study of structural and magnetic properties of
TRANSCRIPT
STUDY OF STRUCTURAL AND MAGNETIC PROPERTIES OF
INTERMETALLIC THIN FILMS
by
EZHIL ARASAN MANOHARAN
GARY MANKEY, COMMITTEE CHAIR
PATRICK LECLAIR
RAINER SCHAD
DEAN TOWNSLEY
RAMANA REDDY
A DISSERTATION
Submitted in partial fulfillment of the requirements
for the degree of Doctor of Philosophy
in the Department of Physics
in the Graduate School of
The University of Alabama
TUSCALOOSA, ALABAMA
2016
Copyright Ezhil A. Manoharan 2016
ALL RIGHTS RESERVED
ii
ABSTRACT
Intermetallic thin films have tunable magnetic properties. The magnetic phases of
intermetallic thin films were tuned by changing the alloy composition of the intermetallic
system. L10 Fe50Pt50 thin film has high magnetic anisotropy which makes them ideal candidates
for the thin film recording media. Magnetic phases of Fe50Pt50 can be tuned by the addition of
third element like Mn by forming Fe50-x Mnx Pt50 ternary alloy system. In this work magnetic
phases of ordered Fe rich Fe50-xMnxPt50 and Mn rich Fe50-xMnxPt50 thin films of Fe50-x Mnx Pt50
alloy system is investigated. Fe rich Fe50-xMnxPt50 thin films are epitaxially grown on a- Al2O3
and MgO (100) substrates, while Mn rich Fe50-xMnxPt50 thin films are grown on MgO (100)
substrates. The change in the magnetic properties in Fe rich Fe50-xMnxPt50 thin films due to
presence of tetragonal phase and the prediction of a the presence of a new low temperature phase
in the Mn rich Fe50-xMnxPt50 thin films is verified. These intermetallic films are produced in a
Ultra High Vacuum sputtering system with Reflective High Energy Electron Diffraction and
Auger electron spectroscopy. RHEED is used to verify epitaxy and Auger electron spectroscopy
measures chemical composition.
iii
DEDICATION
This dissertation is dedicated to everyone who helped me and guided me through the
trials and tribulations of creating this manuscript. In particular, my family and close friends who
stood by me throughout the time taken to complete this masterpiece.
iv
LIST OF ABBREVIATIONS AND SYMBOLS
AFM Antiferromagnet
FM Ferromagnet
F/C Mixed ferromagnetic and antiferromagnetic phase
G/C Non collinear mixed antiferromagnetic phase
F/G Non collinear mixed ferromagnetic and antiferromagnetic phase
RGA Residual gas analyzer
UHV Ultra high vacuum
XRD X-ray diffraction
XRR X-ray reflectivity
RHEED Reflection high energy electron diffraction
VSM Vibrating sample magnetometer
PPMS Physical property measurement system
CMA Cylindrical mirror analyzer
SQUID Super conducting quantum interference device
v
ACKNOWLEDGMENTS
I thank all the people who were involved in my Ph.D. years. I thank Dr. Mankey for
teaching me vacuum science. I am really grateful for this opportunity in the University of
Alabama. I thank my committee members for believing in me. I would like to take thank physics
machine shop people Joe Howell, David Key, Danny Whitcomb and Jason Kyukendall for their
help in building and modification of the vacuum system. I also thank Dr. Patrick LeClair for his
great advice and support. I thank Amit Singh for helping me during PPMS measurements.
vi
CONTENTS
ABSTRACT .......................................................................................................................................... ii
DEDICATION ...................................................................................................................................... iii
LIST OF ABBREVIATIONS AND SYMBOLS ......................................................... iv
ACKNOWLEDGMENTS .............................................................................................................. v
LIST OF TABLES .......................................................................................................................... viii
LIST OF FIGURES ........................................................................................................................... ix
CHAPTER 1 INTRODUCTION ............................................................................................... 1
CHAPTER 2 EXPERIMENTAL TECHNIQUES ...................................................... ..4
I. Auger Spectroscopy ........................................................................................................................ 4
II. Cylindrical Mirror Analyzer ................................................................................................ 10
III. Reflective High Energy Electron diffraction ......................................................... 17
IV. Sputtering ........................................................................................................................................ 20
V. Substrate Cleaning ..................................................................................................................... .29
VI. Flux Calibration .......................................................................................................................... 32
VII. X-ray Diffraction ..................................................................................................................... 37
CHAPTER 3 FE50PT45Rh5 FILMS ......................................................................................... 43
CHAPTER 4 FE25PT75 ..................................................................................................................... 56
CHAPTER 5 FE RICH Fe50-X MNXPT50 ............................................................................ 64
CHAPTER 6 MN RICH FEe50-X MnXPT50 ..................................................................... . 80
vii
CHAPTER 7 CONCLUSION .................................................................................................... 94
REFERENCES ...................................................................................................................................... 96
viii
LIST OF TABLES
I. Angles and intensities for (111) texture during Cr diffusion for the region I ...... 47
II. Angles and intensities for the region II during Cr diffusion .......................................... 48
III. Angles and intensities for the region III during Cr diffusion ...................................... 48
IV. Angles and intensities for the region IV during Cr diffusion ...................................... 49
V. Angles and intensities for the region V during Cr diffusion ........................................ 49
VI. Angles and intensities for the region VI during Cr diffusion ....................................... 49
VII. Angles and intensities for the region VII during Cr diffusion ................................... 50
VIII. Angles and intensities for (111) texture during no Cr diffusion
for the region I ............................................................................................................................................ 54
IX. Order parameter S for various compositions of Fe50-xMnxPt50 thin films
on a-plane sapphire .................................................................................................................................... .69
X. Composition, out of plane and in plane lattice parameters, c/a ratio,
saturationmagnetization and coercivity of Fe50-xMnxPt50 films ..................................... 77
XI. Expected moment and the observed moment for the Mn rich
Fe50-xMnxPt50 for different compositions ...................................................................................... 93
ix
LIST OF FIGURES
1.1 Calculated phase diagram of Fe50-xMnxPt ...................................................................... ...... 2
1.2. Spin diagrams for the different magnetic phases ..................................................... ...... 2
2.1 Schematic of Auger process ..................................................................................................... ...... 5
2.2 Schematic of Auger system used by Harris .................................................................. ...... 7
2.3 Auger of alloy steel reported by Harris ............................................................................ ...... 8
2.4 CMA developed by Palmerberg, Bohn and Tracy ................................................... ...... 9
2.5 CMA ......................................................................................................................................................... ...... 10
2.6 Inelastic mean free path of electron .................................................................................... ...... 13
2.7 Calculated Intensities of Fe/Ni bilayer ............................................................................. ...... 15
2.8 Influence of inter mixing of Ni/Nb bilayer at different temperatures ........ ...... 16
2.9 RHEED pattern for a GaAs (100) substrate (left)
and Cr deposited on GaAs substrate at 550 Co (left) ............................................... ...... 18
2.10 RHEED pattern for (left) the sapphire (1-120) and (right) 6 nm Cr /
14 nm Pt deposited on the sapphire substrate ............................................................ ...... 18
2.11 RHEED pattern for 30 nm polycrystalline Fe50 Pt45 Rh5 thin film............ ...... 19
2.12 Schematic of magnetron sputtering ................................................................................. ...... 20
2.13 Deposition rate vs. sputtering power .............................................................................. ...... 22
2.14 Schematic of “RASCAL” sputtering system ............................................................ ...... 23
2.15 Pressure versus time during the baking process ............................................................ 25
2.16 Mass spectrum of unbaked system .......................................................................................... 26
x
2.17 Mass spectrum of baked system ......................................................................................... ...... 27
2.18 Sample heater and sample holder assembly ............................................................. ...... 28
2.19 Presence of carbon in single crystal sapphire (1-120) ........................................ ...... 30
2.20 Sapphire substrate after 5 mins of sputter cleaning at 0.5 mTorr
at 0.5 kV beam voltage .................................................................................................................... 31
2.21 Kiessig plot of Cr deposited for 200 seconds
at 4 mTorr Ar pressure .............................................................................................................. ...... 34
2.22 XRR scan of a FePt on Si thin film for thickness calibration
purposes(a) and linear fit to obtain film thickness (b) ............................................... 35
2.23 XRR scan of a MnPt on Si thin film for thickness calibration
purposes (a) and linear fit to obtain film thickness (b) ............................................... 36
2.24 X-ray diffraction ............................................................................................................................. ...... 37
2.25 Definition of angles in the X-ray goniometer ........................................................... ...... 38
2.26 (a) FCC crystal structure (b) L10 crystal structure ................................................ ...... 39
2.27 Atomic factor for different metals .................................................................................... ...... 40
3.1 X-ray diffraction spectra of a-plane Al2O3/ 6 nm Cr/ 14 nm Pt/
30 nm Fe50Pt45Rh5 (111)Intensities of Fe/Ni bilayer ..................................................... 44
3.2 Auger spectra of Al2O3/ 6 nm Cr/ 14 nm Pt/30 nm
Fe50Pt45Rh5 (111) shows the presence of Cr diffusion .......................................... ...... 45
3.3 RHEED shows the polycrystalline Al2O3/ 6 nm Cr/ 14 nm Pt/
30 nm Fe50Pt45Rh5 thin film....................................................................................................... ...... 46
3.4 Pole figure of Al2O3/ 6 nm Cr/ 14 nm Pt/30 nm Fe50Pt45Rh5 (111) ..................... 47
3.5 Anomalous interfacial magnetization in a single layer
and trilayer FeRhPd films bilayer ................................................................................................ 51
3.6 Auger spectra of Al2O3/ 3 nm Cr/ 12 nm Pt/30 nm Fe50Pt45Rh5 (111)
shows the absence of Cr diffusion ....................................................................................... ...... 52
3.7 RHEED of Al2O3/ 3 nm Cr/ 12 nm Pt/30 nm Fe50Pt45Rh5 (111) ................... ...... 52
xi
3.8 Phi scan of Al2O3/ 3 nm Cr/ 12 nm Pt/30 nm Fe50Pt45Rh5(111)
shows the six fold symmetry ............................................................................................................. 53
3.9 Pole figure of Al2O3/ 3 nm Cr/ 12 nm Pt/30 nm
Fe50Pt45Rh5 (111) ...................................................................................................................................... 53
3.10 RHEED of epitaxial Al2O3/
30 nm Fe50Pt45Rh5 thin film ............................................................................................................ 55
4.1 XRD for Fe25Pt75 films with significant (111) texture .................................................. 59
4.2 (a) 220 (b) 110 pole figure for Fe25Pt75 film showing anomalous twelve
fold symmetry and some chemical ordering as well for the sample 1 ............... 60
4.3 (a) 220 (b) 110 pole figure for Fe25Pt75 film showing
poor epitaxy and no chemical ordering for the sample 2 ............................................. 61
4.4 XRD for Fe25Pt75 film with less (111) texture ..................................................................... 62
4.5 (a) 220 (b) 110 pole figure for Fe25Pt75 film showing good epitaxy and
no chemical ordering for the sample ....................................................................................... 63
5.1 Rocking curves for x=0, 3 and 6 thin films ........................................................................... 67
5.2 Rocking curves for x=9 and 12 thin films .............................................................................. 68
5.3 Out of plane hysteresis loops for various compositions of
Fe50-xMnxPt5 thin films on a-plane sapphire .......................................................................... 70
5.4 Coercivity as a function of composition the Fe50-xMnxPt50 thin films ................. 71
.
5.5 220 pole figure for Fe50-xMnxPt50 thin films ........................................................................... 74
5.6. RHEED Pattern for, x = 12 (top) and x= 15 (bottom) ................................................... 75
5.7 XRD 2θ-θ scans of 45nm Fe50-xMnxPt50 thin films ........................................................... 76
5.8 Auger spectra of Fe50-xMnxPt50 thin films ................................................................................ 77
5.9 Out of plane hysteresis loops for the Fe50-xMnxPt50 thin films .................................. 78
5.10 Saturation magnetization and coercivity as a function of composition for
the Fe50-xMnxPt50 thin films ............................................................................................................. 79
6.1 XRD of different Fe50-x Mnx Pt50 thin films ........................................................................... 81
xii
6.2 220 Pole figure for different concentrations of Mn rich Fe50-xMnxPt50 ............. 82
6.3 RHEED pattern for the different Fe50-x Mnx Pt50 thin films ........................................ 83
6. 4. Magnetization as a function of temperature for different
Substrates ....................................................................................................................................................... 86
6.5 Moment vs. Temperature measurement for ferromagnetic material,
FePt .................................................................................................................................................................... 87
6.6 Moment vs. Temperature measurement for Mn rich Fe50-xMnxPt50 ...................... 88
6.7 Hysteresis loops at different temperatures for MgO (100) substrate .................... 89
6.8 Hysteresis loops at different temperatures for x= 40 (top) and
and x=42 (bottom) ................................................................................................................................... 90
6.9. Hysteresis loops at different temperatures for x= 44 (top) and x=46
(bottom) ............................................................................................................................................................ 91
1
CHAPTER 1
INTRODUCTION
L10 Fe50-xMnxPt50 alloys have high anisotropy and magnetization, which makes them
ideal candidates for ultra high density magnetic recording media [1], as high anisotropy makes
information stable. The magnetic properties of Fe50 Pt50 can be tuned by the addition of other
elements. Tuning of magnetic phases of Fe50Pt50 by the addition of Mn is of particular interest as
magnetic properties can be engineered to find the optimum materials parameters like high
anisotropy and thermal stability for magnetic recording applications. Fe rich Fe50-xMnxPt50 alloys
are promising materials for heat assisted magnetic recording, and thus understanding the
temperature dependence of these alloys is very important to predict their thermomagnetic
properties. The ideal candidates will have a high anisotropy and a relatively low Curie
temperature. It has been suggested that materials with a saturation magnetization Ms of 900
emu/cm3, anisotropy K of 5 x 10
7 erg/cm
3 and a Curie temperature Tc of 600-650 K to achieve
high areal density of hard disk drive up to 4 Tb/in2 data storage [2]. Fe-rich Fe50-xMnxPt50 alloys
were predicted to show an increase in magnetic moment by 2% and magnetocrystalline
anisotropy 33% as compared to Fe50Pt50 [4]. A calculation of the magnetic phase diagram for
Fe50-xMnxPt50 published in 2015 [3] is shown in Fig. 1.1, where as the composition changes the
associated magnetic phase changes from purely ferromagnetic F, to mixed ferromagnetic and
antiferromagnetic phase F/G, non collinear canted antiferromagnetic phase G, mixed
2
antiferromagnetic phase G/C, mixed ferromagnetic and antiferromagnetic phase F/C and an
antiferromagnetic phase C. The symbols denote the experimental data, while the solid and
dashed lines shows the second order and first order transitions. The red lines shows the original
phase diagram due to the absence of non Heisenberg terms published in 1987 [56]. Fig. 1.2
show spin diagrams for the different phases.
Fig.1.1 Calculated phase diagram of Fe50-xMnxPt (adapted from ref.3) which shows the predicted
location of the new low temperature ferromagnetic phase F/C.
Fig. 1.2 Spin digrams of the phases: a) G-type, b) C-type, c) F/G, d) G/C, e) F/C. Large spheres
show Fe/Mn sites while small spheres show Pt sites , the green and brown arrows denote the
spins of Fe/ Mn atoms with non collinear magnetic ordering (ref 3)
3
High temperature ( over 800º C) and higher annealing time were known to induce
structural changes ( tetragonal phase ) in ordered L10 Fe50Pt50 [5] [6] [7] [8]. These strucutural
changes contributed to the increase in the coercivity in the L10 Fe50 Pt50 [9]. It has also been
shown that high temperature annealing ( over 650 º C) of L10 Fe50Pt50 coupled with the addition
of elements like Cu, Ag and Au were known to induce larger coercivity [10]. In this work we
show that by the addition of Mn to Fe50Pt50 at high temperature and higher annealing time could
result in the either increase or decrease in the coercivity. Fe50-xMnxPt50 thin films deposited on
sapphire substrates show a decrease in coercivity, while the Fe50-xMnxPt50 thin films deposited on
MgO substrates shows both decrease and increase in the coercivity, this increase considing with
structural changes in the Fe50-xMnxPt50 thin films. It was predicted that in FeMnPt alloy system,
the increase in magnetic properties is due to the presence of ferromagnetic alignment of Mn
atoms while the decrease in magnetic properties is due to presence of antiferromagnetic
alignment of Mn atoms [55]. We also investigate the prediction of low temperature
ferromagnetic phase in the Mn rich end of Fe50-xMnxPt50 magnetic phase diagram. In our
investigation we found that conventional magnetometry is not a reliable method to detect the
presence of low temperature ferromagnetism.
4
CHAPTER 2
EXPERIMENTAL TECHNIQUES
I. Auger Spectroscopy
Auger Electron Spectroscopy (AES) is a standard analysis technique in surface and interface
physics [12]. In this work, it was used predominantly to check the cleanliness and diffusion of
the thin films. In general, it can also be used in the studies of film growth and surface-chemical
composition (elemental analysis) as well as depth profiling of the concentration of particular
chemical elements.
AES is an electron core-level spectroscopy, in which the excitation process is induced by
a primary electron beam from an electron gun. The Auger process results in the ejection of
secondary electrons of relatively sharply-defined energy, which are analyzed and detected by a
standard electron analyzer. A Cylindrical Mirror Analyzer (CMA) is widely used in Auger
spectroscopy. As with all other electron spectroscopies, AES is surface sensitive because of the
limited escape depth of electrons. Typical probing depths in AES are in the range 10–30 Å. The
principle of the Auger process is explained in Fig. 1. The primary electron bombardment
produces an initial hole by ionization of a core level (K or L shell). Both primary electron and
core electron then leave the atom with some energy. This transition may be accompanied by the
emission of a characteristic x-ray photon, or Auger transition, in which the energy gained by the
electron that “falls” into the deeper atomic level is transferred to another electron of the same or
a different shell. This latter electron is then emitted with a characteristic Auger energy, thereby
5
leaving the atom in a double-ionized state. Thus for an Auger process at least three electrons are
needed in an atom.
Fig 2.1. Schematic of Auger process.
6
In the late 1940s and early 1950s, it was widely understood that due to the difference in
the surface and bulk properties of the materials, there will be striking differences in the device
performances, this idea triggered the need to analyze the surface of thin layered materials in the
late 1950s and 60s. Auger technique became a widely used tool for surface analysis due to the
work of Harris in 1968 [13]. He worked on a phase sensitive detection system in which a small
oscillating voltage is superimposed on a constantly increased voltage, thus by measuring the
collected electrons; it was possible to detect a small perturbation in the number of electrons at a
given energy level. Harris used a technique of electrical differentiation where a small modulating
voltage is applied to the grid and the collector is tuned to the frequency of the applied signal
(Fig. 2) to improve the signal to noise ratio of the measurement, this proved to be ground
breaking in making Auger spectroscopy as practical tool. The signal from the collector is used as
input to a lock-in amplifier. The lock-in amplifier is tuned to the frequency of the modulation
voltage. The lock-in output is then proportional to the derivative of the distribution curve. Fig. 3
shows the plot of derivative of the distribution curve with respect to electron energy taken with
alloy steel by Harris. After pioneering work by Harris, Palmerberg, Bohn and Tracy (Fig.4)
greatly improved the signal to noise ratio of the instrument by implementation of the cylindrical
mirror analyzer (CMA) [14]. The Omicron CMA used in this work closely resembles the one
developed by Palmberg, et al.
7
Fig. 2.2 Schematic of Auger system (ref.13) based on an electrostatic sector analyzer
configuration.
8
Fig.2. 3 Auger of alloy steel reported by Harris (ref.13).
9
Fig.2. 4 CMA developed by Palmerberg, Bohn and Tracy (ref.14).
10
II. Cylindrical Mirror Analyzer
Fig 2.5 Cylindrical Mirror Analyzer.
The CMA consists of coaxial cylinders and the Channel Electron Multiplier (CEM). In
the apparatus a pair of coaxial metal cylinders, each with length c and radius a and b for the inner
and outer cylinders (for our system CMA 100 Fig. 5, b = 4.3 cm and a = 1.86 cm, respectively,
are used to energy filter electrons emitted from a surface. A coaxial electron source is used as an
excitation source and electrons are emitted from the surface following a cos angular
distribution, where is the angle between the surface normal and emission direction. In this case,
electron diffraction effects are considered to be small and the emission can be considered to have
azimuthal symmetry. The outer cylinder is held at zero (or ground) potential and the inner
cylinder is placed at a positive potential, V, to select electrons emitted at a specific kinetic
11
energy, T, expressed in electron volts, In the apparatus, V is varied to select a value of T. The
apparatus is designed such that only electrons emitted within an angular acceptance of around
an average angle of emission of = 42º are detected. The angular acceptance is determined by
the dimensions of an aperture placed in front of the electron multiplier detector. For the purpose
of discussion, assume that the electrons enter the region between the cylinders at a radial distance
of r(0) = 1.25a with b = 2.3a and c ~ 2b (CMA100 geometry). The electrons that are detected
will exit the region at the same radial distance as they entered. For an infinitely long cylinders, it
is straightforward to show that the electrostatic potential between the cylinders is given by
and the corresponding radial component of the electric field is then
.
The radial force on an electron in this region is then
.
To find the trajectory, the motion in the axial (z) and radial (r) directions are considered
independently with no force in the z-direction, so the solution of the differential equation in the
r-direction determines the trajectory with the appropriate scale factor to convert time to
horizontal distance. Newton’s second law gives a nonlinear differential equation:
)()(
)/(2.1
trtr
meVr
12
subject to the initial conditions r(0) = 1.25a and initial velocity component in the r-direction of
.
sin2
)0(.
m
Tr
where T is the kinetic energy
There is no acceleration in the axial direction, so the scale factor to convert time to
distance in the axial direction is:
t
m
Ttz
m
Ttzz
z
cos2
)(
cos2
)()0(
0
.
The initial condition is given by
cz
z
)(
0)0(
where τ is the time for electron to reach the detector, while c is the distance from source to the
detector. Thus the time for electron to reach the detector is given by
cos2
m
T
c
substituting this time into the equation for r(t) gives a unique value of T for each eV. In our
system the two are linearly related with T = 1.65 eV, so by scanning voltage V, specific kinetic
energies are selected.
Electrons generally lose energy as they travel through the top few layers of the solid, this
inelastic scattering can be described by inelastic mean free path, in 1979 Seah and Dench [15]
described the function of inelastic mean free path and electron energy which is called “universal
13
curve” (Fig. 6), according to their model the inelastic mean free path with electron energy E
eVcan be described as
EBE
AE
2)(
where A=3000 eV2 Å and B = 0.5 Å/eV
1/2
For E > 100 eV, the equation can becomes
EBE )(
Fig 2.6 Inelastic mean free path of electron
.
The Auger electrons detected come from the topmost layers in the sample.In situ Auger
characterization of the sample provides the information of the diffusion process. The energy of
Auger electrons depend on the sample composition, energy of the incident electron, inelastic
mean free path of Auger electron and emission angle. A simple model is used to describe how
the intensities depend on film structure. Using Beer-Lambert’s law, the attenuation of Auger
1 10 100 10001
10
100
1000
Universal curve
(Ä)
Kinetic energy (eV)
14
intensity, I, for a single layer with a thickness d and inelastic mean free path (E0), (E) is given
by
cos)(/)(/
00 EdEd
eeII
where E0 is the incident electron energy (3 keV) and E is the measured energy, the detector is
collects electrons emitted at an angle θ = 42º. Thus for our experimental set up the attenuated
Auger intensity I from the substrate with an effective inelastic mean free path is given by
/
0
deII
Where
cos)()(
cos)()(
10
10
EE
EE
This model can be extended to n layers with constant separation, with a concentration
profile ci and relative Auger sensitivity factor, Si, for each layer, thus Auger intensity is given by
the sum over all the n layers
Fig. 2.7 shows the calculated intensities of Fe film on a Ni substrate. The ratio of Auger
intensities can be used to find the Fe thickness provided that the Fe grows as a continuous layer
on top of the Ni substrate. If there is an interface mixing of the top layers of the surface, then
intensity of the bottom layer will be higher, this can used to detect the diffusion process in a
multilayer thin film. Fig. 2.8 shows the influence of inter mixing of Ni/Nb bilayer at different
0
,)(
expn i
niiiE
ndcSI
15
temperatures. It can be seen that the bottom Nb layer segregates to the surface at higher
temperatures as the ratio of the peak heights of Nb/Ni increases with the temperature.
Fig. 2.7 Calculated intensities of Fe/Ni bilayer.
16
200 400 600 800 1000
T= 556 C
T= 523 C
T= 32 C
T= 676 C
Kinectic Energy (eV)
T= 836 C
dN/dE
Nb Ni
Ni/Nb
Fig. 2.8 Influence of inter-mixing of Ni/Nb bilayer at different temperatures.
17
III. Reflection High Energy Electron Diffraction
Thin films grown on single crystal substrates with well defined geometric relation to the
substrate atoms are called epitaxial thin films. RHEED is one of the electron based surface
diffraction techniques used for in situ characterization of epitaxial thin films in ultra high
vacuum [10-19]. For a single crystal substrate the RHEED shows specular dots and lines on the
phosphor screen, while epitaxial thin films grown on the single crystal substrate shows
characteristic streak pattern on the phosphor screen and if the film is polycrystalline the RHEED
pattern shows concentric rings because random orientation of the crystal produces continuous
angular distribution of spots. The RHEED set up consists of 15.2 keV electron gun (with De-
Broglie wavelength e=0.1 Å) facing the sample at an glancing angle, this low angle geometry
ensures diffracted electrons forming a series of streaks on the phosphorus screen for an epitaxial
thin film. The diffraction condition is given by
222
sin2
lkh
ad
d e
where d is the lattice spacing, is the angle of incidence, a is the lattice constant and h,k,l are
the miller indices. The lattice spacing can be related to the RHEED geometry which is given by
x
y
2tan
Where y is the observed spot spacing and x is the distance between the screen and point of
incidence. In this work, the main goal was to make highest epitaxial thin films and RHEED was
used as tool to check the quality of the substrate surface and epitaxial nature of the thin films.
18
Fig. 2.9 RHEED pattern for a GaAs (100) substrate (left) and Cr deposited on GaAs substrate at
550 Co (left).
Fig. 2.10 RHEED pattern for (left) the sapphire (1-120) and (right) 6 nm Cr / 14 nm Pt deposited
on a sapphire substrate.
Fig. 2.10 shows the RHEED pattern for (left) the sapphire (1-120) and (right) 6 nm Cr /
14 nm Pt deposited on the sapphire substrate at 730 Co and annealed at 875 C
o for 1.5 hrs.
19
RHEED was used to understand the evolution of intermetallic epitaxial thin films with respect to
the Cr seed layer and Pt buffer layer. One might expect a thin film to grow epitaxially on the 6
nm Cr/ 14 nm Pt bilayer (FIG. 2.10 right) but 30 nm Fe50 Pt45 Rh5 thin film grown on this seed
and buffer layer was polycrystalline. This can be seen in the Fig 2.11 where RHEED pattern for
30 nm Fe50 Pt45 Rh5 thin film show a ring pattern due to the polycrystalline nature of the thin
film.
Fig. 2.11 RHEED pattern for 30 nm polycrystalline Fe50 Pt45 Rh5 thin film.
20
IV. Sputtering
Sputtering originated from the Latin word sputare means to emit saliva with noise. It was
first found by Grove (1852) when a thin wire deposited a film on a silver surface at a pressure of
0.5 Torr [20]. In 1854 Faraday found a thin film deposited on the discharge tube, Plucker (1858)
also reported film deposition inside a gas discharge tube [21]. In late 1800s people were more
concerned about the cathode disintegration and sputtering slipped under the radar. Berkhardt and
Reineke (1939) improved the sputtering process by enhancing the ionization of atoms by using a
magnetic field across the target [22]. This type of sputtering is called magnetron sputtering. In
the 1960s vacuum technology improved and sputtering found a new lease on life. At present
magnetron sputtering is most cost effective technique used for thin film deposition.
Fig. 2.12 Schematic of Magnetron sputtering.
21
Magnetron sputtering enables a lowering of the operating pressure while achieving higher
deposition rate. These magnets are protected by cooling water from overheating as NdFeB
magnets have a maximum service temperature of 250ºC. In this process atoms are knocked out
of the target by the argon ions and deposited on the substrate. Generally lowest argon pressure is
desired to achieve stable plasma. For our experiment we found Ar pressure of 3.2 to 3.5 mTorr
produced stable plasma. Copper films with constant thickness were deposited at various powers.
The sputtering rate is proportional to the power which can be seen from the figure where copper
films of constant thickness were deposited at different thickness. This is very important property,
as the alloy composition of the thin films can be controlled by controlling the power of the
magnetron guns. For this work the thin films were made using multiple magnetron guns aimed at
a common focal point, this type of magnetron configuration is called confocal magnetron
sputtering. This configuration enables uniform deposition and ability to control the both
stoichiometric and non stoichiometric alloy compositions.
22
Fig. 2.13 Deposition rate vs. sputtering power.
0 20 40 60 80 1000
2
4
6
8
10
R
ate(
Å/s
)
Power (watts)
Slope = 0.094 Rate / Power
Cu Deposition Rates vs. Power
23
Fig. 2.14 Schematic of “RASCAL” sputtering system (ref. 25).
The sputtering of the thin films for this work was carried out in “RASCAL” sputtering
chamber at the MINT Center at the University of Alabama. Characterization of the thin films
was also performed in the main chamber. The ultra high vacuum is achieved by the rotary pump,
24
turbo pump and cryo pump. A load lock chamber connected to the main chamber is used to
transfer the substrates. The load lock has a quick access hinge door with a knob and is sealed
with a Viton gasket. The load lock is isolated from the main chamber during the sample loading
by a pneumatically controlled valve. The pressure of the load lock is monitored by the
thermocouple gauge MKS Model 917, while the pressure of the main chamber is monitored by a
convection enhanced Pirani gauge MKS Model 345 and a ion gauge. A SRS model residual gas
analyzer (RGA) is connected to the main chamber to monitor the presence of gases in the
vacuum chamber. In order to achieve ultra high vacuum in the range of 10-10
torr, it is imperative
that excess water and other adsorbed gases in the chamber should be removed by an out gassing
procedure called baking. This is achieved by wrapping heat tapes around the metallic parts of the
vacuum chamber (without the over lapping of tapes), this ensures the vacuum chamber
temperature rises to 165oC, the chamber is heated for 36 to 48 hrs, during this period the pressure
and temperature are monitored.
25
0 5 10 15 20 25 30 35 40 45 50 55 60 65
2.0x10-7
4.0x10-7
6.0x10-7
8.0x10-7
1.0x10-6
Tchamber
= 165 C0
Tchamber
= 151 C0
Pre
ssure
(T
orr
)
Hours
Tchamber
= 105 C0
Fig. 2.15 Pressure versus time during the baking process.
Fig. 2.16 and Fig.2.17 shows the partial pressure versus mass to charge ratio spectrum of the
unbaked and baked system measured by residual gas analyzer. In the unbaked case, the vacuum
chamber contains significant amount of water, after the baking procedure the partial pressure of
water is less than 1.2 × 10-9
torr. The main chamber is equipped with a manipulator which houses
the heater assembly (600 W halogen bulb) vacuum banana plugs. The heater can reach
temperatures exceeding 900oC. The substrate temperature is measured using thermocouple and a
pyrometer. Titanium sheets were used to reduce heat radiation and also to reduce heat loss. The
Z movement of the manipulator is motorized while the azimuth rotation is accomplished by
bellows-sealed rotary feed through. The manipulator side also houses required wiring for sample
heating and thermocouple for sample temperature measurement. The samples were mounted on a
26
custom built sample holder frame by means of thin tantalum wire using a spot welder. Sample
holder consists of four vacuum banana plugs mounted on the corners of the platen engages with
the sockets of the heater assembly on the manipulator. The sample transfer is accomplished by a
rotary linear drive manufactured by Transfer Engineering Inc. During the sputtering, the argon
gas flow rate was adjusted to a constant pressure of 3.4mTorr using a type MK50 Mass flow
controller. The flow measurement is based on the differential heat transfer between the
symmetrically attached temperature sensors in the gas flow path.
Fig. 2.16 Mass spectrum of an unbaked vacuum chamber with partial pressure of water at 2.6 x
10-7
torr.
27
Fig. 2.17 Mass spectrum of baked vacuum chamber with partial pressure of water at
1.1 x 10-9
torr.
28
Fig. 2.18 Sample heater and sample holder assembly.
29
V. Substrate Cleaning
It is very important that the substrates should be clean before the deposition. The presence of
impurities like carbon or oxygen is detrimental to the epitaxial growth. Substrates can be
contaminated in every step of substrate storage or preparation process. Sputter cleaning is done
using a Perkin-Elmer 20-045 Ion Gun. Fig. 2.19 shows the presence of carbon in a “clean”
degassed substrate probed using Auger analysis. It’s quite evident that the normal degassing
procedure is not enough to make contamination free substrates. Cleaning of the substrates
without damaging them by removing its contaminants is achieved using ion beam etching. In this
process the surface of the substrate is etched in a vacuum using inert gas ions like Ar+. This
process produces an atomically clean surface. In order to produce a clean substrate, a two step
process was used in all our experiments. The substrates are left in load lock vacuum for at least 8
hours and cleaned in the main chamber using ion beam etching. Ion beam is produced when
argon gas is leaked into the main chamber and ionized using an ion gun source at 0.5 kV. The
argon gas pressure is kept constant at 0.5 mTorr; the resulting Ar+ ions were accelerated and
focused on the substrate surface. Fig. 2.20 shows an Auger spectrum for the clean sapphire
substrate after 5 mins of ion beam etching.
30
Fig. 2.19 Presence of carbon on single crystal sapphire (Al2O3, 11-20).
0 200 400 600 800 1000
dN
/dE
Energy (eV)
Carbon presence in
Al2O
3 (11-20) substrate
Al
C
O
31
0 200 400 600 800 1000
dN
/dE
Energy(ev)
Al
O
Fig. 2.20 Sapphire substrate after 5 mintues of sputter cleaning at 0.5 mTorr at 0.5 kV beam
voltage.
32
VI. Flux Calibration
The deposition rate for the targets were calculated using a Quartz Crystal Microbalance
(QCM) from Inficon XTC and X-ray reflectivity (XRR). The QCM has a thin wafer of quartz,
the resonant frequency decreases when a thin film is deposited on the quartz wafer, this change is
frequency δf is directly proportional to mass of material deposited on the wafer. The thickness d
of material deposited on wafer is given by
2
22
2)1(
31
2 f
fZ
f
f
f
fVd
where f is the initial resonant frequency of the sensor, ρ is the density of the film, ρq the density
of quartz, Vq the velocity of sound in quartz and V the velocity of sound in the film. The quantity
ρV is referred to as the acoustic impedance; Δf is the total change in frequency of the QCM,
accumulated over its lifetime. The Z-factor is the ratio of the acoustic impedance of the
deposition material to that of the quartz is given by
qqV
VZ
The tooling factor is a correction for the difference in material deposited on the quartz sensor
versus the substrate. The thickness of thin film deposited on the substrate is compared with x-ray
reflectivity measurements (XRR) to calibrate the quartz crystal monitor, thus by adjusting the
tooling factor to match with XRR. Once the instrument is calibrated, the mass density and Z-
factor of the film material are entered and the microprocessor in the controller solves for d.
33
All our x-ray measurements were carried out using a Phillips X’pert x-ray
diffractometer. It has a four axis system with and where is the angle between the
incident beam and the diffracted beam, is the angle between the incident beam and sample
surface which is equal to is the rotation angle about the sample normal and is the tilt angle
about the sample surface The filament is normally operated at 40 kV and 45 mA. The Cu target
gives a Kα with a wavelength 1.542 Å, which is the average of Kα1 and Kα2. X-rays reflected from
a thin film undergo interference and forms fringes called Kiessig fringes. The square of the
angles θ2 m are plotted against the square of order of interference m2, the thickness t of the thin
film is found from the slope of the Kiessig plot [23].
order
542.1
densitylength scattering
angle critical
order m for the maximaintensity the toingcorrespond angles
thickness
4
4
22
th
2
2
2
2222
m
Å
t
t
mslope
t
m
c
c
m
cm
34
The above equation is (equations are) used to calibrate the flux for our experiments. The
scattering length density obtained from the NIST website [24] for various elements are used to
calculate the critical angles, these expected values are compared to the measured critical angles
using linearly fitted Kiessig plots ( Fig. 2.21). The XRR and Kiessig plot for the Fe50Pt50 and
Mn50 Pt50 is shown in figs 2.22 and 2.23 respectively.
0 50 100 1500
1
2
3
5
exp
2
52
1036.4
1021.3
ectedc
c
(ra
d2) Rate= 2.61 Å/s
m2
Kiessig Plot of Cr for 200s
Linear fit
Fig. 2.21 Kiessig plot of Cr deposited for 200 seconds at 4 mTorr Ar pressure.
35
2 4 6 8
0.01
0.1
1
10
100
1000
10000
100000
24 = 6.88º
23 = 5.24º
22 = 3.82º
21 = 2.32º
log(I
)
2(deg.)
FePt
0 2 4 6 8 10 12 14 16 18
0.0009
0.0018
0.0027
0.0036
218 sec FePt/Si
60 Watt, 3.2 mTorr
slope = 2.14x 10-4
t =51.9 Å
2 m (
x10
-3)
m2
Fig. 2.22 XRR scan of a Fe50Pt50 on Si thin film for thickness calibration purposes (a) and linear
fit to obtain film thickness (b).
36
2 4 6 8
0.01
0.1
1
10
100
1000
10000
25 = 7.04º
24 = 5.8º
23 = 4.46º
22 = 3.22º
21 = 2º
log(I
)
2(deg.)
MnPt
0 5 10 15 20 250.000
0.001
0.002
0.003
0.004
218 sec MnPt/Si
60 Watt, 3.2 mTorr
slope = 1.445x 10-4
t = 63.2 Å
2 m (
x10
-3)
m2
Fig. 2.23 XRR scan of a Mn50Pt50 on Si thin film for thickness calibration purposes (a) and linear
fit to obtain film thickness (b).
37
VII. X-ray diffraction
Fig.2.24 Geometry of X-ray diffraction.
X-ray diffraction is an analytical technique that tells the information about the crystal
structure of the thin film. The scattered intensity of the x-rays is observed as the function of
incident and scattered angle. The wavelength of the x-ray beam is close to the lattice spacing of
the crystal so the x-rays undergo constructive interference when the path difference is an integral
number of wavelengths in specific directions which satisfies the following Bragg’s equation.
2𝑑 𝑠𝑖𝑛𝜃 = 𝑛𝜆
where n is the integer, 𝜆 is the wavelength, and d is the distance between atomic planes. This is
used to determine the structure of the crystal. Pole figure measurement is a measurement
38
technique where the diffraction angle is fixed to particular plane spacing, say (220), and the
diffracted intensity is collected by varying the tilt angle and the rotation angle with respect
to the sample’s normal direction.
Fig. 2.25 Definition of angles in the X-ray goniometer
The effective sum of the amplitude of scattered waves by all the atoms in the unit cell is
expressed in terms of the structure factor F, which can be written as
)(2
1
nnn lwkvhuiN
nhkl efF
where f is the atomic scattering factor, thus the structure factor contains the information about
the locations (u,v,w) of atoms within a unit cell. Thus depending on the structure of a unit cell
the certain reflections will be allowed while some reflections will be forbidden. All the materials
studied in this work are based on the FCC and L10 crystal structure. In a FCC structure, the unit
cell consists of four atoms, the position of the atoms are given by
39
2
1,
2
1,0
2
1,0,
2
1
0,2
1,
2
1
0,0,0
444
333
222
111
wvu
wvu
wvu
wvu
Thus the structure factor F for FCC is given by
]1[ )()()( lkilhikhi eeefF
F=4f when h,k and l are all even or all odd
F=0, when h,k, and l are mixed (i.e. one even, two odd or one odd or two even). This means
(111), (200) diffractions are allowed while (001) are not allowed.
In L10 we see (100) diffractions because the atoms at specific sites are not equivalent as they
have alternate planes of different atoms.
Fig. 2.26 (a) FCC crystal structure (b) L10 crystal structure
40
0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1.0 1.1 1.20
11
22
33
44
55
66
77
Pt
Pd
Rh
Fe
f
Sin / [1 / Å]
Mn
Fig. 2.27 Atomic scattering factor for different metals (excluding dispersion corrections).
During the scattering process the x-rays may result in dispersion (separation of wave into
components of different frequencies). This can be accounted by dispersion correction factor. The
thermal vibration can change the intensity of x-rays, this temperature dependence of intensity can
be corrected using Debye-Waller temperature factor. Thus atomic scattering factor can be written
as
2sin
)(
BM
efifff M
corrected
41
where fcorrected is the corrected atomic scattering factor, f is the atomic factor, f and f are the
real and imaginary dispersion correction respectively, while Me is the Debye Weller
temperature factor with B is the constant which depends on each element.
The corrections for geometry and the polarization is accounted by the Lorentz
polarization factor which decreases the intensity at the intermediate angles and increase the
intensity in the forward and backward direction.
cossin
2cos1 2
X-ray intensity also depends on the film thickness, this can be corrected by the absorption
factor.
sin
2
1
t
eG
where μ is the absorption coefficient, which can be extracted from the International tables for x-
ray crystallography [17] and t is the total thickness of the thin film.It is possible to estimate the
chemical order parameter S of the films from the integrated peak intensities of 100 and 200
planes. The x-ray intensity of the diffracted beam from the 100 or 200 plane is given by [23]
sin
222
200/100 1cossin
2cos1t
eFI
42
theoryI
I
II
S
200
100
exp200
100
2
where I100 and I200 are the integrated intensities of the superlattice (100) peaks and (200)
fundamental peaks. Thus order parameter S can be estimated from the experimental and
calculated intensities ratio for perfect chemical order.
43
CHAPTER 3
FE50PT45RH5 FILMS
Intermetallic alloys of FePt with metals like Rh, Pd or Mn have tunable magnetic
properties. The magnetic phase change properties of these intermetallic alloys has been studied
for future technological applications like thermally assisted recording media [25,26, 30]. In this
work an effective epitaxial growth procedure for intermetallic thin films has been developed. The
surfaces of the films were analyzed using in situ Auger and RHEED techniques. We used a seed
layer which is chemically different from the intermetallic alloy. The films are deposited at
different temperatures to understand the effects of temperature on the diffusion of the seed layer.
It is shown that even with a very thin seed layer, inter diffusion is possible if the substrate
temperature is too hot; this was quite evident when the Cr seed layer of both 3nm and 6 nm
diffused to the surface of single layer of 30 nm Fe50 Pt45 Rh5 layer. The X-ray results show that
diffused single layer shows texture. The goal was to grow an epitaxial intermetallic thin film like
Fe50Pt45Rh5 without surface diffusion as this leads to the polycrystalline structure, as indicated by
the RHEED analysis which showed the signature ring shaped pattern. FexPt50-xRhx were known
to grow epitaxially on a-plane sapphire substrates at 600-700º with 6 nm Cr seed layer and 14 nm
Pt buffer layer [27, 28, 29, 31, 32]. 30 nm Fe50Pt45Rh5 films were prepared in the UHV
sputtering chamber with a base pressure lower than 6 x 10-10
torr, prior to deposition, the surface
of the substrates were cleaned with 5 minutes of ion sputtering at 5 x 10-4
argon pressure and
44
followed by 700ºC degassing for few hours. The sputtering rates of each target were calibrated
with quartz crystal microbalance and XRR. Fig 3.1 shows the X-ray diffraction of a 30 nm
Fe50Pt45Rh5 (111). Fig. 3.2 shows surface analysis with Auger spectra reveals the diffusion of the
Cr seed layer to the surface, while the RHEED (Fig.3.3) shows the presence of polycrystalline
material on the surface. It is interesting to note that X-ray measurements of this thin film show
the presence of certain degree of epitaxy. Fig. 3.4 shows the phi angle measurement, the thin
shows the presence of six-fold symmetry while the Fig.3.5 pole figure measurement of (111)
shows the presence of textured domains which coincides with RHEED measurements. The
angles for the different polycrystalline domain regions are summarized in the tables.
39 40 41 42 430
100
200
300
plane Al2O
3/ 6 nm Cr/ 14 nm Pt/30 nm Fe
50Pt
45Rh
5(111)
Inte
nsity (
Cps)
Scattering angle (2)
Fig. 3.1 X-ray diffraction spectra of a-plane Al2O3/ 6 nm Cr/ 14 nm Pt/30 nm Fe50Pt45Rh5 (111)
45
200 400 600 800 1000
6 nm Cr/14 nm Pt/ 30 nm Fe50
Pt45
Rh5
dN
/dE
Kinetic energy (eV)
C
Cr
FePt Rh
Tdep
= 725 C
Fig. 3.2 Auger spectra of Al2O3/ 6 nm Cr/ 14 nm Pt/30 nm Fe50Pt45Rh5 (111) shows the presence
of Cr diffusion.
46
Fig.3.3 RHEED shows the polycrystalline Al2O3/ 6 nm Cr/ 14 nm Pt/30 nm Fe50Pt45Rh5 thin
film.
47
Fig.3.4 Fe50Pt45Rh5 (111) Pole figure of Al2O3/ 6 nm Cr/ 14 nm Pt/30 nm for a polycrystalline
thin film.
Table I. Angles and intensities for (111) texture during Cr diffusion for the region I
deg.
deg.
Max. Intensity
(cps)
29.5 69 7.7
29.5 70 6.2
30.5 69 12.33
30.5 70 9.63
31.5 69 13.35
31.5 70 10.4
48
Table II. Angles and intensities for the region II during Cr diffusion
Table III. Angles and intensities for the region III during Cr diffusion
deg.
deg.
Max. Intensity
(cps)
-0.5 14 9.63
0.5 15 9.63
-0.5 14 9.13
0.5 15 9.13
deg.
deg.
Max. Intensity
(cps)
-0.5 52 37.25
-0.5 53 42.07
0.5 52 28.26
0.5 53 35.53
49
Table IV. Angles and intensities for the region IV during Cr diffusion
Table V. Angles and intensities for the region V during Cr diffusion
Table VI. Angles and intensities for the region VI during Cr diffusion
deg.
deg.
Max. Intensity
(cps)
4.5 69 2.67
7.5 69 2.54
deg.
deg.
Max. Intensity
(cps)
10.5 54 2.34
11.5 55 3.45
deg.
deg.
Max. Intensity
(cps)
22.5 61 3.23
22.5 62 2.47
23.5 61 2.47
24.5 62 2.28
50
Table VII. Angles and intensities for the region VII during Cr diffusion.
Diffusion of the seed layer to the top layer may alter the magnetic properties of
intermetallic thin film. This was recently reported regarding the anomalous interfacial
magnetization in the FeRhPd thin films [36]. It should be noted that they used a Rh seed layer
which is chemically similar to FeRhPd thin film which could have induced the anomalous
interfacial magnetization.
deg.
deg.
Max. Intensity
(cps)
33.5 78 4.59
33.5 80 4.0
34.5 78 4.0
34.5 80 3.76
51
Fig. 3.5 Anomalous interfacial magnetization in a single layer and trilayer FeRhPd films ( ref.36)
In order to minimize/stop the diffusion of seed layer, we tried three routes; the first route
is to reduce the size of the seed layer and buffer layer to 3 nm Cr and 12 nm Pt, the second one is
to reduce the temperature and the third option is to stop using the seed and buffer layer. Fig. 3.6
shows the absence of surface diffusion of Cr in a 30 nm Fe50Pt45Rh5 thin film grown on a 3 nm
Cr seed layer and 12 nm Pt buffer layer at a relatively low temperature of 5110C.
52
200 400 600 800 1000
Growth temperature = 510ºC
3 nm Cr/ 12nm Pt/30 nm Fe50
Pt45
Rh5
dN
/dE
Energy
Fe
PtRh
Fig. 3.6 Auger spectra of Al2O3/ 3 nm Cr/ 12 nm Pt/30 nm Fe50Pt45Rh5 (111) shows the absence
of Cr diffusion
Fig. 3.7 RHEED of Al2O3/ 3 nm Cr/ 12 nm Pt/30 nm Fe50Pt45Rh5 (111)
RHEED (Fig. 3.7) shows the mixture of rods and rings which is the characteristic feature
of epitaxial and polycrystalline nature. It is interesting to note that according to the phi scan
measurement (Fig. 3.8) this sample is epitaxial in nature while pole figure measurement (Fig.
53
3.9) reveals the presence of epitaxial (111) texture (region I) in the sample. The angles and
intensities for (111) texture are summarized in the Table I.
-40 -20 0 20 400
5
10
15
20
253 nm Cr/12 nm Pt/ 30 nm Fe
50Pt
45Rh
5 (111)
Inte
nsi
ty
Azimutal angle
Fig. 3.8 Phi scan of Al2O3/ 3 nm Cr/ 12 nm Pt/30 nm Fe50Pt45Rh5(111) showing six fold
symmetry.
Fig. 3.9 shows the Fe50Pt45Rh5 (111) pole figure of Al2O3/ 3 nm Cr/ 12 nm Pt/30 nm thin film.
54
Table VIII. Angles and intensities for (111) texture during no Cr diffusion for the region I
It was found (Fig. 2.10) that the seed layer and buffer layer grow epitaxially on sapphire
substrate at 7500C and 1.5 hrs of annealing at 880
0C. 30 nm Fe50Pt45Rh5 is deposited directly on
a plane sapphire Al2O3 at 7500C and 1.5 hrs of annealing at 880
0C. Fig. 3.10 shows the RHEED
pattern for the 30 nm Fe50Pt45Rh5 with characteristic epitaxial nature.
deg.
deg.
Max. Intensity
(cps)
29.5 69 30.6
29.5 70 25.3
30.5 69 32.67
30.5 70 27.52
31.5 69 27.95
31.5 70 24.2
55
Fig. 3.10 RHEED shows the (left) buffer and seed layer annealing (right) direct deposition of
epitaxial Al2O3/30 nm Fe50Pt45Rh5 thin film.
It can be seen that the seed and buffer layer are not necessary with a high temperature annealing.
Thus direct deposition and high temperature annealing is sufficient to grow epitaxial thin films.
56
CHAPTER 4
FE25PT75
Chemically ordered Fe25Pt75 can display two coexisting antiferromagnetic phases with
TN1~160 K and TN2~100 K, while the ferromagnetic state can be seen in the chemically
disordered state [37]. Thick films of ordered Fe25Pt75 when irradiated with 15 keV He ions,
results in a bi-layer of disordered Fe25Pt75 and ordered Fe25Pt75 [38]. The energy of He ions is
such that they only penetrate approximately half way through the thick Fe25Pt75, thus exchange
bias can be created between the disordered ferromagnetic layer and ordered anti-ferromagnetic
layer. Three samples of thick epitaxial Fe25Pt75 of 280 nm thickness were prepared by
cosputtering using Fe50Pt50 and Pt. The first two samples were prepared without any buffer and
seed layer while the third sample was prepared using Fe seed layer and CrPt3 buffer layer. The
desired stiochiometry for any alloy thin film can be grown by co sputtering with pure targets or
from alloy targets or with the combination of both [30,31]. Let us consider the conditions
required to fabricate Fe100-x Ptx thin films from Fe and Pt targets. The rate of mass m deposited
per unit time are using these targets is given by
PtPtPt
FeFeFe
Rm
Rm
where R and ρ are the deposition rate and density of the elemental targets, while the mass for the
elemental targets in one mole of Fe100-x Ptx thin films is given by
57
PtPt
FeFe
Mxm
Mxm
)100(
where M is the atomic weight of the elemental targets, now equating the mass ratios we ge
Pt
Fe
Pt
Fe
m
m
m
m
PtFe
FePt
Pt
Fe
Pt
Fe
PtPt
FeFe
Mx
Mx
R
R
Mx
Mx
R
R
)100(
)100(
The deposition rate R is directly proportional to the power P supplied to the sputtering gun. It can
be written as
PcR
If the deposition rates for the elemental targets are known for a given power, the constant c can
be determined. Then the deposition rate of one of the target is kept constant and the deposition
rate for the other target can be solved to achieve the desired composition of the Fe100-x Ptx thin
films. To fabricate Fe100-x Ptx thin films from Fe50 Pt50 and Pt, we have
xxjii PtFePtPtFe 100
This implies
xji
xi
100
This gives xj 2100
Now one can solve for the deposition rate R ratio for Fe50Pt50 and Pt to give Fe25Pt75
58
FePt
Pt
Pt
FePt
Pt
FePt
x
x
M
M
R
R
2100
where mFePt and mPt is the atomic weights of Fe50Pt50 and Pt, RFePt and RPt is the deposition rate
for Fe50Pt50 and Pt, while Pt and FePt are the density of Pt and Fe50Pt50. The deposition rate for
each target determines the stoichiometry of the film. The power for the targets is set according to
the above equation to get the desired stoichiometry of the film, since the deposition rate is
directly proportional to the power.
During the epitaxial growth, the film orientation is determined by the substrate surface.
During the deposition the thin film adapts to in plane lattice parameters but for some values if
pressure or temperature, the thin film may adopt to grow two or more in plane orientation crystal
planes due to the strain. This phenomenon is called twinning. In this work we show that direct
deposition of Fe25Pt75 on MgO (100) gives rise to (111) texture. We also show that this twinning
of (111) planes can be reduced insignificantly by a BCC (110) textured seed layer like Fe. This
alignment of BCC (110) along the direction of FCC unit cell is a feature of a Bain path [48]
which significantly reduces the (111) twinning structure. Sample 1 and sample 2 were deposited
at 770ºC and annealed at 930ºC for 1 hr. The X-ray diffraction for these two samples shows
significant 111 texture. It is interesting that the sample 1 shows twelve fold symmetry while the
sample 2 shows four-fold symmetry. In order to reduce the 111 texture, 1nm Fe seed layer and 2
nm CrPt3 seed layer was used for the sample 3, the figure 4.4 shows that 111 texture has been
reduced when compared to the samples without seed and buffer layers. The pole figure also
shows four-fold symmetry. The pole figures of 220 peak verifies the epitaxial nature of the thin
film while 110 peak indicate the L10 ordering as it is forbidden in FCC structure.
59
20 25 30 35 40 45 50
10
100
1000
10000
100000
sample 1
FePt3(002)
FePt3(111)
Inte
nsi
ty (
cps)
Scattering angle
FePt3(001)
MgO (100)/ 280 nm FePt3
MgO(002)
20 25 30 35 40 45 50
10
100
1000
10000
100000
MgO (100)/ 280 nm FePt3
FePt3(002)
FePt3(001)
FePt3(111)
Inte
nsi
ty (
cps)
Scattering angle
sample 2
MgO (002)
Fig. 4.1 XRD for Fe25Pt75films with significant 111 texture and order parameter less than 0.2 ( S
< 0.2)
60
Fig. 4.2 (a) 220 (b) 110 pole figure for Fe25Pt75 film showing anomalous twelve fold symmetry
and some chemical ordering as well for the sample 1.
61
Fig. 4.3 (a) 220 (b) 110 pole figure for Fe25Pt75 film showing poor epitaxy and no chemical
ordering for the sample 2.
62
20 25 30 35 40 45 501
10
100
1000
10000
100000
FePt3(001)
FePt3(002)
FePt3(111)
Inte
nsi
ty (
cps)
Scattering angle
sample 3Mgo(100)/1nm Fe/2nm CrPt3/280 nm FePt
3
MgO (002)
Fig. 4.4 XRD for Fe25Pt75 film with less 111 texture
63
Fig. 4.5 (a) 220 (b) 110 pole figure for Fe25Pt75 film showing good epitaxy and no chemical
ordering for the sample 3.
64
CHAPTER 5
FE RICH FE50-X MNXPT50
Ordered Fe50Pt50 thin films have high anisotropy [41,42] and they are a promising
candidates for recording media like heat assisted magnetic recording [1]. The recording media
requires small grains to achieve high density, and for the signal to noise ratio to be high, the
grains should also possesses high anisotropy energy. However, high anisotropy energy also
makes it harder for writing media. Fe50-xMnxPt50 is a magnetic chameleon system where small
change in the Mn concentration can lead to a significant change in the magnetic properties [2],
which make them ideal candidates for thin film recording media. Developing methods to control
the magnetization and anisotropy are needed to engineer the material for optimal properties [4].
Growing a ferromagnetic film directly on a given substrate inherently induces structural change
in the ferromagnetic film due to stress caused by the lattice mismatch.
These structural changes can be a very important in tuning the magnetic properties of
ferromagnetic films like coercivity and magnetization as they lead to changes in electronic
structure which may alter the magnetic properties. Ferromagnetic properties like coercivity were
shown to change with respect to the thickness of the thin film in Fe50Pt50 [43-47]. It was shown
that by exploiting the tetragonal distortion, magnetic properties like magneto crystalline
anisotropy can be maximized [48]. Tuning the structural properties in a magnetic chameleon
system like Fe50-xMnxPt50 opens a range of tunable magnetic properties. In L10 Fe50Pt50, the high
anisotropy is due to the combination of tetragonal distortion, exchange effects and spin orbit
moment of Pt [49]. It was predicted that maximum anisotropy for Fe50Pt50 would be achieved
65
at c/a ratio close to 0.96 [50]. Thus tetragonal distortion is one of the important factors in
maximizing the magneto crystalline anisotropy. Buker et al. [3] predicted that magnetic
anisotropy of Fe50-xMnxPt50 increases by 33% and magnetic moment increase of 2% at x=12,
however the prediction does not agree with subsequent experimental results [52-54]. Meyer [52]
attributed the significant reduction of magnetization and anisotropy due to the antiparallel
alignment of Fe and Mn moments observed in x-ray magnetic circular dichroism (XMCD). It
should be noted that the previous experimental results didn’t show a tetragonal distortion in the
XRD 2θ- θ scans. Thus the absence of a tetragonal distortion coincides with the decrease in the
magnetic properties in the Fe50-xMnxPt50 system. This can be attributed to the presence of
antiferromagnetic phase (AFM) of Mn species from the previous experimental results which
contributed to decrease in the magnetic properties. Cuadrado et al. [55] explained that in L10
Fe50-xMnxPt50 bulk alloy, Mn atoms in AFM alignment led to the decrease in the total magnetic
anisotropy and Mn atoms in FM alignment led to the increase in the total magnetic anisotropy.
66
Epitaxial Fe1-xMnxPt50 (x=0, 3, 6, 9, 12) thin films were prepared by co-sputtering Fe50Pt50
and Mn50Pt50 on to Al2O3 (1-120) single crystal substrates at 3.2 mTorr Ar with the thickness
fixed at 50 nm. These thin films are deposited directly on the sapphire substrates at 780ºC and
were annealed at 920ºC for 1 hr. The rocking curves for these samples are shown in the figure
5.1 and 5.2. The rocking curves for the concentration x=0, 3 and 6 show ordered phase. Order
parameter S for various compositions shown in the Table IX ( the dispersion corrections were
ignored in the calculations), while the x=9 show reduced intensity of 100 peak but x=12 show
almost no 100 peak. Both are characteristic features of a disordered phase. By comparing the
integrated intensities of the rocking curves of the ordered Fe1-xMnxPt50 to for a perfectly ordered
thin film, the amount of order can be determined. The structure factor F for the Fe-Mn-Pt system
is given by
22 )(2 MnFePt ffF for the (100) peak
22 )(2 MnFePt ffF for the (200) peak
Where the atomic factor for fFe/Mn is given by
MnFeMnFe fx
fx
f
5050
)1(
67
FIG. 5.1 Rocking curves for x=0, 3 and 6 thin films
6 8 10 12 140
5
10
15
20
25
30
A12O
3 (1-120)/ 52 nm Fe
50Pt
50(100)
Inte
nsi
ty (
cps)
18 20 22 24 26 28
0
5
10
15
20
25
30
A12O
3 (1-120)/
52 nm FePt(200)
Inte
nsi
ty (
cps)
10 15
10
20
30
40
50
60
Scattering angle ()
Al2O
3 (1-120)/ 52 nm Fe
47Mn
3Pt
50 (100)
Inte
nsi
ty (
cps)
20 25
10
20
30
40
50
60
Al2O
3 (1-120)/ 52 nm Fe
47Mn
3Pt
50 (200)
Inte
nsi
ty (
cps)
Scattering angle ()
10 150
10
20
30
40
50Al
2O
3 (1-120)/ 52 nm Fe
44Mn
6Pt
50 (100)
Inte
nsi
ty (
cps)
scattering angle ()
20 250
10
20
30
40
50
Al2O
3 (1-120)/ 52 nm Fe
44Mn
6Pt
50 (200)
Inte
nsi
ty (
cps)
scattering angle ()
68
FIG. 5.2 Rocking curves for x=9 and 12 thin films.
10 15 20 250
5
10
15
Al2O
3 (1-120)/ 52 nm Fe
41Mn
9Pt
50 (100)
Inte
nsi
ty(c
ps)
Scattering angle ()
15 20 25 30 35 400
5
10
15
Al2O
3 (1-120)/ 52 nm Fe
41Mn
9Pt
50 (200)
Inte
nsi
ty(c
ps)
Scattering angle ()
20 30
1
2
3
4
5
Al2O
3 (1-120)/ 52 nm Fe
38Mn
12Pt
50 (200)
Inte
nsi
ty (
cps)
Scattering angle ()
10 20
1
2
3
4
5
Al2O
3 (1-120)/ 52 nm Fe
38Mn
12Pt
50 (100)
Inte
nsi
ty (
cps)
Scattering angle ()
69
TABLE IX. Order parameter S for various compositions of Fe50-xMnxPt50 thin films on a-plane
sapphire.
The out of plane hysteresis loops for various composition (figure 5.3) shows decrease in
corecivity which agrees with the pervious experimental results. FePt (x=0) has a coercvity of 12
kOe For x=0.06 the corecivity decrease to 9.06 kOe while the x = 0.12 and x = 0.18 has the
coercivity of 6.7 and 7 kOe respectively. For the composition x=0.24 the sample shows drastic
decrease in the coercivity of 2.5 kOe this can also be attributed to the disordered phase as found
by XRD. Coercivity as a function of composition for the Fe1-xMnxPt50 thin films on a-plane
sapphire is shown in the fig. 5.4. This increase in coercivity can be attributed to the tetragonal
structure and not the sublattice ordering of Mn atoms in FM alignment as implied by Cuadrado
[55].
Fe1-xMnxPt50 100 200
x fFe-Mn fPt G fFe-Mn fPt G S
x=0 21.6 69.78 10.19 0.0383 16.92 58.84 3.94 0.0410 0.82
x=3 21.09 68.86 8.70 0.0394 16.87 58.84 3.90 0.0421 0.85
x=6 21.06 68.86 8.86 0.0392 17.18 59.73 4.038 0.0419 0.84
x=9 - - - - - - - - 0
x=12 - - - - - - - - 0
70
FIG. 5.3 Out of plane hysteresis loops for various compositions of Fe1-xMnxPt50 thin films on a-
plane sapphire
-20 0 20
-2
0
2
Al2O
3 (1-120)/ 52 nm Fe
47Mn
3Pt
50 (111)
mo
men
t (m
emu
)
Applied field(kOe)
x=0.06
x=0
-20 0 20
-2
0
2
Al2O
3 (1-120)/ 52 nm Fe
44Mn
6Pt
50 (111)
mo
men
t (m
emu
)
Applied field (kOe)
x=0
x=0.12
-20 0 20
-2
0
2
Al2O
3 (1-120)/ 52 nm Fe
41Mn
9Pt
50 (111)
mo
men
t (m
emu
)
Applied field (kOe)
x=0.18
x=0
-20 0 20
-2
0
2
Al2O
3 (1-120)/ 52 nm Fe
38Mn
12Pt
50 (111)
mo
men
t (m
emu
)
Applied field (kOe)
x=0
x=0.24
71
0 5 10 150
5
10
15
Fe50-x
MnxPt
50 on Al
2O
3(11-20)
Co
erci
vit
y (
kO
e)
X (Mn concerntration)
FIG.5.4. Coercivity as a function of composition for the Fe50-xMnxPt50 thin films on a-plane
sapphire.
72
Epitaxial Fe50-xMnxPt50 (x=0, 6,9,12 and 15) thin films were prepared by co-sputtering
Fe50Pt50 and Mn50Pt50 on to MgO (100) single crystal substrate at 3.5 mTorr Ar and the thickness
was fixed at 45 nm. The thin films were directly deposited on the MgO substrate at 780ºC to
avoid diffusion and the samples were annealed at 920ºC for 1 hr. Before deposition the substrates
were sputter cleaned at 0.5 mTorr Ar for 5 mins. Figure 5.5 shows the 220 pole figure for the
different Fe50-xMnxPt50 thin films where the four fold symmetry can be seen. Figure 5.6 shows
the presence of 100 and 200 in the RHEED pattern for the Mn concentrations (x= 12 and 15)
which is consistent with chemical ordering. XRD 2θ- θ scans for different concentrations (x=0,
6,9,12 and 15) are shown in the figure 5.7. The shift in the 001 and 002 peaks shows the
expansion of c-axis. For x= 9, 12 and 15 the XRD scans shows the splitting of 200 peaks due to
tetragonal distortions, as the reflection from peaks like 002 and 200 are no longer equivalent in
the tetragonal structure. The distance d between the planes (h,k,l) in a tetragonal lattice is given
by
2
2
2
22
2
1
c
l
a
kh
d
Substituting this in the Bragg’s equation, we get
2
2
2
22
2
2
2222
sin4
sin4
sin2
c
l
a
kh
nd
nd
The out of plane c and in plane lattice parameter a were solved using the doublet peaks,
while for the singlet peaks out of plane c were solved from 002/001 peaks and used to calculate
73
the in plane lattice parameter a in the tetragonal lattice for 202 peaks, c and a values are
summarized in the Table X. The auger measurement for these samples is shown in the figure 5.8.
Figure 5.9 shows the out of plane hysteresis measurement at 300K. For the two films x=6 and x=
9, the magnetization decreases which is consistent with the previous results but x=9 shows
significant increase in the coercivity which coincides with the onset of the 200 peak. The films
x= 12 and 15 shows significant increase in both magnetization and coercivity which is different
from the previous results, these films show prominent peak splitting due to tetragonal distortion.
The decrease in the coercivity and magnetization is due to the addition of Mn in Fe50Pt50 system
could be attributed to the presence of AFM phase in Mn atoms which agree with the previous
results, while the increase in the magnetic properties due to the addition of Mn could be
attributed to the presence of FM phase in Mn atoms at the same or different atomic planes. So
the presence of peak splitting in the L10 Fe50-xMnxPt50 thin films could be due to the appearance
of FM phase. Figure 5.10 shows the relation between coercivity, saturation magnetization with
respect to the composition for the Fe1-xMnxPt50 thin films on MgO (100).
74
FIG. 5.5. 220 Pole figure for Fe50-xMnxPt50 thin films.
75
FIG.5.6. RHEED pattern for, x = 12 (top) and x= 15 (bottom)
76
22 23 24 25 46 47 48 49 50
001
001
001
001
200
200
002
002
002
002
002200001
x=0
x=6
x=9
x=12
x=15
X-r
ay i
nte
nsi
ty (
a.u
.)
Scattering angle (deg.)
Fe50-x
MnxPt
50
FIG. 5.7. XRD 2θ-θ scans of 45nm Fe50-xMnxPt50 thin films
77
200 400 600 800 1000
Fe
Fe
Fe
Pt
Pt
Pt
dN
/dE
Kinetic energy (eV)
x=0
x=3
x=9
x=12
Fe50-x
MnxPt
50
FePt
FIG. 5.8. Auger spectra for Fe50-xMnxPt50 thin films
TABLE X. Composition, out of plane and in plane lattice parameters, c/a ratio, saturation
magnetization and coercivity of Fe50-xMnxPt50 films, all at 300 K
Fe50-xMnxPt50 c ( Å ) a ( Å)
(200)
c/a a ( Å)
(202)
Ms
(Oe)
Hc
(Oe)
0 3.711 - 0.958 3.875 1046 1890
3 3.715 - 0.961 3.865 1030 1753
6 3.719 3.842 0.968 - 990 2629
9 3.742 3.851 0.971 - 1140 2767
12 3.747 3.854 0.972 - 1061 2137
78
-20 -10 0 10 20
-1000
-500
0
500
1000
Mag
net
izat
ion
(em
u/c
c)
Applied field(kOe)
Fe44
Mn6Pt
50
Fe50
Pt50
x = 0.12
-20 -10 0 10 20
-1000
-500
0
500
1000
x = 0.18
Mag
net
izat
ion
(em
u/c
c)
Applied field(kOe)
Fe50
Pt50
Fe41
Mn9Pt
50
-20 -10 0 10 20
-1000
-500
0
500
1000
x = 0.24
Mag
net
izat
ion
(em
u/c
c)
Applied field(kOe)
Fe50
Pt50
Fe38
Mn12
Pt50
-20 -10 0 10 20
-1000
-500
0
500
1000
x = 0.3
Mag
net
izat
ion
(em
u/c
c)
Applied field(kOe)
Fe50
Pt50
Fe35
Mn15
Pt50
FIG.5.9. Out of plane hysteresis loops for the Fe50-xMnxPt50 thin films
79
0 2 4 6 8 10 12 140
200
400
600
800
1000
Composition (x)
Ms
Hc
Co
erc
ivit
y (
Oe)
Satu
rati
on
Mag
neti
zati
on
(em
u/c
c)
0
500
1000
1500
2000
2500
3000
FIG.5.10. Saturation magnetization and out of plane coercivity as a function of composition for
the Fe50-xMnxPt50 thin films on MgO (100)
80
CHAPTER 6
MN RICH FE50-XMNXPT50
The Mn rich end of bulk Fe50-x MnxPt50 phase diagram was published in 1987 [56]. The
authors investigated ordered powered samples of the ternary alloy system of Fe50-x MnxPt50
through neutron scattering and concluded that the Mn rich end of the phase diagram is
antiferromagnetic structure. In 2011, a first principles calculation also found that Mn rich
Fe50Pt50 [57]. In 2015, it was predicted that Mn rich end of Fe50-x MnxPt50 ternary alloy system
may have a low temperature ferromagnetic phase which was not observed in previous
experiments [2]. In this work four samples of L10 Mn rich concentrations (x = 40, 42, 44 and 46)
epitaxial films of 45 nanometers where prepared on MgO (100) single crystal substrates to verify
the prediction.
X-ray diffraction for the four samples is shown in the figure 6.1. The presence of 001
peaks shows that a portion of the films are ordered. The four fold symmetry of the four samples
is shown through the (220) pole figures in the figure 6.2. This is consistent with the RHEED
pattern for these films but the intensity of (001) peak is very small for all four samples which
suggest poor ordering. The lattice parameter for the samples increased from 3.916 to 3.927 Å
which is close to the reported value for the Mn50Pt50 lattice parameter of 4.0 Å [58].
81
24 45 46 47 48
200
200
200
200
001
001
001
x=46
x=44
x=42X-r
ay i
nte
nsi
ty (
a.u
)
Scattering angle (deg.)
x=40
Fe50-x
MnxPt
50
001
FIG. 6.1. XRD of different Fe50-x Mnx Pt50 thin films as a function of x.
82
Fig. 6.2. 220 Pole figure for different concentrations of Mn rich Fe50-xMnxPt50.
83
FIG.6.3. RHEED pattern for the different Fe50-x Mnx Pt50 thin films ( Top (left) x=40, top (right)
x= 42, bottom (left) x=44, and bottom (right) x=46
It has been claimed in the past that many non magnetic materials show magnetic behavior
at low temperature which turned out to be controversial due to reproducibility [59-63]. Spurious
results arise from experimental difficulties in measuring small moments, These artifacts are not
so important in bulk samples but in thin films it becomes prominent and the choice of substrate,
84
the handling of the substrate, preparation methods of thin films like annealing temperature and
annealing time, the choice of sample holders, Teflon tapes which were used to fix samples to the
sample holder of the instrument, the presence of very small amount of oxygen in the instrument
are only a few sources of small false magnetic signals [64]. This is because at low temperature
the walls of the chamber, the sample holder and the sample itself act as a cryopump, hence the
impurities condense on these cold surfaces, there by inducing paramagnetic contribution to the
measurements which can result in incorrect moment values and temperature dependence. Hence
the discoveries of new magnetic effects at low temperature require extreme carefulness. X-ray
magnetic circular dichroism (XMCD) or magneto optical methods are more reliable methods to
detect the presence of ferromagnetism instead of conventional magnetometry measurements [65-
66]. It was shown that nearly all substrate show small ferromagnetic like signal from room
temperature to low temperatures in conventional magnetometry measurements like SQUID [11].
The author has shown that MgO in particular is more magnetic than any other substrate. This can
be seen from the magnetization as a function of temperature in fig. 6.4. In this work PPMS is
used to detect the presence of ferromagnetism at low temperature. If the low temperature
ferromagnetic phase is present, then as the temperature is lowered the moment should increase.
Therefore our first experiment was to put a magnetic material and substrate to see the behavior at
low temperature; this is shown in the figure 6.5. It can seen from the figure that the
ferromagnetic thin film on MgO shows the same signature as the substrate, at around 50 K the
moment show a signature bump due to the presence of molecular oxygen as it undergoes
antiferromagnetic transition at 43 K [67]. Then we measured the four samples of Mn rich Fe50-
xMnxPt50, the moment versus the temperature is shown in the figure 6.6. Remarkably they
showed same behavior as the ferromagnetic material like Fe50Pt50 and the substrate. In other
85
words the instrument couldn’t distinguish between a substrate or magnetic/non-magnetic
material in a moment versus temperature measurement. So we decided to measure the M(H)
curves at different temperatures from 300K to 100 K. Figure 6.7 shows the hysteresis loops
(M(H)) for the MgO (100) substrate, where the moment is not a perfectly linear as expected for a
diamagnetic material as the slopes at positive and negative field are different, which indicate the
presence of small magnetic signals. M(H) loops for the four samples at different temperatures
showed (Figures 6.8 and 6.9) no temperature dependence, the low temperature mixed
ferromagnetic phase was expected to show a temperature dependence in M(H) loops.
86
Fig. 6. 4. Magnetization as a function of temperature for different substrates (Ref. 11)
87
100 200 300-2
0
2
4
Applied field: 6T
mo
men
t (m
emu
)
Temperature(K)
MgO (100)
100 200 300-2
0
2
4
x=0
Fe50
Pt50
/ MgO (100)
mo
men
t (m
emu
)
Temperature
FIG.6.5. Moment vs. Temperature measurement for ferromagnetic thin films on (top) MgO
substrate alone, (bottom) Fe50Pt50 film on MgO
88
50 100 150 200 250 300-2
0
2
4
Applied field: 6T
Fe50-x
MnxPt
50
mo
men
t (m
emu
)
Temperature(K)
x=40
x=42
x=44
x=46
FIG. 6.6. Moment vs. Temperature measurement for the substrate and Mn rich Fe50-x Mnx Pt50
thin films.
89
-20 0 20
-0.4
-0.2
0.0
0.2
0.4
mo
men
t (m
emu
)
Applied field (kOe)
300 K
260 K
220 K
200 K
180 K
160 K
140 K
100 K
FIG. 6.7. Hysteresis loop (M(H) curves) at different temperatures for MgO (100) substrate.
90
-20 0 20
-0.04
-0.02
0.00
0.02
0.04
mom
ent
(mem
u)
Applied field (kOe)
300 K
260 K
220 K
200 K
180 K
160 K
140 K
100 K
Fe50-x
MnxPt
50
x=40
-20 0 20
-0.04
-0.02
0.00
0.02
0.04
mom
ent
(mem
u)
Applied field (kOe)
300 K
260 K
220 K
200 K
180 K
160 K
140 K
100 K
Fe50-x
MnxPt
50
x = 42
FIG. 6.8 Hysteresis loop at different temperatures for x= 40 (top) and x=42 (bottom)
91
-20 0 20
-0.04
-0.02
0.00
0.02
0.04
x=46
Fe50-x
MnxPt
50
mom
ent
(mem
u)
Applied field (kOe)
300 K
260 K
220 K
200 K
180K
160 K
140 K
100K
FIG. 6.9. Hysteresis loop at different temperatures for x= 44 (top) and x=46 (bottom)
-20 0 20
-0.04
-0.02
0.00
0.02
0.04
x=44
Fe50-x
MnxPt
50
mo
men
t (m
emu
)
Applied field (kOe)
300 K
260 K
220 K
200 K
180 K
160 K
140 K
100 K
92
The hysteresis behavior of the substrate at different temperature is shown in the figure
6.7. The moment of thin film can be found by the following equation
substratesubstratefilmfilm mmm
where
mfilm+ substrate = moment of the total film and substrate normalized with respect to its area
msubstrate= moment of the substrate normalized to its area
The expected moment expected for the predicted low temperature phase can be calculated using
tAMm ected exp
where M is the magnetization, A is the area of the sample, t is the thickness of the thin film which
is 45 nm. The magnetization of the Fe50Pt50 is 1000 emu/cc, if we set an upper limit of 1000
emu/cc and a lower limit of 500 emu/cc, the expected moment for the four samples is between 3
to 6 x 10-4
emu. The magnetization of the substrate Msubstrate is given by
mBM msubstrate
where B is the applied field (3 T for M(H) loops and 6 T for M (T) curves), χm is the mass
diamagnetic susceptibility and m is the mass of the substrate. The moment from the substrate for
our experiment is of the order of 4 x 10-3
emu. Thus the moment of the substrate signal is about
ten times stronger than the expected moment for the low temperature ferromagnetic phase. As
the substrate signal overwhelms the expected signal it is very hard to detect ferromagnetism at
low temperature.
mexpected/ thin film<< msubstrate , for low temperature measurements this is not critical for Fe rich
Fe50-x MnxPt50 ( x < 0.5 ) with large magnetic signals but when measuring Mn rich Fe50-x MnxPt50
( x > 40 ) with small magnetic signals ( ~ 10-4
emu ) the presence of small artifacts will be
comparable to the signal from the film. The hysteresis loops for the four samples (fig. 6.8 and
93
6.9) shows the presence of magnetic signal but the measured moment is ten times smaller than
the expected moment. The results are summarized in the table XI.
X
Area
cm2
(substrate
0.11 )
mexpected
lower limit
( at 200K)
emu
mexpected
higher
limit
( at 200K)
emu
mobserved
(200 K)
emu
mobserved
(100 K)
emu
Observed
Magnetization
(100 K)
emu/cc
40
0.13 3.12 x 10-4
6.24 x 10-4
1.2 x 10-5
1.4 x 10-5
22.3
42
0.15 3.6 x 10-4
7.2 x 10-4
1.94 x 10-5
2.1 x 10-5
29.2
44
0.12 2.88 x 10-4
5.76 x 10-4
0.36 x 10-5
0.55 x 10-5
9.54
46
0.062 1.49 x 10-4
2.98 x 10-4
2.25 x 10-5
1.82 x 10-5
61.2
Table XI. Expected moment and the observed moment for the Mn rich Fe50-x Mnx Pt50 for
different compositions.
94
CHAPTER 7
CONCLUSION
Ordered epitaxial intermetallic thin films on a plane sapphire and MgO (100) have been
fabricated using sputtering technique and characterized using various techniques. It has been
shown that the deposition temperature causes the diffusion of seed layer to the top surface which
may contribute to abnormal magnetic properties. Determination of optimal low deposition
temperature and direct deposition were one of the possibilities to avoid diffusion. Fe rich end of
the Fe50-x Mnx Pt50 phase diagram was studied. The films were grown on a plane sapphire (11-20)
and MgO (100). These thin films show striking difference in their magnetic properties. The thin
films grown on sapphire agree with the previous experimental results (AFM phase) while the
thin films grown on MgO (100) shows behavior opposite to what was predicted recently. Thus it
implies that AFM alignment of the Mn moment decreases the magnetic properties, while the FM
alignment of Mn Moments shows increase in the magnetic properties.
We have shown that this increase in magnetic properties coincide with the tetragonal
distortion. At x=12 the coercivity has increased by 46.4 % when compared to Fe50Pt50. Thus for
the concentrations at the vicinity of x=12.5 the magnetic properties of Fe50-xMnxPt50 increases,
which agrees with the previous predictions. We conclude that high temperature and longer
annealing time contribute to the structural transformations in the Fe50-xMnxPt50 thin films on
MgO. For the Mn rich side of the Fe50-x Mnx Pt50 phase diagram we attempted to verify the recent
95
prediction of low temperature ferromagnetic phase. The low temperature hysteresis measurement
for the samples at the various concentrations x=40, 42, 44 and 46 showed the presence of
ferromagnetism but we encountered the experimental difficulties to measure these magnetic
moments as the predicted magnetic moment is the order of 10-4
emu, ten times smaller than the
substrate background of our magnetometry measurements. The standard procedure of sample
handling for the magnetic measurements induce the appearance of magnetic signals of the order
of 10-4
emu which is of the same magnitude of the predicted low temperature magnetic phase,
while these signals are negligible when measuring magnetic materials with higher saturation
magnetization like Fe rich end of the Fe50-xMnx Pt50 (x < 15) but in case low temperature
magnetic phase the conventional magnetometry measurements pose incredible difficulties when
measuring low magnetic moments of the order of ~ 10-4
emu as the substrate and other spurious
signals overwhelm the actual signal. Magneto optical methods and XMCD are more reliable
methods for measuring the predicted new low temperature ferromagnetic phase in Fe50Mn50-xPt50
system.
96
REFERENCES
1. K. Shibata, Materials Transactions, Vol. 44, No.8, 1542 (2003)
2. D. Weller, G. Parker, O. Mosendz, E. Champion, B. Stipe, X. Wang, T. Klemmer, G. Ju, and
A. Ajan, IEEE Trans. Magn. 50, 3100108 (2014)
3. B. S. Pujari, P. Larson, V. P. Antropov, and K. D. Belashchenko, Phys. Rev. Lett. 115,
057203 (2015)
4. T. Burket, O. Eriksson, S. I. Simak, A. V. Ruban, B. Sanyal, L. Nordstrom, and J. M. Willis,
Phys. Rev. B 71, 134411 (2005)
5. Jihoon Park, Yang-Ki Hong, Seong-Gon Kim, Li Gao, and Jan-Ulrich Thiele, J. Appl. Phys.
117, 053911 (2015)
6. N. Zotov, R. Hiergeist, A. Savan, A. Ludwig, Thin Solid Films 518, 4977 (2010)
7. M. Schilling, P. Ziemann, Z. Zhang, J. Biskupeck, Beilstein J. Nanotechnol. 7, 591 (2016)
8. K. M. Hyie and I. I. Yaacob, Proc. of World Cong. on Eng. July 2-4, 2008, London, U.K.
9. K. Yano, V. Nandwana, N. Poudyal, C. Rong, and J. P. Liu, J. Appl. Phys. 104, 013918
(2008)
10. C. L. Platt, K. W. Wierman, E. B. Svedebrg, R. V. Veerdonk, J. K. Howard, A. G. Roy and
D. E. Laughlin, J. Appl. Phys. 92, 10 (2002)
11. M. Khalid, A. Setzer, M. Ziese, P. Esquinazi, D. Spemann, A. Poppl and E. Goering, Phys.
Rev. B 81, 214414 (2010)
12. P. Auger, Compt. Rend. 180, 65 (1925)
13. L. A Harris, J. Appl. Phys. 39, 1428 (1968)
14. P. Palmberg, W, Bohn, G. K, and Tracey, J. C, Appl. Phys. Lett. 15, 254 (1969)
15. M. P. Seah and W. A. Dench, Surface and Interface Analysis 1, 2 (1979)
97
16. P. A. Maksym, J. L. Beeby, Surface Science, 110, 2 (1981)
17. S. Ino, Japn. J. Appl. Phys. 16, 6 (1977)
18. W. Braun, Applied RHEED: reflection high-energy electron diffraction during crystal
growth, Springer, New York (1999)
19. P. A. Maksym, Surface Science, 149, 1 (1985)
20. W. R. Grove. "VII. On the electro-chemical polarity of gases". Philosophical Transactions of
the Royal Society. 142(I), (1852)
21. J. Plucker, “Observations on the Electrical Discharge Through Rarefied Gases,” The
London, Edinburgh and Dublin Philosophical Magazine, 16, 409 (1858)
22. Rudolf Reinecke and. Wilhelm Berkhardt, USP 2,157,478
23. B. D. Cullity and S. R. Stock , Elements of X-ray diffraction, Prentice Hall, NJ (2001)
24. http://physics.nist.gov/PhysRefData/FFast/html/form.html
25. N. T. Nam, W. Lu, and T. Suzuki, J. Appl. Phys. 105, 07D708 (2009)
26. I. Suzuki, T. Koike, M. Itoh, T. Taniyama, and T. Sato, J. Appl. Phys. 105, 07E501 (2009)
27. P. Mani, Krishnamurthy VV, Robertson JL, Klose F, Mankey GJ, J. Appl. Phys. 99, 08C109
(2006)
28. D. Lott, F. Klose, H. Ambaye, G. J. Mankey, P. Mani, M. Wolff, A. Schreyer, H. M.
Christen, and B. C. Sales, Phys. Rev. B 77, 132404 (2008)
29. T. Saerbeck, F. Klose, D. Lott, G. J. Mankey, Z. Lu, P. R. LeClair, W. Schmidt, A. P. J.
Stampfl, S. Danilkin, M. Yethiraj, and A. Schreyer, Phys. Rev. B 82, 134409 (2010)
30. W. Lu, N. T. Nam, and T. Suzuki, J. Appl. Phys. 105, 07A904 (2009).
31. H. Lee, The magnetic and chemical structural property of the epitaxially grown multilayered
thin film, PhD dissertation (2012), University of Alabama
32. P. Mani, Probing spin ordering in Fe-Pt based antiferromagnetic films using neutron
diffraction, PhD dissertation (2005), University of Alabama
33. Zhihong Lu, Magnetic anisotropy graded media and Fe-Pt alloy thin films, PhD dissertation
(2009), University of Alabama
34. J. U. Thiele, L. Folks, M. F. Toney, and D. K. Weller, J. Appl. Phys. 84, 5686 (1998).
98
35. A. C. Sun, P. C. Kuo, J.-H. Hsu, H. L. Huang, and J. M. Sun, J. Appl. Phys. 98, 076109
(2005).
36. S. P Bennett, H Ambaye, H. Lee, P. LeClair, G.J. Mankey, V. Lauter, Sci. Rep. 5, 9142
(2015)
37. A.T. Heitsch, D.C. Lee, B.A. Korgel, J. Phys. Chem. C, 114, 2512 (2010)
38. S. Maat, A. J. Kellock, D. Weller, J. E. E. Baglin, Eric E. Fullerton, J. Magn. Magn.
Materials, 265, 1 (2003).
39. D. Weller, G. Parker, O. Mosendez, E. Champion, B. Stipe, X. Wang, T. Klemmer, G. Ju,
and A. Ajan, IEEE Trans. Magn. 50, 3100108 (2014).
40. J. A. Aboaf, S. R. Herd, and E. Klokholm, IEEE Trans. Magn. 20, 1642 (1984)
41. B.M. Lairson, M.R. Visokay, R. Sinclair, and B.M. Clemens, Appl. Phys. Lett. 62, 639
(1993)
42. K. R. Coffey, M. A. Parker, and J.K. Howard, IEEE Trans. Magn. 31, 2737 (1995)
43. S. R. Lee, S. Yang, Y. K. Kim, and J. G. Na, Appl. Phys. Lett. 78, 4001 (2001)
44. M. H. Hong, K. Hono, and M. Watanabe, J. Appl. Phys. 84, 4403 (1998)
45. S. Okamoto, N. Kikuchi, O. Kitakami, T. Miyazaki, Y. Shimada, and K. Fukamichi, Phys.
Rev. B. 66, 024413 (2002)
46. T. Shima, K. Takanashi, Y. K. Takahashi and K. Hono, Appl. Phys. Lett. 81, 1050 (2002)
47. Zhihong Lu, M. J. Walock, P. LeClair, W. H. Butler, and G. J. Mankey, J. Vac. Sci. Technol.
A 27, 1067 (2009)
48. J. Buschbeck, I. Opahle, M. Richter, U. K Rolber, P. Klaer, M. Kallmayer, H. J Elmers,
G.Jakob, L. Schultz and S. Fahler, Phys. Rev. Lett. 103, 216101 (2009)
49. P. Ravichandran, A. Kjekshus, H. Fjellvag, P. James, L. Nordstrom, B. Johansson, and
O. Eriksson, Phys. Rev. B 63, 144409 (2001)
50. A. Sakuma. J. Phys. Soc. Jpn. 63, 3053 (1994)
51. D. B. Xu, J.S. Chen, T.J. Zhou, and G. M. Chow, J. Appl. Phys. 109, 07B747 (2011)
52. G. Meyer and J. U. Thiele, Phys. Rev. B 73, 214438 (2006)
99
53. C.J. Sun, D.Xu, S.M. Heald, J. Chen, and G.M. Chow, Phys. ReV. B 89, 140408 (2011)
54. F.T. Yuan, L.S. Lee, S.K. Chen, T.S. Chin, W.C. Chang, J. Magn. Magn. Materials 272,
1164 (2004)
55. R. Cuadrado, Kai Liu, Timothy J. Klemmer, and R. W. Chantrell, Appl. Phys. Lett. 108,
123102 (2016)
56. A. Z. Menshikov, V. P. Antropov, G. P. Gasnikova, Yu. A. Dorofeyev, and V.A. Kazantsev,
J. Magn. Magn. Materials 64, 159 (1987)
57. Markus E. Gruner and Peter Entel, Beilstein J. Nanotechnol., 2, 162–172 (2011)
58. Marcios M. Soares, M. Santis, Helico C.N. Tolentino, Aline Y. Ramos, Mohammad El
Jawad, and Yves Gauthier, F. Yildiz and M. Przybylski, Phys. Rev. B. 85, 205417 (2012)
59. A. Brinkman, M. Hujiben, M. Van Zalk, J. Hujiben, U. Zeitler, J.C. Maan, W. G.Van Der
Weil, G. Rijnders, D. H. A. Blank, and H. Hilgenkamp, Nature Mater. 6, 493 (2007)
60. S. D. Yoon, Y. Chen, A. Yang, T. L. Goodrich, X. Zuco, D. A. Arena, K. Ziemer, C.Vittoria,
V. Harris, J. Phys. Condens. Matter 18, L355 (2006)
61. A. Sudarsen, R. Bhargavi, N. Rangarajan, U. Siddesh, and C.N.R. Rao, Phys. Rev. B. 74,
161306 (2006)
62. A.Shipra, A. Gomathi, A. Sudarsen, and C. N. R. Rao, Solid State Commun. 142, 685 (2007)
63. I. K. Schuller and Y. Brynserade, Solid state Commun. 30, 75 (2005)
64. M. A. Garcia, E. Fernandez Pinel, J. Valenta, A. Quesada, V. Bouzas, J. F. Fernandez, J. J.
Romero, M. S. Martin Gonzalez, and J. L. Costa Kramer, J. Appl. Physics (2009)
65. J. Stohr, J. Magn. Magn. Mater. 200, 470 (1999)
66. Z. Vager, I. Carmeli, G. Leitus, S. Reich and R. Naaman, J. Phys. Chem. Solids 65, 713
(2004)
67. S. Gregory, Phys. Rev. Lett. 40, 723 (1978)