 substrates. Data taken at
. 4.5 K. The solid and dashed lines are fits of the experimental data
with the flux creep model


185
films, a more complete description of the mechanisms by which thermally induced
stresses were transferred through the barrier layers was obtained by comparing
the resistance values of YBa2Cu307.x/barrier layer films grown on (1102) A1203
with or without an Al6Si2013 layer. A complicating factor was that Al6Si2013 was a
source of Si, which favors formation of the tetragonal phase when it diffuses into
YBa2Cu3O7.jp and thereby reduces the transition temperature. Previous data
showed the presence of tetragonal YBa2Cu307_x will lower T0 (see figure 4-1), but
only increase the normal state resistance values slightly. On the other hand, the
resistance values of cracked YBa2Cu307.x films are about a factor of 10 higher
than those of uncracked films. This study was useful for distinguishing whether
YBa2Cu307.x degradation was caused by tensile cracking, or Si diffusion into the
YBa2Cu307.x film.
Addition of Al6Si2013 into the barrier layer structures significantly altered the
electrical performance and microstructures of YBa2Cu307.x films deposited on Si
substrates. Resistance versus temperature data (figure 4-71) show the normal
state resistances for both the YBa2Cu307_x/YA103/Al6Si2013 and YBa2Cu307.x/
LaA103/Al6Si2013 films on Si are metallic, but the resistance values for the film
with a LaA103 layer are less than one-half of those observed in the film with a
YA103 layer. The transition temperatures were very similar, with T0 = 61.5 and
59 K for the samples with a YA103 or LaA103 barrier layer, respectively. SEM
data (figures 4-72 and 4-73) show the morphologies of both films are virtually
identical, with unconnected crack networks running through the films. The x-ray


RESISTANCE (OHMS)
155
Figure 4-45. Resistance versus temperature for a YBa2Cu307.x/LaA103 film on
(100) Si.


210
oriented.93 In addition, Raman spectra from YBa2Cu307.x/Y-Zr02 films deposited
on (1102) and (0001) A1203 show the peak at 500 cm'1 is small, which suggests
both YBa2Cu307.x films were highly (001) oriented; the non-symmetric lineshape
of the 335 cm'1 peaks indicate the films were highly oxygenated. The explanation
for this apparent anomaly is given by the SEM micrograph of YBa2Cu307_x/Y-
Zr02 on (0001) A1203 (figure 4-32). This picture shows a faint series of cracks in
the film, and the degradation of the electrical and superconducting properties are
attributed to these cracks. Like the YBajQ^O^/Y-Zr02 films on Si, these cracks
form because of the difference in thermal expansion coefficients between
YBa2Cu307_x and the substrate. The thermal expansion coefficient of A1203 is
anisotropic (table 4-1), with the lowest values observed in directions parallel to
the (0001) plane, and the largest values being perpendicular to (0001). Thus we
conclude that for YBa2Cu307.x/Y-Zr02 films on A1203, the thermally induced
stresses are largest in YBajCujO^ films deposited on the (0001) plane. This
accounts for the cracking and degradation of electrical properties in these films.
In general, film cracking dramatically increased the normal state electrical
resistances of YBa2Cu307.x films. Since the electrical resistance across a crack is
not very sensitive to changes in the temperature, resistance versus temperature
profiles of YBa2Cu307_x films were less metallic than those of the uncracked films,
and were sometimes semiconducting. For cases in which cracking was the
dominant defect, any excess resistance caused by intragranular damage was
masked by the cracking.


INTENSITY (ARBITRARY UNITS)
124
Figure 4-19. X-ray diffraction from Y-Zr02 and YBa^O? J Y-ZrO, films
deposited on (1102) LaA103.


167
Figure 4-55. Cross-sectional view of YBa2Cu307.x/YA103 film on (100) Si, taken
with SEM.


245
likely to occur in this system were inferred from the known YxQiyOz compounds.
The quaternary flow chart is shown in figure A5. This chart chart lists the four-
phase equilibria (determined in figures Al A4) in the columns marked "ternary",
and the five-phase equilibria in the "quaternary" column.


17
the superconducting penetration depth is much less than the normal metal skin
depth, and it is only at (o > 1012 Hz that attenuation in a superconducting
stripline approaches a for normal metals (figure 2-6).16,17
In semiconductor devices, the reduced attenuation and Joule heating provided
by superconducting interconnect lines would permit interconnect lines to be scaled
down considerably, thereby reducing chip to chip propagation delays.
Figure 2-6. Surface resistances of YBajQigOy.* (dotted line) and copper (solid
line) films deposited on LaA103 substrates. Reference 17.