correlation between the structural, optical, and magnetic
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Correlation Between the Structural, Optical, and Magnetic Properties
of CoFeB and CoFeB Based Magnetic Tunnel Junctions
Upon Laser or Oven Annealing
Fakultät für Naturwissenschaften
der Technischen Universität Chemnitz
Dissertation zur Erlangung des akademischen Grades
eingereicht am 7. Januar 2020
vorgelegt von M. Sc. Apoorva Sharma
geboren am 02. Mai 1987 in Bikaner, Indien
Gutachter:
Prof. Dr. Dr. h. c. Dietrich R.T. Zahn
Prof. Dr. Georgeta Salvan
“ Every search begins with beginner’s luck. And every
search ends with the victor’s being severely tested.”
― Paulo Coelho
Bibliografische Beschreibung
Sharma, Apoorva
Correlation Between the Structural, Optical, and Magnetic Properties of CoFeB and
CoFeB Based Magnetic Tunnel Junctions Upon Laser or Oven Annealing
Technische Universität Chemnitz Dissertation, 2020
Diese Dissertation befasst sich mit der Untersuchung der maßgeblichen
Herausforderungen der heutigen TMR-Präparation (tunneling magnetoresistance) für
beispielsweise Magnetfeldsensor- oder auch Speichertechnologie (MRAM – magnetic
random access memory).
Im ersten Teil der Arbeit werden die elektronischen, strukturellen und magnetischen
Eigenschaften der ferromagnetischen Elektrode eines typischen magnetischen
Tunnelkontaktes, z.B. CoFeB, erforscht, wobei spektroskopische Ellipsometrie, magneto-
optische Spektroskopie, Röntgendiffraktometrie und Messverfahren für den spezifischen
elektrischen Widerstand zum Einsatz kommen. Weiterhin wurde der Einfluss der
Temperatur einer thermischen Behandlung auf die optischen und magneto-optischen
Merkmale untersucht, wobei eine starke Korrelation zwischen den beobachteten
spektralen Merkmalen und der Kristallisation von CoFeB nachgewiesen wurde. Die
(magneto-) optische Spektroskopie bietet somit eine zerstörungsfreie und besonders
sensitive Validierungsmethode für die Dünnschichtkristallisation, die in die moderne
CMOS Herstellungstechnologie integriert werden kann.
Der zweite Teil der Arbeit befasst sich mit dem lokalen Tempern unter Verwendung eines
fokussierten Laserstrahls, mit dem Ziel die Referenzmagnetisierung in einem
magnetischen Tunnelkontakt definiert einzustellen und die Wirkung der hierfür
notwendigen thermischen Behandlung auf die übrigen Schichten im Schichtstapel zu
untersuchen. Hierzu wurden zahlreiche Parameter für das laserbasierte lokale Tempern
variiert, um die optimale Austauschfeldstärke im magnetischen Referenzsystem
einzustellen, idealerweise ohne den gegebenen Schichtstapel zu schädigen. Schließlich
wurde der Einfluss des laserbasierten Temperns (als auch des Ofentemperns) auf die
Unversehrtheit der Schichten und Grenzflächen, insbesondere auf die Diffusion
verschiedener Elemente, mittels Röntgen-Photoemissionsspektroskopie untersucht.
Keywords
Exchange bias, Laser irradiation, Magnetic tunnel junction, Magneto-optical Kerr effect
spectroscopy, Magnetometry, Spectroscopic ellipsometry, Spintronic, Tunnel
magnetoresistance and 3d transition metal boride.
Table of content
Chapter 1: Introduction ................................................................................................. 1
1.1 This thesis ............................................................................................................. 2
Chapter 2: Theory.......................................................................................................... 5
2.1 Magnetism ............................................................................................................. 5
2.2 Spintronics ............................................................................................................15
2.3 Optical spectroscopy ............................................................................................18
Chapter 3: Experimental ..............................................................................................31
3.1 Sample preparation ..............................................................................................31
3.2 Annealing ..............................................................................................................32
3.3 Measurement techniques .....................................................................................35
Chapter 4: 3d-transition metal boride layers: structural, electronic, and magnetic
properties ......................................................................................................................46
4.1 Introduction ...........................................................................................................46
4.2 Thick films .............................................................................................................47
4.3 Thin films ..............................................................................................................60
4.4 Conclusion ............................................................................................................65
Chapter 5: Setting exchange bias using laser vs oven annealing techniques .......67
5.1 Introduction ...........................................................................................................67
5.2 Magnetisation reversal of complex MTJs layer stack ............................................69
5.3 Application of FORC analysis: from single CoFeB layer to MTJ layer stack .........81
5.4 Potential of direct-write laser annealing technique................................................84
5.5 Conclusion ............................................................................................................86
Chapter 6: Exchange bias and diffusion processes in laser annealed CoFeB/IrMn
thin films .......................................................................................................................88
6.1 Introduction ...........................................................................................................88
6.2 Magnetometry investigations ................................................................................89
6.3 XPS-depth profiling ...............................................................................................90
6.4 Structural analysis ................................................................................................94
6.5 Topographic characterisation ...............................................................................95
6.6 Conclusion ............................................................................................................96
Chapter 7: Summary and outlook ...............................................................................97
Appendix A .................................................................................................................100
Bibliography ...............................................................................................................101
List of figures .............................................................................................................111
List of tables ...............................................................................................................118
Abbreviations .............................................................................................................119
Erklärung ....................................................................................................................120
curriculum vitae .........................................................................................................121
Scientific contributions .............................................................................................122
Acknowledgements ....................................................................................................126
Introduction Chapter 1
1
Chapter 1: Introduction
High-precision magnetic field sensors have become indispensable in a wide variety of
modern scientific equipment and industrial devices. Their fields of applications range from
the medical and health care sector to the technology of everyday life. For medical
purposes, for example, the magnetic signature of a tracer particle may be tracked in vivo
along all three spatial dimensions as it passes the intestinal tract of a patient, thereby
revealing any anomalies or functional disorders1,2. Also, direct monitoring of the
biomagnetic signals from the heart and the brain has been realised by ultrahigh-precision
magnetic field sensors1,3,4. In automobiles and industrial equipment, to provide another
example, magnetic field sensors are employed for reliable, high-precision, length, angle
and position measurements under challenging environmental conditions5,6. In the field of
power generation and transport, especially within the emerging context of dynamic,
distributed “green” power generation by wind and water, trustworthy contact-free
measurements of electric currents using magnetic field sensors are of great importance7.
The severity of such measurements increases manyfolds in the internet of things (IoT)
applications, where ever so small leakage currents need to be detected by adequate
sensor technologies to save energy8,9. Furthermore, magnetic field sensors have recently
entered the market of mobile communication as their increasingly high sensitivity now
allows smartphone navigation based on the electronic reading of the earth’s natural
magnetic field10.
Aforementioned fields of application, as well as many others, have in common that they
are experiencing a continuous replacement of conventional Hall and anisotropic
magnetoresistance (AMR) sensors by more precise magnetoresistive components. In
particular, the so-called magnetoresistive effects, i.e. the giant (GMR)11 and tunnelling
(TMR)12 magnetoresistance, in which besides the electron charge also its spin is used as
an information carrier, exhibit much higher signal amplitudes and sensitivities compared
to Hall and AMR technology13. The application of the GMR effect already led to
revolutionary progress in hard disk drive technology with such an impact that A. Fert and
P. Grünberg were awarded the Nobel Prize in Physics in 2007. In the last two decades,
the GMR-based technology has been slowly replaced by TMR-based devices, particularly
after the introduction of the first TMR-based read-heads in 2005. Their extremely high
signal output at room temperature, along with their very low energy consumption and their
Introduction Chapter 1
2
much higher miniaturisation capabilities as compared to the sensors based on the Hall
effect, AMR or GMR effect, promises a significant improvement of sensor quality. The
remarkable achievements in the evolution of the magnetoresistive sensor are shown in
Figure 1.
Figure 1. The milestones in the evolution of magnetoresistive devices.
1.1 This thesis
The research work presented in this thesis was performed as a part of a joint project
between the Semiconductor Physics group at Chemnitz University of Technology (HLPH-
TUC), Back-End-of-Line group at the Fraunhofer Institute for Electronic Nano Systems
(BEOL-Fraunhofer ENAS) and the Laser Microtechnologies at the Laser Institute of the
Hochschule Mittweida (LHM). The project “Interfacial perpendicular magnetic
anisotropy for next-generation monolithic 3D TMR sensors” was financially
supported by the Deutsche Forschungsgemeinschaft with the Grant No.: 282193534. The
objective of this project was the monolithic integration of a three-dimensional magnetic
field sensor system based on TMR in CoFeB / MgO / CoFeB for quantitative
measurements of the magnetic field. The primary emphasis was laid on a higher
sensitivity, minimum signal hysteresis, low overall power consumption as well as a high
degree of miniaturisation. Although the project was tightly knitted among the three project
partners, the work packages can be broadly classified as sample/material
characterisation, microfabrication and laser annealing performed by HLPH-TUC, BEOL-
Fraunhofer ENAS and LHM, respectively.
Introduction Chapter 1
3
The performance of TMR sensors vastly depends on the thermal treatment for the
crystallisation of the ferromagnetic electrodes as well as the MgO barrier and for setting
the reference magnetisation via exchange bias. The typical method for thermal treatment
involves vacuum oven annealing in the presence of a magnetic field. However, this
inherits some limitations, such as non-selective alignment of the magnetisation direction
of the reference layer for each sensor on a wafer and time consumption. Often, due to the
limited thermal budget, the vacuum oven annealing is incompatible with CMOS
technologies. A laser-based annealing approach overcomes the shortcomings of the
conventional annealing method. This method allows not only the selective magnetisation
alignment of the reference layer of the single sensor but also reduces the annealing time.
This thesis focuses on studying the various aspects of laser annealing on the magnetic
and non-magnetic layers comprised of CoFeB / MgO / CoFeB magnetic tunnel junctions
(MTJ). Properties such as magnetisation, exchange bias, crystallisation and surface
morphology in conjunction with varying laser exposure parameters were investigated.
Additionally, (magneto-) optical spectroscopy methods were used to understand the
electronic structure of the ferromagnetic layers. For benchmarking purposes, studies were
made using the conventional vacuum oven annealing method.
Chapter 2 of this thesis describes the necessary theoretical background required to
explain and understand the results presented in this work. The chapter begins with a brief
overview of the fundamentals of magnetism and explains characteristic features of
ferromagnets, namely, magnetic hysteresis, coercive field, remanence magnetisation etc.
The following sections describe the phenomena that are important for this research, such
as magnetic anisotropy, exchange bias and tunnel magnetoresistance in-depth. It also
sheds light on the advantages of optical spectroscopic techniques and how these can be
exploited to understand material properties. In chapter 3, the sample preparation and
characterisation techniques that have been used in the framework of this thesis are
presented. The basic principle and the technical details for the equipment used are
explained. Additional specifications regarding the individual experiments are discussed
along with the experiment in each following chapter. Chapter 4 presents the optical and
magneto-optical response of the materials investigated in this thesis. Additionally, the
investigation of the crystallisation of CoFeB ferromagnetic layers using (magneto-) optical
methods along with X-ray diffraction as well as electrical four-point probe method is
discussed. Chapter 5 brings a proof-of-concept that the laser annealing is a suitable
Introduction Chapter 1
4
method for heat-treating the magnetic tunnel junction devices. It also aims to establish the
optimal laser annealing parameters based on the magnetic properties of the MTJ layer
stack and presents the study with different seed layers, which were tested in order to
improve the exchange bias fields. Chapter 6 deals with the more fundamental
investigation of the influence of laser annealing on a reference layer stack. The main focus
here is to characterise and understand the diffusion and the type of migrating species in
the layer stack induced by the laser annealing. Finally, a summary of all the studies is
compiled in chapter 7.
Theory Chapter 2
5
Chapter 2: Theory
This chapter lays the theoretical foundation for the interpretation of the results presented
in this thesis. It also serves the purpose to familiarise the reader with the scientific terms
used throughout this thesis. At first, the fundamentals of magnetism are explained,
followed by a brief review about spintronics and the principles behind the TMR effect,
along with the development road map of TMR sensors. The second part draws attention
toward the light-matter interaction with a focus on (magneto-) optical spectroscopic
techniques for characterisation and analysis of magnetic materials. A section of this
chapter is published (with minor changes) by Royal Society of Chemistry ISBN-13:
9781839162107, ISBN-10:1839162104.
2.1 Magnetism
The term magnetism in ordinary life is associated with the physical property of
ferromagnetic materials to be attracted by a magnetic field. For instance, the spontaneous
magnetisation of the lodestone (magnetite) is already known to humanity for several
centuries. However, it was only in the early 19th century that Hans-Christian Oersted
bridged the magnetism with electricity and triggered more systematic studies of
magnetism. Later on, the discovery of the Faraday and Kerr effect by Michael Faraday
and John Kerr revolutionised the concept of interaction between light and magnetic
materials14. A more detailed description of these effects is presented in section 2.3.5. The
discovery of Faraday and Kerr effect laid the foundation of the famous Maxwell equations
developed by James Clark Maxwell. In the following, the mathematical equations relevant
for this work along with their units, are explained. The units associated with the physical
quantities are shown in appendix A.
The relation between the magnetic induction (B ) and the applied magnetic field (H ) in
vacuum is given by14,15
= 𝜇0 2-1
μo is the magnetic permeability of the vacuum, also known as the magnetic constant.
Theory Chapter 2
6
In order to describe the magnetic field inside matter, a third vector term called
magnetisation (M ) was introduced. The magnetisation M of a material can be defined as
the vector sum of all magnetic moments (m ) located within the material volume (V):
=1
𝑉∑ 2-2
The total magnetic induction will account for the contribution of the magnetisation
accordingly:
= 𝜇0(𝐻 + 𝑀) 2-3
Hence, the relation between the applied magnetic field and the magnetisation is given by
= 𝜒 2-4
where, χ is the magnetic susceptibility, a dimensionless quantity.
Based on the magnetic susceptibility (or permeability), all elements and their alloys can
be classified into three categories, namely as diamagnetic, paramagnetic or magnetically
ordered (ferromagnetic, antiferromagnetic, ferrimagnetic, etc.). For diamagnetic
materials, the magnetic susceptibility is negative (χ < 0). In this case, the atomic
magnetic moments in the absence of the magnetic field are zero and an external magnetic
field H induces atomic magnetic moments, which leads to a reduction of B . On the other
hand, paramagnetic materials have χ > 0; an intrinsic atomic magnetic moment is
already present in the absence of the external magnetic field, but the atomic moments do
not interact with each other. When an external magnetic field H is applied, the magnetic
moments tend to align with the applied field, leading to a linear increase of B with H .
Paramagnetic materials consist of atoms or ions with unpaired electrons, which results in
an individual magnetic moment associated with the atoms or ions. However, these
individual magnetic moments do not interact with each other and have no long-range
ordering, resulting in a zero net magnetisation. Hitherto, for the ease of presentation, the
vector notations for these quantities will be dropped.
Ferromagnetic (FM) materials show magnetic order below a critical temperature, with χ
being several orders of magnitude larger than for the paramagnetic state. Such materials
Theory Chapter 2
7
exhibit magnetic hysteresis and persistent magnetisation at zero fields, the so-called
remanent magnetisation or remanence. Due to these properties, ferromagnetic materials
have gained a lot of attention since centuries and are still an interesting subject for current
research activities. Another peculiar magnetic ordering case is the so-called
antiferromagnetism (AFM), where the neighbouring atomic magnetic moments are
coupled antiparallel to each other, leading to zero remanence. The exchange bias
phenomenon often observed in AFM/FM layered systems represents a significant part of
investigations in this thesis, as the exchange bias not only determines the directional
sensitivity but also improves the signal to noise ratio in magnetic field sensors. A
schematic representation of the classes of magnetically ordered materials (including
cases that are not addressed in this thesis) is shown in Figure 2.
Figure 2. The classification of magnetic materials [Image adapted from14].
2.1.1 Ferromagnetism
As mentioned above, ferromagnetic materials are materials that possess a spontaneous
magnetisation below the Curie temperature. The spontaneous magnetisation is the net
magnetisation that exists inside a magnetised microscopic volume in the absence of an
Theory Chapter 2
8
external magnetic field. The Curie temperature (TC) is the temperature above which
ferromagnetic materials lose their spontaneous magnetisation and become
paramagnetic. Iron (Fe), Cobalt (Co), and Nickel (Ni) are the three elementary
ferromagnetic metals at room temperature. Furthermore, some other rare-earth metals
(4f-elements) also have significantly contributed to magnetism, especially when alloyed
with transition metals. A very prominent example is NdFeB, which is used in permanent
magnets.
Two theories of ferromagnetism have been successful in explaining many of its properties:
the Weiss theory16 and the Stoner band theory of ferromagnetism17. In classical physics,
Weiss postulated that an internal “molecular field” acts in ferromagnetic materials to align
the magnetic moments parallel to each other. Below TC, the molecular field is strong
enough to keep a magnetisation even in the absence of an externally applied field. At high
temperatures, the thermal energy KBT, is larger than the energy of the molecular field,
resulting in the random orientation of the magnetic moments and thus paramagnetic
behaviour with a small but positive susceptibility. This theory has proven to describe the
experimentally observed Curie-Weiss behaviour for the temperature dependence of the
susceptibility for many ferromagnetic materials in the paramagnetic region above the
Curie temperature. However, there are two limitations to the Weiss theory. First, it
assumes that the number of magnetic dipoles below and above the Curie temperature is
the same (ferro to para-magnetic phase transition), while this is in contrast to experimental
evidence. Second, the magnetic dipole moment is an integer number of the Bohr
magneton, which is not observed experimentally18. To explain this discrepancy, the Stoner
band theory of ferromagnetism is needed.
The band theory of ferromagnetism was first proposed by E. Stoner19 and then
independently by J. Slater20–22. It has been successful in explaining non-integer values of
atomic magnetic moments and predicting some aspects of the magnetic behaviour of the
3d metals and alloys evaluated from the imbalance of spin-up and spin-down electrons
by integrating the respective spin-polarised density of states15.
𝑚 = 𝜇𝐵(𝑛↑ − 𝑛↓) 2-5
where, B = 9.27×10-24 J·T-1 is the Bohr magneton and n↑ denotes the number of spinup
and n↓ the number of spindown electrons.
Theory Chapter 2
9
The Stoner model assumes that the magnetisation occurs due to the exchange interaction
between the electrons with the same spin orientation with neighbouring atoms in a crystal
lattice, thus causing an effective imbalance in spin-up and spin-down electron density.
This difference in concentration of spin-up and spin-down electrons results in a band
splitting known as exchange splitting (Δ). The existence of ferromagnetism can then be
expressed by the Stoner criterion15.
𝑆𝐷(𝐸𝐹) ≥ 1 2-6
here, S is the Stoner constant and D(EF) is the density of states.
Thus, the ferromagnetism can be explained by an electron density of states, where the
imbalance due to spin orientations of electrons is addressed by relocating the
corresponding energy densities with respect to each other. The difference between the
upper edge of the 3d band of majority spin electrons and the maximum energy occupied
by an electron at 0K (also known as Fermi energy EF) is the Stoner gap. Figure 3 a shows
the Stoner model for ferromagnetic metals with a 3d shell. The shaded and unshaded
areas represent the occupied and unoccupied states, respectively. The band with higher
occupancy of electrons is called the majority band, and the corresponding electrons are
named majority spin electrons, while the term minority refers to the band with lower
occupancy and the related electrons.
Figure 3. The Stoner model of ferromagnetic metals illustrated for the 3d shell, and nomenclature used for the band description of magnetism (a.). Occupied electron states below the Fermi energy EF are shaded, unoccupied states above EF are shown unshaded. Hext, m and Δ denote the external magnetic field, magnetic moment and exchange splitting, respectively. The calculated density of states of Fe, Co, Ni, and Cu (b.) [Images are taken from15].
Theory Chapter 2
10
2.1.2 Antiferromagnetism
The theory of ferromagnetism in the 3d-transition metals, Fe, Co and Ni, is based on the
postulation of positive exchange interactions between neighbouring atoms/ions. However,
this assumption does not explain the paramagnetic and antiferromagnetic (negative
exchange) behaviour for remaining 3d metals (Sc, Ti, V, Cr, Mn). In the year 1936, Louis
Néel developed the concept of antiferromagnetism14. He observed that below a certain
critical temperature, the atomic magnetic moments in such metals/alloys are arranged
alternately parallel and antiparallel. Above this temperature, the moments are disordered
similar to ferromagnets above TC. This critical temperature was later named after him as
Néel temperature (TN) of antiferromagnets. The Néel temperature is defined as the
transition temperature at which an antiferromagnet becomes paramagnetic. This theory
was later mathematically modelled by Van Vleck23.
A Néel antiferromagnet lattice comprises two intertwined sublattices with equal and
opposite magnetisation, hence resulting in zero net magnetisation. For example, an IrMn3
lattice has two magnetic sublattices A and B, shown in Figure 4, with the magnetisation
MA = -MB. According to the Weiss theory, the inter-sublattice molecular field coupling can
be given as nAB and nBA14. Similarly, the intra-sublattice coupling is given by nAA and nBB.
Therefore, the net molecular field of each sublattice is
𝐻𝐴 = 𝑛𝐴𝐴𝑀𝐴 + 𝑛𝐴𝐵𝑀𝐵 + 𝐻 2-7
𝐻𝐵 = 𝑛𝐵𝐵𝑀𝐵 + 𝑛𝐵𝐴𝑀𝐴 + 𝐻 2-8
here, nAA = nBB, nAB = nBA, and H is the external magnetic field.
In the absence of external field H = 0, the net magnetisation M = MA + MB = 0 and the
magnetisation of each sublattice becomes zero at TN, i.e. MA(TN) = MB(TN) = 0. The
spontaneous magnetisation of each sublattice in the paramagnetic region
(MA = MB = χHB = χHA) for an antiferromagnet can thus be given by:
𝑀𝐴 = (𝐶∗
𝑇) (𝑛𝐴𝐴𝑀𝐴 + 𝑛𝐴𝐵𝑀𝐵 + 𝐻) 2-9
Theory Chapter 2
11
𝑀𝐵 = (𝐶∗
𝑇) (𝑛𝐵𝐵𝑀𝐵 + 𝑛𝐵𝐴𝑀𝐴 + 𝐻) 2-10
where, C∗ = (μ0(n/2)meff2 )/3kB and n is the number of magnetic ions per unit volume with
n/2 on each sublattice14.
Equations 2-9 and 2-10 must have a nontrivial solution in order to fulfil the condition of
spontaneous magnetisation of each sublattice in the absence of a magnetic field. For this,
the determinant of the coefficient of MA and MB must be zero resulting in
[(𝐶∗
𝑇)𝑛𝐴𝐴 − 1]
2
− [(𝐶∗
𝑇)𝑛𝐴𝐵]
2
= 0 2-11
and rearranging the terms on the right-hand side and left-hand sides
(𝐶∗
𝑇)𝑛𝐴𝐴 − 1 = [(
𝐶∗
𝑇)𝑛𝐴𝐵] 2-12
This yields TN = C*·(nAA - nAB). The susceptibility of the antiferromagnet above TN can be
described by adding the sublattice coupling in the Curie-Wiess law.
𝜒 =
𝐶
𝑇 − 𝜃𝑃 2-13
where, θp = C*·(nAA + nAB) and C = 2C*.
The Néel temperature of disordered bulk IrMn3 is around (960 ± 10)K, and increases with
increasing iridium content in the alloy24. Additionally, TN varies from bulk antiferromagnetic
material to thin films. In this we focused on IrMn3 stoichiometry, hence hereafter this
stoichiometry will be referred to as IrMn.
Figure 4. The magnetic lattice structure of IrMn superimposed over the crystal lattice (a.). Magnetic sublattice showing inbound (A-yellow) and outbound magnetisation (B-green) with zero net magnetisation (b.) [images adapted from24].
Theory Chapter 2
12
2.1.3 Magnetic hysteresis
A ferromagnetic material, when subjected to an external magnetic field, results in an
alignment of its magnetisation in accordance with the applied magnetic field direction. The
magnetisation may increase from the remanent magnetisation (Mr) (which also can be
zero if the material was demagnetised before) up to the saturation magnetisation (Ms) at
the maximum when the saturation field (Hs) is applied. If the field is now reduced to zero,
the magnetisation decreases from Ms to Mr. To reduce the magnetisation to zero
(demagnetisation) an external magnetic field in the opposite direction is required, known
as the coercive field or coercivity (HC). The plot of the magnetisation as a function of
applied magnetic field strength for ferromagnetic material results in a closed-loop called
M(H) hysteresis loop, as shown in Figure 5. The suitability of a ferromagnetic material for
a particular application is widely determined by the characteristics shown by its hysteresis
loop (Ms, Mr, HC). For example, a squared shaped hysteresis loop, revealing two stable
states, might be suitable for data storage, whereas materials with small coercive fields
accompanied by a linear evolution of M(H) between remanence and saturation are the
preferred candidates for magnetic field sensor applications.
Figure 5. Exemplary M(H) hysteresis loop of the ferromagnetic material Co60Fe20B20 after 30 min annealing at 600°C: Mr is the remanent magnetisation at H=0; Ms denotes the saturation magnetisation and HC the coercivity.
2.1.4 Magnetic anisotropy
The magnetic anisotropy energy is a measure of the dependence of the ground-state
energy on the direction of the magnetisation. This anisotropy defines the preferential
(easy), the intermediate, and the magnetically hard directions of magnetisation in
Theory Chapter 2
13
materials. The overall magnetic anisotropy is a combination of several contributions.
There are three main contributing anisotropies in the materials studied throughout this
thesis, namely, magnetocrystalline anisotropy, shape anisotropy, and exchange
anisotropy14,15,25.
Magnetocrystalline anisotropy: The magnetocrystalline anisotropy energy is defined as
the energy required to rotate the magnetisation from the easy to a hard direction14. This
occurs due to the spin-orbit coupling, i.e. the coupling between the spin and orbital
momentum of the electrons. When an external field tries to reorient the electron’s
magnetic moment, the orbit of that electron also has to get reoriented, but due to strong
coupling to the crystal lattice, the magnetisation rotation might be hindered.
Shape anisotropy: Shape anisotropy is the result of the demagnetising field (Hd)
generated by the samples due to its own magnetisation14. If the sample is spherical and
a magnetic field is applied, the field will magnetise the sample to the same extent in all
directions. As the magnetisation is uniform, the demagnetisation field is distributed equally
in all directions. A non-spherical sample, for example, a long bar is easier to be
magnetised along the long axis than the short axis, due to the non-uniform demagnetising
field generated by the rod. The shape anisotropy must not be confused with
magnetocrystalline anisotropy as shape anisotropy is not an intrinsic property of the
material and solely depends on the shape of the FM sample.
Exchange Anisotropy: In the year 1956, Meiklejohn and Bean discovered a new form of
anisotropy in magnetic material systems and named it to exchange anisotropy26. They
stated that such anisotropy is the consequence of the interaction between the
ferromagnetic and the antiferromagnetic layers in a FM/AFM bilayer system when cooled
through the Néel temperature in the presence of the static magnetic field. This results in
a shift (or biasing) of the M(H) loop, generally in the opposite direction of the cooling field.
Such a shift is termed as the exchange bias field (HEB). See Figure 6 for a schematic
representation of the effect on the hysteresis loop.
Theory Chapter 2
14
Figure 6. M(H) hysteresis loop of CoFeB / IrMn bilayer at room temperature (RT). The green line represents the M(H) response when the sample is cooled from 200°C to RT in the presence of a magnetic field (120 mT), and the red line is the response of the as-deposited, demagnetised state.
Along with the shift in the hysteresis loop, exchange biased systems exhibit an angular
dependence kud·cosθ resulting in a unidirectional anisotropy in addition to the common
uniaxial anisotropy kua·sin2θ observed in a FM (see Figure 7). Here, kud and kua are the
unidirectional anisotropy and FM uniaxial anisotropy constants. The angle between the
magnetisation and the anisotropy axis is denoted with θ. The energy of such a coupled
bilayer system is the summation of both anisotropy energy per unit area.
𝐸 = −𝜇𝑜𝑀𝐹𝑀𝐻𝑡𝐹𝑀 𝑐𝑜𝑠 𝜃 − 𝜎𝑒𝑥 𝑐𝑜𝑠 𝜃 + 𝑘𝑢𝑎 𝑡𝐹𝑀𝑠𝑖𝑛2 𝜃 2-14
where, MFM is the magnetisation of the ferromagnetic layer with a thickness tFM and
σex = kud·tFM.
In this equation, the first energy term is due to the presence of the external magnetic field,
the second term represents the exchange bias energy, and the last term is related to the
uniaxial anisotropy energy in ferromagnets. Here it must be noticed that the overall
exchange bias field strength varies inversely with the thickness of the ferromagnetic layer
(its magnetisation, to be accurate) for thin films (typically up to 10 nm) and is independent
of the film thickness for thick films14.
In general, for an exchange bias system as the temperature increases and approaches
TN, the exchange bias field strength decreases and vanishes ultimately. However, it is
often observed in thin FM/AFM bilayers that the exchange bias effect disappears already
well below TN. The temperature at which the exchange bias field becomes zero (HEB = 0)
Theory Chapter 2
15
is, therefore, called blocking temperature (TB) and can be significantly lower than TN. For
example, IrMn has TN = 690 K and TB = 540 K27.
Figure 7. Uniaxial magnetic anisotropy induced through the exchange bias effect in CoFeB / IrMn after cooling from 200°C in the presence of 120 mT magnetic field. The coercive fields (a.) and the exchange bias fields (b.) are plotted as determined by angle-resolved longitudinal MOKE magnetometry at room temperature.
The exchange interaction, as well as further related effects, have been vastly investigated
for possible applications in magnetoresistive devices. For instance, the appropriate use
of an AFM material (layer) in a magnetoresistive layer stack assists in tuning the sensitivity
range of a magnetic field sensor and in suppressing the Barkhausen noise. The extensive
experimental research on AFM/FM layers stacks has also stimulated theoretical
descriptions, particularly concerning the microscopic origin of the effect28. In general, the
spins of FM are parallel to each other. However, at the interface to the AFM layer,
localised net moments arise from interfacial roughness, frustrated exchange bonds
(interfacial spins coupled antiferromagnetically), stress and dislocations, giving rise to the
exchange bias27,29,30.
2.2 Spintronics
Spintronics, as the name suggests, is the combination of the term spin and electronics,
whereby both, charge and spin of the electron(/s) are exploited. An often-proposed
scheme of a spintronic device is based on the magnetoresistive effects (MR). In particular,
magnetoresistive devices, based on the giant magnetoresistance (GMR) and tunnelling
magnetoresistance (TMR) effect, are comprised of a non-magnetic layer (conducting in
case of GMR and insulating for TMR) sandwiched between two FM layers.
Theory Chapter 2
16
In the year 1971, P.M. Tedrow and R. Meservey et al. laid the foundation of the TMR
effect by discovering the spin-polarised tunnelling for a Al / Al2O3 / Ni junction31. M.
Julliere, in 1975, proposed the spin conservation theory and presented a quantitative
explanation for the response of the magnetic tunnel junction (MTJ)12. M. Bowen et al.
emphasised the importance of the electronic and crystal structure of the entire electrode/
barrier/ electrode system in the year 2001 and measured a TMR resistance change of
60% at 30K for a Fe(001) / MgO(001) / FeCo(001) MTJ32. This MR response was four
times higher than any other contemporary Al2O3-based MTJ studied until that date.
However, the very low operational temperature limited the concept to find any device
application. Later on, Ikeda et al. presented in 2008 a breakthrough result for the
CoFeB / MgO / CoFeB system with an extraordinary TMR ratio of 603% at room
temperature33.
2.2.1 Tunnelling magnetoresistance effect
Tunnelling magnetoresistance is a magnetoresistive effect that occurs in a magnetic
tunnel junction, which is a multilayer structure consisting of two ferromagnets separated
by an ultra-thin non-magnetic (NM) insulator acting as a tunnel barrier. The charge
transfer through the tunnel barrier occurs due to quantum tunnelling and due to the
regulation of spin-polarised current controlled by the relative orientation of the
Figure 8. The development road map of the magnetoresistive devices [image was taken from34].
Theory Chapter 2
17
magnetisation of two ferromagnetic layers. The current through the two ferromagnetic
electrodes consists of two separate spin current channels for spin-up and spin-down
electrons (Mott`s two-current model35). The two channels encounter a different electrical
resistance in ferromagnetic materials due to the spin polarisation of the FM. The electrode
spin polarisation arises from the imbalance of the density of states of the spin-up and spin-
down electrons near the Fermi level in the ferromagnetic layers and thus from their
magnetic anisotropy in the magnetised state, see Figure 9.
Figure 9. Schematic representation of the tunnel magnetoresistance in the case of two identical ferromagnetic layers separated by a non-magnetic insulating barrier such as MgO. The tunnelling process conserves the spin. When the electronic states on each side of the barrier are spin-polarised, the electrons will more easily find free states to tunnel through the barrier if the magnetisations are parallel (a.) than if they are antiparallel (b.) to each other due to the ratio of the density of states of both electrons (spin-up, spin-down) at the Fermi level. The arrows in green and red show the higher and lower tunnelling probability of spin-polarised electron through a tunnel barrier, respectively. The yellow balls represent the electrons with their intrinsic spin orientation in grey [image redrawn from36].
In an FM/NM/FM structure, the overall electrical resistance depends then on the mutual
magnetisation directions of both FM:
𝑇𝑀𝑅 =2𝑃1𝑃2
1 − 𝑃1𝑃2=
𝑅↑↓ − 𝑅↑↑
𝑅↑↑ 2-15
where, P1 and P2 represent the spin polarisations of two ferromagnetic layers and R↑↓ and
R↑↑ are the resistance of two FM in parallel and antiparallel magnetisation.
Theory Chapter 2
18
2.3 Optical spectroscopy
As discussed previously, the electronic properties of the MTJs ferromagnetic electrodes
play a decisive role in MTJs. Hence it is very important to probe and understand these
properties. There are many experimental techniques to measure the electronic states of
materials, namely: photoelectron spectroscopy (PES), inverse photon-electron
spectroscopy (IPES), electron energy loss spectroscopy (EELS), optical spectroscopy,
etc. Among these methods, the optical spectroscopy techniques hold the high figure of
merit, as they are relatively simple and non-invasive techniques. Optical spectroscopy
techniques are based on the principle of light-matter interaction and can be broadly
classified on the basis of the wavelength of electromagnetic wave used to probe the
material (THz, visible, UV, X-ray etc.). In the present work, two optical spectroscopy
techniques namely, spectroscopic ellipsometry (SE) and magneto-optical Kerr effect
(MOKE) spectroscopy in the NIR (SE only), visible, and UV spectral ranges were used to
assess the electronic as well as magnetic properties of ferromagnetic materials.
Additionally, in this framework, the optical spectroscopies were used to investigate
structural changes in the ferromagnetic layer and their influence on its optical properties.
Both spectroscopic techniques follow the principles of polarimetry, where the sample
under investigation is illuminated with light of known polarisation. The changes in the
polarisation state of the reflected light are recorded, thus providing information about the
electronic states of the sample. The information that can be deduced from these two
methods includes the complex refractive index, the layer thickness, roughness, magnetic
hysteresis, exchange coupling, and crystallinity etc.
2.3.1 Polarisation of light
Light is an electromagnetic wave with the electric and magnetic fields vectors oscillating
perpendicular to each other and also perpendicular to the direction of propagation. It is
sufficient to discuss only the electric field vector to describe the polarisation of the light.
Thus, an electromagnetic wave travelling in the z-direction can be mathematically
represented as a superposition of two orthogonal components of the electric field:
(𝑧, 𝑡) = (𝐸𝑥 + 𝐸𝑦)𝑒𝑖(𝑘𝑧−𝜔𝑡) 2-16
Theory Chapter 2
19
here, Ex and Ey are the components of the electric field projected onto the x- and y- axes.
The quantities X and Y are unit vectors with magnitude one and direction pointing along
their respective axes.
For example, Ey = iEx, where Ex is a real number, means that the y-component of the
electric field is phase-shifted with respect to the x- component by π/2. Substituting the
value in the above equation and considering the real part we obtain:
𝑅𝑒[𝐸(𝑧, 𝑡)] = 𝑅𝑒[𝐸𝑥𝑒𝑖(𝑘𝑧−𝜔𝑡)] + 𝑅𝑒 [𝑒
𝑖𝜋2 𝐸𝑥𝑒
𝑖(𝑘𝑧−𝜔𝑡)]
2-17 = 𝐸𝑥 𝑐𝑜𝑠(𝑘𝑧 − 𝜔𝑡) +𝐸𝑥 𝑐𝑜𝑠 (𝑘𝑧 − 𝜔𝑡 +𝜋
2)
= 𝐸𝑥[𝑐𝑜𝑠(𝑘𝑧 − 𝜔𝑡) − 𝑠𝑖𝑛(𝑘𝑧 − 𝜔𝑡)]
In the above equation, it is clear that the y component lags the x component by a quarter
cycle. Thus, the net electric field vector maintains a constant magnitude and appears to
rotate in a circular pattern in the x-y plane. Various polarisation states with the
corresponding phase differences are shown in Figure 10.
Figure 10. Various polarisation states of light occurring as a result of the various phase difference (𝛿𝑥 −𝛿𝑦 and 𝛿𝑦 − 𝛿𝑥) between the two components of the electric field (along x- and y- axes) of equal amplitude
(𝐸𝑦 = 𝐸𝑥) [image taken from37].
Since Cartesian coordinates x, y and z are interchangeable and are relative directions, it
is standard to define the electric field components respective to the sample plane. The
plane perpendicular to the surface of the sample that contains the vector pointing the
direction of propagation of light is called the plane of incidence. Perpendicular to the
Theory Chapter 2
20
propagation vector is two mutually perpendicular vector components of the electric field
vector of (E), defining the polarisation of light (the x-y plane). These components of E are
therefore parallel and perpendicular to the plane of incidence and are named as “parallel”
(p) and perpendicular “senkrecht” (s), respectively.
2.3.2 Jones formalism representation of polarisation
R. Clark Jones introduced in 1941 a matrix algebra to describe the polarisation of light
and also the influence of optical elements on the light polarisation38. As discussed in the
previous section, the polarisation state can be represented by superimposing two
electromagnetic waves oscillating along x- and y- axes, respectively. The Jones vectors
define the amplitude (E) and phase (ϕ) of the electric field along the x and y axes (p and s).
𝐽 = (𝐸𝑥𝑒
𝑖𝜙𝑥
𝐸𝑦𝑒𝑖𝜙𝑦
) 2-18
Since Jones matrix algebra is applicable only for polarised light, the following discussion
is restricted for linearly polarised s and p waves. A linearly polarised light wave incident
onto a sample is denoted as s (or p) polarised when the electric field vector is
perpendicular (or parallel) to the plane of incidence (x-z). For example, the change in the
polarisation of the reflected light (Jr) from a reflecting sample can be written as the matrix
product of the reflection matrix R to the Jones vector of the incident light Ji.
where, R is the reflection matrix, and r is the reflection coefficient of s and p polarised
light.
𝐽𝑟 = 𝑅 × 𝐽𝑖 2-19
𝑅 = [𝑟𝑠𝑠 𝑟𝑠𝑝𝑟𝑝𝑠 𝑟𝑝𝑝
] 2-20
(𝐸𝑟𝑥𝑒
𝑖𝜙𝑟𝑥
𝐸𝑟𝑦𝑒𝑖𝜙𝑟𝑦
) = [𝑟𝑠𝑠 𝑟𝑠𝑝𝑟𝑝𝑠 𝑟𝑝𝑝
] × (𝐸𝑖𝑥𝑒
𝑖𝜙𝑖𝑥
𝐸𝑖𝑦𝑒𝑖𝜙𝑖𝑦
) 2-21
Theory Chapter 2
21
2.3.3 Light-matter interaction
Often, two physical quantities are used to describe the interaction of light and matter,
namely, the complex refractive index (N) and the dielectric function (ε). Both of them are
complex numbers.
𝑁 = 𝑛 + 𝑖𝑘 2-22
and
휀 = 휀1 + 𝑖휀2 2-23
Here n (real part) is the refractive index of the material and describes the phase velocity
of the light travelling through the material, and k (imaginary part) is called extinction
coefficient, describing the loss of wave energy to the material. Similarly, in the case of the
dielectric function, the real part denotes the dielectric capacity of material and the
imaginary part determines the dielectric losses. These two physical quantities are related
by
휀 = 𝑁2 2-24
In general, the optical response of the material is a direction-dependent property. In an
optically anisotropic material, ε can also be described by a second-rank tensor called
dielectric tensor or permittivity tensor ε.
휀 = (
휀𝑥𝑥 휀𝑥𝑦 휀𝑥𝑧
휀𝑦𝑥 휀𝑦𝑦 휀𝑦𝑧
휀𝑧𝑥 휀𝑧𝑦 휀𝑧𝑧
) 2-25
For a (non-magnetised) medium with a biaxial symmetry, the dielectric tensor can be
reduced to the principal dielectric constants (or principal refractive indices) given as
휀 = (
휀𝑥𝑥 0 00 휀𝑦𝑦 0
0 0 휀𝑧𝑧
) = (
𝑁𝑥𝑥2 0 0
0 𝑁𝑦𝑦2 0
0 0 𝑁𝑧𝑧2
) 2-26
Theory Chapter 2
22
For an optically isotropic (non-magnetised) material εxx = εyy = εzz. The dielectric tensor for
isotropic materials is given by
휀 = (
휀𝑥𝑥 0 00 휀𝑥𝑥 00 0 휀𝑥𝑥
) 2-27
The diagonal components of this tensor can be determined experimentally, e.g. by SE.
Many optical spectroscopy techniques, including SE and MOKE work in reflection
geometry. The diagonal components of the reflection matrix (equation 2-20) are directly
related to the principal component of the dielectric tensor (and to the optical constants)
and are determined by spectroscopic ellipsometry, while the off-diagonal components of
the reflection matrix are related to the off-diagonal components of dielectric tensor and
determined using magneto-optical Kerr effect spectroscopy or by generalized
spectroscopic ellipsometry or Müller matrix ellipsometry. Each of these polarimetry
techniques is discussed in detail in the following sections.
Selection rules
Another way to explain the interaction of light with matter is by taking into account the
absorption of a single photon by a single electron through electronic transitions. Like any
other physical process, the laws of conservation have to be followed in the process of
photons absorption. These conservation laws are often narrated as selection rules39.
Table 1. The selection rules of (magneto-) optical transitions
Energy 𝐸𝑓 − 𝐸𝑖 = ħ𝜔 The energy of the photon absorbed must be equal to the energy difference of the initial and final states
Momentum ħ𝜔
𝑐≈ 0
As a photon has negligible linear momentum compared to the electron, the linear moment of the electron must be conserved → “vertical transition.”
Spin ∆𝑠 = 0 Since photons carry no spins, during an electronic transition, the spin of the electron must be conserved.
Orbital angular momentum quantum number
∆𝑙 = ±1 The total orbital momentum is conserved. As a photon has 𝑙 = 1ħ, it means that the only allowed transition are s to p, p to d, d to f, etc.
Magnetic quantum number
∆𝑚𝑙 = ±1, 0
The total orbital momentum along z-direction (𝑚) must be conserved. The type of the absorbed photon is thus circularly left (-1), or right (+1) polarised, or linearly (0) polarised.
Theory Chapter 2
23
2.3.4 Spectroscopic ellipsometry
When linearly polarised light is incident on a sample surface at an oblique incidence angle,
the reflected light is in general elliptically polarised. In spectroscopic ellipsometry, the
change in the polarisation of the reflected with respect to that of the incident light is
measured. The name “ellipsometry” is related to the analysis of the elliptical polarisation
of the reflected light. Figure 11 shows the principle of ellipsometry.
Figure 11. The schematic of spectroscopic ellipsometry [image redrawn from40]
In a typical rotating compensator ellipsometer, the incident light is linearly polarised. Such
polarisation can be explained as the combination of p and s components of the electric
field oscillating in-phase. The amplitude of each component determines the orientation in
the quadrant. The reflected light from the surface of the sample is elliptically polarised,
which means that both the amplitude and the phase of the p and s components are altered
so that the projection of the net electric field vector forms an ellipse onto a plane
perpendicular to the light propagation direction.
The light interacts with the material following Maxwell’s equations, which in turn, give the
boundary conditions at each interface between two media. The boundary conditions
provide a predictable solution of the interaction of p and s components of the electric field
to the samples and can be used to obtain the Fresnel equations.
𝑟𝑝 =𝐸𝑟𝑝
𝐸𝑖𝑝=
𝑛𝑡 𝑐𝑜𝑠 𝜃𝑖 − 𝑛𝑖 𝑐𝑜𝑠 𝜃𝑡
𝑛𝑡 𝑐𝑜𝑠 𝜃𝑖 + 𝑛𝑖 𝑐𝑜𝑠 𝜃𝑡 2-28
𝑡𝑝 =𝐸𝑡𝑝
𝐸𝑖𝑝=
2𝑛𝑖 𝑐𝑜𝑠 𝜃𝑖
𝑛𝑡 𝑐𝑜𝑠 𝜃𝑖 + 𝑛𝑖 𝑐𝑜𝑠 𝜃𝑡 2-29
Theory Chapter 2
24
𝑟𝑠 =𝐸𝑟𝑠
𝐸𝑖𝑠=
𝑛𝑖 𝑐𝑜𝑠 𝜃𝑖 − 𝑛𝑡 𝑐𝑜𝑠 𝜃𝑡
𝑛𝑖 𝑐𝑜𝑠 𝜃𝑖 + 𝑛𝑡 𝑐𝑜𝑠 𝜃𝑡 2-30
𝑡𝑠 =𝐸𝑡𝑠
𝐸𝑖𝑠=
2𝑛𝑖 𝑐𝑜𝑠 𝜃𝑖
𝑛𝑖 𝑐𝑜𝑠 𝜃𝑖 + 𝑛𝑡 𝑐𝑜𝑠 𝜃𝑡 2-31
where, rp, rs, tp, and ts are the reflection and transmission coefficients for p and s
component θi and θt are the angle of incidence and transmittance, respectively; ni and nt
are the refractive index of incidence and transmission medium, respectively.
Often the SE samples have multiple interfaces, resulting in multiple reflections. The
superposition of these multiple reflections at the interface generates an interference
depending on the relative phase of each reflected light wave (see Figure 12).
Figure 12. The schematic of an optical model for an ambient / thin film / substrate structure, showing the reflected and refracted light at each interface. Using the Fresnel coefficients, the contribution of reflections from each interface can be calculated [image was taken from40].
The detector then measures this change in polarisation in terms of the ellipsometric
parameters Ψ and Δ. These parameters can be related to the Fresnel reflection
coefficients of p and s polarised light, resulting in the fundamental equation for
ellipsometry:
𝜌 =𝑟𝑝𝑝
𝑟𝑠𝑠= 𝑒−𝑖∆𝑡𝑎𝑛𝜓 2-32
Typically, SE is used to determine the optical constants (N) of a material and/or thickness
of thin films. Every material has unique optical constants which are directly related to the
electronic structure of the material. Hence optical constants are considered as a
fingerprint of materials. In ellipsometry, a suitable model describing the sample has to be
Theory Chapter 2
25
developed to fit the experimental data Ψ and Δ. Using Fresnel’s equations, predictive Ψ
and Δ spectra from the model are calculated. If the optical constants or the thickness of
the film are unknown, the estimated values or range are considered as input parameters.
A non-linear regressive approach (for example the Levenberg-Marquardt method) is then
used to find the best match of the calculated Ψ and Δ spectra to the measured ones. The
quantitative comparison of the best match between the calculated spectra and the
measured ones is expressed in terms of Mean Squared Error (MSE). The lower the MSE
is, the better the fit.
𝑀𝑆𝐸𝑁𝐶𝑆 = √1
3𝑛 − 𝑚∑[(
𝑁𝐸𝑖 − 𝑁𝐺𝑖
0.001)2
+ (𝐶𝐸𝑖 − 𝐶𝐺𝑖
0.001)2
+ (𝑆𝐸𝑖 − 𝑆𝐺𝑖
0.001)2
]
𝑛
𝑖=1
2-33
where, n is the number of wavelengths in the measured spectral range, m is the number
of fit parameters in the optical model, and N = cos(2Ψ), C = sin(2Ψ)·cos(Δ),
S = sin(2Ψ)·sin(Δ) are the Mueller matrix components for an isotropic material. The
discussion of the Mueller matrix is beyond the scope of this work and can be found in
reference37. Typical precision in measuring the N, C, and S parameters is ~0.001, hence
this factor is included in the MSE definition, implying that ideal data modelling will have
an MSE of ~1.
2.3.5 Magneto-optical Kerr effect spectroscopy
Michael Faraday, in 1846, discovered a magneto-optical effect, which was later called the
Faraday effect. He observed that a linear polarised light beam changes its polarisation
state after passing through a piece of glass in the presence of a magnetic field. Later,
John Kerr reported a similar effect for the reflected light from a piece of polished iron and
found that this effect is proportional to the magnetisation in the sample41,42. Today, the
characterisation techniques based on the Faraday and magneto-optical Kerr effect
(MOKE) are prominent measurement tools in the field of magnetism.
In MOKE (or Faraday effect) based measurement techniques, the polarisation change of
the reflected (transmitted) light is measured in terms of the complex Kerr (Faraday)
rotation (θ). Practically, θ is measured as the ratio of off-diagonal components to either
one of the diagonal components of the reflection coefficient matrix. (equation 2-34).
Theory Chapter 2
26
𝜃 =𝑟𝑝𝑠
𝑟𝑠𝑠 or 𝜃 =
𝑟𝑠𝑝
𝑟𝑝𝑝 2-34
For the ease of presentation, later only the rp component is considered for the calculations.
However, the mathematical equations remain identical for the rs component, as well.
Additionally, all the MOKE experiments presented in this thesis were conducted with
incident light of p polarisation. In general, the reflection coefficient and the dielectric
function (optical constants) of a material are complex numbers. Therefore, the Kerr
rotation (ΘK) is also expressed in terms of complex numbers. The real part of MOKE (or
Faraday) rotation represents the tilt of the polarisation ellipse (θK or θF) and the imaginary
part denotes its ellipticity (ηK or ηF). η is caused by the difference in the absorption of right
circularly and left circularly polarised light, whereas θ is caused by the difference in the
phase velocity between the left circular and right circular polarised light in a magnetised
medium.
𝛩𝐹 = 𝜃𝐹 + 𝑖𝜂𝐹 2-35.
𝛩𝐾 = 𝜃𝐾 + 𝑖𝜂𝐾 2-36.
Based on the relative orientation of magnetisation and incident plane of light, MOKE can
be classified into three categories: polar (P-MOKE), longitudinal (L-MOKE), and
transversal (T-MOKE). These geometries correspond to the magnetisation aligned normal
to the sample surface (P-MOKE), or in the sample plane but either in the plane of
incidence (L-MOKE) or perpendicular to it (T-MOKE). Figure 13 summarises the basic
geometries of MOKE.
Figure 13. The three geometrical configurations for MOKE, namely polar (a.), longitudinal (b.) and transversal (c.).
In the case of P-MOKE and L-MOKE, the magnetisation lies in the plane of incidence,
which influences the off-diagonal reflection coefficients rps and rsp. In T-MOKE the
Theory Chapter 2
27
magnetisation is perpendicular to the plane of incidence and influences only the diagonal
component reflection coefficient rpp (and/or rss). As a result, only the reflected light intensity
changes as a function of the magnetisation, but no Kerr rotation or ellipticity is measurable
in this configuration.
Phenomenological explanation of MOKE
According to the selection rules described above, a circularly polarised photon can induce
an electronic transition between two states with the same spin polarisation. If the density
of spin-up and spin-down states are not equal, then a different number of left circular
polarised and right circular polarised photons will be absorbed, which in turn results in the
magneto-optical (MO) signal of ferromagnetic materials.
On an atomic scale, the magnetisation in ferromagnetic materials can be seen as a
perturbation that lifts the degeneracy of the electronic states by spin-orbit coupling and
exchange interaction43. Spin-orbit coupling is the interaction between the electron spin
magnetic moment and the orbital angular momentum. This is similar to the Zeeman effect,
but it is caused by the internal magnetic field generated by the electrons orbital motion.
Figure 14 exemplifies these two effects on the energy diagram of a 3d transition metal
ferromagnet. The left side of Figure 14 shows the lifting in the degeneracy of the d and p
electronic states due to the spin-orbit coupling. The energy distance, ESO, between the
split states corresponds to the strength of the spin-orbit interaction. The exchange
interaction takes place among the spins of electrons in two adjacent atoms, resulting in
parallel spin alignment (Pauli principle) in the case of a ferromagnet. The exchange
interaction splits the spin-up and the spin-down electronic states by the exchange energy
Δex, as shown in Figure 14. All possible magneto-optical electronic transitions between p
and d states of a 3d ferromagnet by blue and red vertical lines (depending on the circular
polarisation state of the impinging photon) are shown in Figure 14.
Theory Chapter 2
28
Figure 14. The schematic energy state diagram of a 3d-ferromagnet, showing optical transitions induced by left (blue) and right (red) polarised photons for a system where only spin-orbit coupling is present (left diagram) and for a system where spin-orbit coupling and exchange interaction is present (middle diagram). The notation in the | ⟩ brackets contains the orbital number (𝑙), magnetic number (𝑚),
and spin orientation (↑ or ↓). The right side diagram shows the corresponding absorption spectra of left and right circular polarised light. [adapted from44]
Using ab-initio calculations based on the Kubo linear response formalism45, it was shown
that the presence of spin-orbit coupling and/or exchange interaction in a material yield to
the appearance of off-diagonal elements in the conductivity tensor. For a magnetised
sample, the dielectric tensor can be written as follows:
휀 = 휀𝑥𝑥 (1 𝑖𝑄𝑍 −𝑖𝑄𝑌
−𝑖𝑄𝑍 1 𝑖𝑄𝑋
𝑖𝑄𝑌 −𝑖𝑄𝑋 1) 2-37
here, Q is the so-called magneto-optical Voigt constant, which, for a given value of the
externally applied magnetic field is material-specific. The Voigt constant can have different
values along the x-, y-, and z- space directions.
In this thesis, only the P-MOKE geometry is employed, and in order to keep the
corresponding calculations simple, the magnetisation is placed parallel to the z-axis. By
assuming that Q is parallel to the direction of magnetisation
Theory Chapter 2
29
(
𝑄𝑋
𝑄𝑌
𝑄𝑍
) = 𝑄 ∥ 𝑀 2-38
the Q vector will only have a component along the z-axis. For an isotropic medium with
the magnetisation pointing along the z-direction, equation 2-37 can be reduced to the
following dielectric tensor
휀 = 휀𝑥𝑥 (1 𝑖𝑄𝑍 0
−𝑖𝑄𝑍 1 00 0 1
) 2-39
The refractive index for the left-circularly polarized light (Nl), and right-circularly polarized
light (Nr) can be given as
𝑁𝑟 = 𝑁(1 −𝑄
2⁄ ) and 𝑁𝑙 = 𝑁(1 +𝑄
2⁄ ) 2-40
As mentioned previously, the Faraday effect arises. Thus the complex Faraday rotation
after transmission through a medium of length L:
𝛩𝐹 = 𝜃𝐹 + 𝑖𝜂𝐹 =𝜔𝐿
2𝑐(𝑁𝑟 − 𝑁𝑙) = −
𝜔𝐿
2𝑐𝑄𝑁
2-41
here, c and ω is the velocity and angular frequency of light. Similarly, the Kerr effect which
is measured in the reflection mode the resultant reflection coefficients for the two circular
polarised state of light is
𝑟𝑟 = 1−𝑁(1−
𝑄2⁄ )
1+𝑁(1−𝑄
2⁄ ) and 𝑟𝑙 =
1−𝑁(1+𝑄
2⁄ )
1+𝑁(1+𝑄
2⁄ ) 2-42
Thus, the complex Kerr rotation angle due to these reflection coefficients, i.e. rr an rl is
Θ𝐾 = θ𝐾 + iη𝐾 = arctan (𝑖𝑟𝑟 − 𝑟𝑙𝑟𝑟 + 𝑟𝑙
) ≈ 𝑖𝑟𝑟 − 𝑟𝑙𝑟𝑟 + 𝑟𝑙
2-43
substituting the reflection coefficient from equation 2-42 in 2-43 one obtains the reflection
on a bulk (half space) sample:
Theory Chapter 2
30
Θ𝐾 = θ𝐾 + iη𝐾 = 𝑖𝑁𝑄
1 − 𝑁2 + 0.25𝑁2𝑄2≈ 𝑖
𝑁𝑄
1 − 𝑁2 2-44
In case of the reflection from a thin magnetic single layer on a non-magnetic substrate for
a thickness less than the wavelength d ≪ λ
Θ𝐾 = θ𝐾 + iη𝐾 == −ω
2C·
𝑁2𝑄𝑑
1 − 𝑁2
2-45
For a layer stack samples, a more complex optical model needs to be used and the
relation between θ and Q will change accordingly (see ref.46).
Often, the results are presented in terms of optical conductivity instead of the dielectric
function, as the optical conductivity is amenable to be calculated from the Kubo linear
response formalism (see, e.g. ref.44,45). The relation between ε and σ is given in
equation 2-46
휀 = 1 +4𝜋𝑖
𝜔𝜎
2-46
where, ε is the dielectric tensor and σ is the conductivity tensor. In the case of P-MOKE
of an optically isotropic sample (as the samples considered in this thesis) the conductivity
tensor (σ) is scaled assuming an isotropic sample.
𝜎 = (
𝜎𝑥𝑥 𝜎𝑥𝑦 0
−𝜎𝑥𝑦 𝜎𝑥𝑥 0
0 0 𝜎𝑥𝑥
) 2-47
Due to the symmetry conditions, the components σxz, σyz, σzx, and σzy can be equated to
zero (σxz = σyz = σzx = σzy = 0). The diagonal components (σxx = σyy = σzz) are symmetric
and represent the optical conductivity of the material, whereas the off-diagonal
components (σxy = -σyx) are accountable for MOKE. As mentioned in the previous section
“Light-matter interaction”, the dielectric tensor elements are complex quantities.
Therefore, the conductivity tensor elements are also complex quantities. Knowing the
components of the conductivity tensor, one can write the complex Kerr rotation as47:
𝜃𝑘 + 𝑖𝜂𝐾 =
𝜎𝑥𝑦
𝜎𝑥𝑥√1 +4𝜋𝑖𝜔 𝜎𝑥𝑥
2-48
Experimental Chapter 3
31
Chapter 3: Experimental
In this chapter, the sample preparation and the measurement techniques that were used
in the framework of this thesis are introduced. The basic principles and the technical
properties of the equipment for the fabrication, processing, and characterisation of the
samples are explained. The intention is only to provide a short overview and further details
regarding any particular method can be found in the listed references.
3.1 Sample preparation
The samples investigated in this thesis were prepared by magnetron sputtering. The
CoFeB thin film samples characterised in the chapter “The electrodes: 3d-transition metal
boride” (chapter 4) with the corresponding capping layers were prepared using the
sputtering facility available under the professorship of “Functional Magnetic Materials”,
headed by Prof. Olav Hellwig. The MTJ layer stacks and the IrMn / CoFeB bilayer
discussed in the chapter “Exchange bias” (chapter 5) were deposited at “Singulus
Technologies AG”.
The fundamental principle of sputtering is based on momentum interchange between the
accelerated ions of inert gas and the target material. Typically, argon (Ar) gas serves for
this purpose, which is ionised to Ar+ ion plasma by the application of a high voltage
between the metallic target as cathode and anode in close vicinity. The energised Ar+ ions
strike the target material atoms and provide them with the necessary kinetic energy to
reach the substrate. In the magnetron sputtering a toroidal magnetic field generated by
permanent magnets underneath the cathode surrounds the target (see Figure 15). Due
to the presence of the magnetic field, a dense plasma is confined over the target, resulting
in a higher yield of the ejected target material in comparison to other physical sputtering
methods.
Due to industrial privacy norms, the sputtering equipment used at “Singulus Technologies
AG” is not discussed in this documentation. However, the equipment available at our
institute is discussed in detail and fundamentally similar to the one at Singulus
Technologies. The sputtering chamber is custom made with four confocal configured
sputtering targets. The magnetrons of each target are powered by a DC(/RF) source that
enables to deposit the layers from a single target or from more targets. For the conformal
deposition of films, the substrate is rotated above the target at a distance of around
Experimental Chapter 3
32
250 mm. This results in a uniform deposition across the substrate wafer with a maximum
diameter of 100 mm. The deposition was conducted under HV conditions, with a base
pressure of ~2×10-4 Pa and argon as the working gas. The argon pressure can be
regulated by the pump rate and the argon flow in the chamber. The pump rate is controlled
with a throttle adjustment at a turbopump. All the depositions were performed at a constant
argon pressure of 0.35 Pa. The nominal deposition rate and the deposited layer thickness
were monitored using a quartz microbalance placed close to the substrate holder. The
resonance frequency of the quartz oscillator shifts due to its increasing mass during the
deposition. The calibration of the quartz crystal microbalance was performed with XRR
(X-ray refractometry) on the deposited samples. The thickness-modulated periodic
oscillations measured in X-ray reflectometry were evaluated with the GenX fitting tool for
extracting the thickness of the sputtered CoFeB film48. The uncertainty of the simulated
thicknesses of the CoFeB films and of the capping layers was in the order of 10% of the
nominal values. This deviation from the nominal thickness is mainly due to the short
sputter time and inherent inaccuracy in measurements and simulation.
Figure 15. Schematic representation of the magnetron sputtering. The argon ions (in red) are responsible for the target etching. The ejected particles (in grey) are sputtered towards the substrate. The direction of the magnetic field used for confining the electrons and the ions close to the target is illustrated by the blue arrows.
3.2 Annealing
In this work, three different types of annealing systems were used for annealing the
samples, namely: in-situ SQUID-VSM system, in-situ macro-MOKE system, and laser
Experimental Chapter 3
33
irradiation system. The two foremost systems can provide ultra-high vacuum (UHV)
conditions. Each of these methods is discussed in the following section.
3.2.1 In-situ SQUID-VSM annealing
The Quantum Design's Superconducting Quantum Interference Device–Vibrating Sample
Magnetometry (SQUID - VSM) Magnetic Property Measurement System-3 (MPMS-3), is
equipped with a sample holder with a resistive element for heating. This option allows
heating from ambient to 750°C under UHV conditions. The central part of the oven is an
insulated heater stick assembly with a sample space of 5 mm in width which is mounted
at the end of a dedicated sample rod and introduced directly into the sample chamber.
The sample is secured on the oven stick holder with zircar cement (ACA5) in conjunction
with tightly wrapped copper foil around it, following the recipe provided in49.
Figure 16. Oven sample mounting platform with a mark at 66 mm to place the sample on oven heater stick [image was taken from49].
3.2.2 In-situ macro-MOKE annealing
A special heating assembly was designed in the frame of this thesis to anneal larger
samples (3 × 2) cm2 in the presence of magnetic field (up to 500 mT) in UHV environment,
unlike in-situ SQUID-VSM where the sample size is strictly restricted to (1 × 1) mm2. This
heater is assembled in the macro-MOKE spectroscopy system, which not only gives the
freedom of annealing in UHV but also allows for depositions and in-situ MOKE
characterisation. However, in-situ MOKE measurements are at this point beyond the
scope of this work. The heater is carved from a solid copper rod, to ensure homogeneous
heating over the whole sample. The thermal energy is provided with a halogen lamp of
250 W mounted in a cavity, as shown in Figure 17, that enables to reach a maximum
temperature of 600°C. For the precise measurement of the temperature, a type-K
Experimental Chapter 3
34
thermocouple (chromel–alumel) is placed right below the sample holder and read with an
Omron (E5EN-H) high-resolution temperature controller. To improve thermal conductivity
between the sample holder and sample a thin layer of silver epoxy was used.
Figure 17. CAD modelled isometric view (a.) and transverse view (b.) of the heater assembly developed for the macro-MOKE system.
3.2.3 Laser annealing
Figure 18 shows the sketch of the experimental setup used for the laser annealing at the
Laserinstitut of the Hochschule Mittweida. At a much higher magnification, the heat profile
induced by the laser beam into an exemplary layer stack is sketched. The laser annealing
experiments were performed by Ms. Sandra Busse. The planning of the experiments was
discussed in the DFG-3D-TMR project team. An Nd: YAG laser with an emission
wavelength of 1064 nm was used to locally heat the samples. The laser system is capable
of supplying both continuous wave (CW) and pulsed wave (PW) laser radiation. The laser
beam was first attenuated and then rapidly deflected by using a galvanometer scanner
and focused by an f-theta objective with a focal length of 80 mm. The resulting focal radius
(86%) of the laser beam was 10 µm. For annealing a larger area raster scanning was
performed with 2.5 µm distance between each scanned line, providing an adequate
overlap of consecutive line scans. The external magnetic field was provided by
electromagnets in the Helmholtz coil configuration, generating a maximum magnetic field
of 420 mT. Often the fluence delivered by laser is given in terms of laser intensity, i.e.,
power per unit area (W·cm-2)50. The following formula can be used to convert the laser
power to laser intensity:
𝐼𝑛𝑡𝑒𝑛𝑠𝑖𝑡𝑦 =2𝑃
𝜋 𝑟2
3-1
where, P is the laser power delivered to the sample, r is the radius of the laser beam, and
multiplier 2 is due to the integration assuming a Gaussian laser profile. All the results for
the laser annealing presented in this thesis are thus given in terms of laser intensities.
a. b.
.
Experimental Chapter 3
35
Figure 18. Schematics of the experimental setup for Nd: YAG laser [image prepared for own publication51].
3.3 Measurement techniques
3.3.1 Structural and topographic characterisation
X-ray diffraction
The atomic stacking has an important implication on the properties of ultrathin and thin
films. For the crystallographic characterisation of the thin films and layer stacks, X-ray
diffractometry (XRD) was used. The thin film thickness was measured by X-ray
reflectometry (XRR). The structural characterisation presented in this work was conducted
using the diffractometer 3000PTS from Seifert-FPM, the Rigaku SmartLab from Rigaku
Corporation, and the KMC-2 Diffraction station52 at Bessy II (Helmholtz-Zentrum Berlin).
Although three different tools were used for the X-ray diffraction measurements, the core
principle of all three types of equipment is the same.
In general, the laboratory X-ray diffraction instrument utilises an X-ray source consisting
of an evacuated tube, in which electrons are emitted from a heated tungsten filament.
These electrons are accelerated by an electric potential (here U = 40 kV at I = 30 mA) to
strike on a water-cooled copper target, producing X-ray radiation with wavelengths
characteristic for the electronic levels in the copper target (mainly copper Kα and Kβ) and
polychromatic background radiation (Bremsstrahlung). For this reason, the instrument is
further equipped with a monochromator crystal to provide beam monochromaticity (with
Experimental Chapter 3
36
solely Cu-Kα radiation). A scintillation detector mounted on the 2θ circle of the goniometer
is then used to measure the diffracted angular dependent intensity from the sample. The
recorded scan consists of constructive interference peaks from adjacent lattice planes
which fulfil the Bragg equation
𝑛𝜆 = 2𝑑 𝑠𝑖𝑛 𝜃 3-2
where, is the wavelength of the X-ray radiation, d is the interplanar spacing, n is an
integer number and θ is the diffraction angle.
The acquired peak positions from a θ-2θ scan can be compared with reference databases
for particular materials and the relevant Miller indices can be assigned to the peaks. In
this study, the X-ray powder diffraction files (PDF) were used as references, which were
published by the International Centre for Diffraction Data (ICDD). The Bragg peak in a θ-
2θ scan from a particular crystallite can be used to estimate the crystallite size (L) using
the Scherrer formula
𝐿 ≈𝐾𝜆𝐶𝑢
𝛥(2𝜃)𝑐𝑜𝑠𝜃 3-3
With a shape factor of crystallites K ≈ 0.9 considering cubic crystallites, the wavelength of
X-ray radiation Cu ≈ 0.154 nm and Δ(2θ) as the full width at half maximum (FWHM) of
the reflex at θ, given in radians53.
Figure 19. Drawing of the diffractometer showing the general scheme of various goniometers and measurement axes [image taken from54].
Experimental Chapter 3
37
An X-ray diffractometer can also be employed to perform XRR measurements. In XRR
measurements, a standard θ-2θ scan is performed but at small angles, below 2θ ≈ 8°.
Due to interference effects affecting the reflected X-ray radiation from the sample surface
and its interfaces, oscillations are observed in the reflectance curve plotted against the 2θ
angle along with exponential decay of the intensity. The measured oscillations contain
information about the thickness, density of the material, and interface roughness. A
careful modelling of the measured reflectivity curves can unwrap the above-mentioned
information about the investigated sample. For this purpose, an open-source software
GenX was used48. The general equation for XRR data fitting is shown below
𝑠𝑖𝑛2 𝜃𝑖 =𝜃𝑐
2 + (𝑛𝑖 + ∆𝑛)2𝜆2
4𝑡2
3-4
θi is the observed position of the maximum or minimum of the ith interference fringe, θc the
critical angle for total reflection, ni is an integer, Δn = 0.5 or 0 for the maximum and
minimum, respectively, is the X-ray wavelength, and t is the film thickness.
Atomic force microscopy
The atomic force microscope was developed by Binning and coworkers55, shortly after the
introduction of the scanning tunnelling microscopy. For atomic force microscopy, a sharp
tip prepared on a cantilever is brought in close proximity to the sample surface. The forces
acting between the tip and the sample surface (van der Waals forces, capillary forces,
Pauli repulsion, and Coulomb forces) induce a deflection of the cantilever when the tip is
scanned over the sample. The deflection of the cantilever is detected optically by a split
photodiode, which measures the position of the reflected laser beam from the cantilever
(the scheme is shown in Figure 20). A typical measurement mode is an intermittent mode,
sometimes also referred to as the AC mode due to the oscillation of the tip close to its
resonance frequency and the alternating contact of the tip with the surface, which takes
place right at the turning point of the oscillation. This allows the best possible resolution
while maintaining the tip durability as well as the sample integrity since any friction
between the tip and the sample is, in principle, avoided. When the cantilever and sample
are close to each other during each oscillation, the tip moves through an interaction
potential created by long-range attractive and short-range repulsive forces between the
sample and tip. This causes changes in the amplitude, phase and resonance frequency
of the oscillating cantilever and thus topography, amplitude and phase information
Experimental Chapter 3
38
(material dependent properties) are collected simultaneously. Further details about atomic
force microscopy can be found, e.g. in55. All the topographical images of the samples
discussed in this work were recorded using an Agilent 5500 scanning probe microscope
enclosed in an isolated acoustic chamber. The atomic force microscope probe consisted
of a silicon tip with an apex diameter of around 10 nm and a cantilever with the resonant
frequency of about 12 kHz and was operated in the AC mode.
Figure 20. A basic principle of atomic force microscopy [image was taken from56].
3.3.2 Optical and magneto-optical spectroscopy
Within the presented work, two spectroscopy methods were employed to characterise the
samples, i.e. spectroscopic ellipsometry (SE) and magneto-optical Kerr effect (MOKE)
spectroscopy.
Variable Angle Spectroscopic Ellipsometry
Spectroscopic ellipsometry measurements were performed using M-2000 ellipsometer
from J. A. Woollam over the same spectral range of 0.7 eV to 5 eV. The complex refractive
index was deduced from the measured Ψ and Δ spectra by the modelling and simulation
tool completeEASE®. For reliance in the modelled results, the spectra are measured for
different angles between 50° - 70° in steps of 5° and different thicknesses of the
investigated layer. The multi-sample analysis with coupled optical constants (n and k) was
performed to fit the model to the data with keeping only the thickness parameter of CoFeB
to vary independently in the close range to XRR measured thicknesses. A similar
procedure was followed for both the annealed and as-deposited samples. The detailed
Experimental Chapter 3
39
description of optical models used in this thesis is presented in the results section (section
4.2 and 4.3) as the optical modelling approach has to be tailored according to the layer
stack of the investigated samples.
Figure 21. The spectroscopic ellipsometry setup M2000 for measuring the complex ratio (𝜌) of the Fresnel reflection coefficients.
Magneto-Optical Kerr Effect Spectroscopy
The polar Kerr rotation (θK) and ellipticity (ηK) spectra were measured with a home-built
MOKE spectrometer. The spectrometer measures both rotation and ellipticity
simultaneously by means of the polarisation modulation method using a piezo-birefringent
modulator, also known as a photoelastic modulator (PEM). A xenon lamp (75 W) is used
to provide the necessary spectral range from near-infrared (0.7 eV) to ultraviolet (5 eV)
regime. The sample is illuminated with the linear polarised light through a bore in the pole
shoe of the magnet. The reflected light is then guided through the PEM followed by an
analyser (rotated by 45° with respect to the polariser), and a monochromator (see Figure
22). The setup is equipped with an InGaAs diode for the IR spectral region (0.7 eV to
1.4 eV) and a photomultiplier for the visible to UV region (1.5 eV to 5 eV). The spectra
discussed in the thesis were measured at RT in the photon energy from 1.5 eV to 5 eV
with an applied magnetic field ~1.7 T, if not specified otherwise. A detailed discussion on
the extraction of θK and ηK from the measured voltage on the photodetector can be found
in reference57.
The off-diagonal component (εxy = ε1xy + i ε2xy), reflecting the magneto-optical response of
the films, is calculated from the recorded θK and ηK, by designing the layered optical model
using the optical constants of the constituent layer and fitting θK and ηK spectra with a
method described elsewhere58.
Experimental Chapter 3
40
Figure 22. Schematic diagram of the magneto-optical Kerr effect spectrometer used in this thesis for measuring Kerr rotation (θK) and ellipticity (ηK).
3.3.3 Magnetometry
The magnetometry measurements were performed with the SQUID-VSM MPMS-3 from
Quantum Design and the NanoMOKETM2 system from LOT-Quantum Design GmbH.
Superconducting Quantum Interference Device–Vibrating Sample Magnetometry
The SQUID-VSM magnetometer is the most sensitive magnetic characterisation
technique developed to date. This technique is based on two principles: quantisation of
the magnetic flux in a superconducting loop (with the magnetic flux quantum ∅ = h/2e)
and the usage of a pickup coil linked with parallel Josephson junctions.
In a SQUID-VSM magnetometer, the sample is vibrated through an external magnetic
field. The change in the magnetic flux created by the moving sample is measured with a
superconducting pick-up coil which converts the changes in magnetic flux to a periodic
voltage signal. This voltage is then precisely measured by using two parallel Josephson
junctions. When a constant biasing voltage is applied on the two parallel Josephson
junctions the current flowing through both arms is equal. The magnetic flux created by the
moving sample imbalances the current flowing through the two junctions. An additionally
applied bias voltage is required to reestablish the equilibrium loop59. The required bias
provides the means to measure the magnetic moment of the sample. The magnetisation
is usually specified as the magnetic moment per volume of the measured sample. The
volume of the sample is determined by taking the area of the substrate and the thickness
of the magnetic layer into account. Figure 23 shows the cross-sectional view of the
Experimental Chapter 3
41
SQUID-VSM. The substrate, the sample holder, and the glue generate a diamagnetic or
paramagnetic background signal, which is usually recorded along with the M(H) loop data.
Such a (dia-) paramagnetic background was subtracted using the data post-processing
software Origin® 2018. Apart from the background signal, the measurements performed
with the MPMS SQUID-VSM reveal some additional artefacts, for instance, the
remanence magnetisation, flux creep being present at low magnetic fields. These
artefacts were tackled by degaussing and/or warming up the superconducting coil
magnets. As the magnetic field in MPMS is not directly measured with a Hall sensor but
via the field-to-current conversion factor (B/I ratio) of the superconducting coil, a regular
calibration is required. Further discussion of these artefacts are beyond the scope of the
present thesis but can be found in detail elsewhere60,61.
Figure 23. A cross-sectional view of the utilised MPMS SQUID - VSM MPMS-3 setup [image taken from62].
Magneto-optical Kerr magnetometry
In order to investigate magnetic the properties of small patterned samples, for instance,
the isolated square patterns of (0.5 × 0.5) mm2 used in the study of various parameters
of laser irradiation, the NanoMOKETM2 magnetometer proved to be beneficial. Such
samples are not suitable for SQUID-VSM measurements since the measured moment is
then averaged over the entire sample. NanoMOKETM2 enables to probe the local
magnetisation with a typical spot size of 100 µm. A schematic image of this setup is
shown in Figure 24.
As discussed in section 2.3.5, the reflected light from a magnetic sample changes its
polarisation state depending on the sample magnetisation. In Nano-MOKETM2, the Kerr
Experimental Chapter 3
42
rotation angle (θK) is estimated using the nulling method. Here, a linear polarised light
beam is incident on the sample and reflected at an angle of 45°. The rotation angle of the
main axis of polarisation ellipse after reflection with respect to the initial polarisation
direction is estimated as follows: initially, the analyser is rotated to minimise the signal on
the photodetector in the absence of magnetic field; subsequently, a magnetic field of up
to 500 mT is applied in the sample plane using an electromagnet. The rotation of the
polarisation axis induced by the sample magnetised in the applied magnetic field will
cause a change in the light intensity on the photodetector. With this technique, only
relative changes in the magnetisation can be measured, but no absolute values.
Moreover, the sensitivity and information depth is subject to the penetration depth of the
light in metal films. The NanoMOKETM2 system is capable of performing both longitudinal,
transverse and polar MOKE measurements. For the experiments discussed in this thesis,
only the longitudinal configuration of MOKE was utilised.
Figure 24. Top view of the Nano-MOKETM2 magnetometer. The yellow arrows point out the optical components in the optical path of the laser shown with the red dashed line.
3.3.4 First order reversal curves (FORC) method
In order to understand the response of the magnetic field sensor, it is crucial to understand
the magnetic properties of the individual constituent magnetic layers. For instance, the
magnetic reversal (switching field distribution) and offset in the hysteresis loop of the free
layer is strongly influenced by the magnetic state of the pinned layer. The conventional
methods such as M(H) loop measurements (performed by SQUID-VSM or MOKE) are
very powerful but can be ambiguous in the case of a complex multilayer stack. In such a
case, the M(H) loop is a combination of several magnetisation reversals of the various
Experimental Chapter 3
43
constituent layers. It is desirable to measure the reversal of the individual layers and to
separate the influence on the magnetic properties of the layers coming from the interaction
between the layers. Such shortcomings of the M(H) loop can be removed by using a type
of hysteresis measurement technique called Magnetic First Order Reversal Curves (M-
FORC)63. The M-FORC technique is extensively used by researchers studying
palaeomagnetism; however, its strength was explored only in a few studies in the field of
nano-magnetism and thin film magnetism. In the year 1999, C.R Pike et al. demonstrated
the application of M-FORC for characterising interactions among fine magnetic particles64.
Nowadays the technique is used as a qualitative fingerprinting65,66 method of the
reversible and irreversible switching67,68, interactions69,70, and coercive field
distributions71, etc. for various types of magnetic samples. This technique is not restricted
to the magnetisation measurements but has been modified and implemented to study the
hysteretic magnetoresistance72 (MR-FORC) and magneto-optical73 (MOKE-FORC)
properties of magnetic multilayer samples. This work focuses only on the magnetisation
measurements for recording the reversal curves. For simplicity reasons, hereafter M-
FORC will be referred to as FORC.
Measurement and analysis of FORC diagram
FORC loops are progressively recorded semi-minor loops of the magnetisation as a
function of the applied magnetic field. The sample is first magnetically saturated by
applying an external magnetic field equal to Hs (cf. Figure 5). The magnetic field is then
reduced to a certain reversal field Ha and swept back to Hs in regular field steps of Hb.
This procedure is repeated for several values of Ha generating a series of magnetisation
loops with M(Ha, Hb), known as FORC loops. Often, the values Ha and Hb are chosen to
be regularly spaced, enabling to plot M(Ha, Hb) on an equally spaced grid of Ha and Hb.
The FORC distribution is the second-order derivative of M(Ha, Hb) with respect to Ha and
Hb, respectively (see Eq .3-5). The calculated derivative when plotted as a function of Ha
(x-axis) and Hb (y-axis) generates a contour plot called FORC plot, but it is common
practice to rotate the plot by changing coordinate axes from Ha and Hb to HC = (Hb–Ha)/2
and HI = (Hb–Ha)/2, which represent coercive field and interaction field distribution.
𝜌𝐹𝑂𝑅𝐶 = 𝜕2𝑀(𝐻𝑎, 𝐻𝑏)
𝜕𝐻𝑎𝜕𝐻𝑏
3-5
Experimental Chapter 3
44
As the FORC plots are the second-order derivative of magnetisation, even minute
changes in the magnetisation will significantly be enhanced. Therefore, very sensitive and
accurate measurements of magnetisation are required. Additionally, a regressive
smoothing factor (SF)74 is implemented in the FORC numerical differentiation algorithm
in order to suppress noise-related artefacts. The FORC results presented in this thesis
were processed with ButterFORC, now called XFORC75. This software was kindly
provided by Dr. Xiang Zhao and was modified to adapt the requirements of the system
studied in thesis work. The FORC loops measured at RT for
Si / SiO2(100 nm) / Ni81Fe19(20 nm) using SQUID-VSM (for Ha = Hb = 0.5 mT) are shown
in Figure 25 a and the FORC contour showing the reversible and irreversible
magnetisation for this sample in Figure 25 b. The maximum (in red) shown in the contour
plot at HC = 2.2 mT (at HI = 0 mT) denotes the shading ‘‘irreversible’’ magnetisation
components of the Ni81Fe19 film. The distribution around this contour presents the
distribution of the coercive field. Also, this ridge is symmetric around HI = 0 mT and
suggests the absence of any exchange bias.
Figure 25. An exemplary first order reversal curve (FORC) recorded for Si / SiO2(100 nm) / Ni81Fe19(20 nm). The reversal magnetic field (Ha) and regular magnetic field (Hb) are shown in red and black solid dots, respectively (a.). The major hysteresis loop (MHL) is shown in the red dashed line. (b.) The FORC distribution calculated from the measured FORC loops. The contour denotes the maxima of the distribution at the “irreversible” located at about HC = 2.2 mT, HI = 0 mT.
Within the course of this thesis, the measurement routines of FORC have been developed
and implemented for the analysis of the reversal behaviour of permalloy layers and, in
addition, for Fe3O4 nanoparticles used as targeted drug delivery systems76. The
advantage of this method will be extended to study the interaction of the free and pinned
layer in MTJ devices as a function of the annealing process.
Experimental Chapter 3
45
3.3.5 XPS depth profiling
The chemical composition of the layers and interfaces in some of the layer stacks
discussed in this thesis was investigated with XPS depth profiling using a Thermo
Scientific NexsaTM surface analysis system. The samples were iteratively etched with a
built-in monoatomic and gas cluster ion source (MAGCIS), mounted at a 60° angle to the
normal incidence. Monoatomic Ar+ ions were accelerated to a constant energy of 1 keV,
onto a scanning area of (3.0 × 1.5) mm², leading to a mean sputter rate of approximately
0.045 nm·s-1. In a multi-element layer stack, the sputter rate cannot be maintained
constant throughout the depth profile; therefore, a fixed argon etching iteration time of
20 s was chosen, with a subsequent collection of XPS spectra. A dwell time of 5s was
given after each sputtering step to settle down the etched material. The iterations were
repeated until the SiO2 / Si substrate was reached. The XPS spectra were recorded using
the monochromatic Al Kα (1486.6 eV) radiation from the X-ray source, which was focused
on a spot diameter of 300 µm in the middle of the sputtered area. The emitted electrons
were collected at the normal angle at the pass energy of 151 eV, providing an energy
resolution of ~1.4 eV (full width at half maximum (FWHM) of the Ag3d5/2 peak). To prevent
charging of the sample, a dual-source argon-ion/electron flood gun with an electron
energy of ~0.3 eV was used. Ta4f, Co2p, Fe2p, B1s, Ir4f, Mn2p, Ru3d, Si2p, and O1s
XPS core-level spectra were collected in SnapShot mode after each sputtering step.
Finally, using Thermo Scientific AvantageTM acquisition and data analysis software, the
recorded data were fitted, taking into account all the investigated elemental peaks. Those
fits were merged to obtain the distribution profile of elements as a function of the sputtering
time (depth).
Figure 26. Process flow diagram of XPS-depth profiling.
The electrodes: 3d transition metal boride Chapter 4
46
Chapter 4: 3d-transition metal boride layers:
structural, electronic, and magnetic properties
In this chapter, layers based on 3d transition metal boride – CoFeB, which should act as
ferromagnetic electrodes in the MTJ, are investigated with spectroscopic techniques in
order to characterise their optical and magneto-optical response and to extract information
about the underlying electronic structure, conductivity, and magnetic properties.
Additionally, it sheds light on the influence of crystallisation of CoFeB alloy on such
properties. For the ease of presentation, the results are categorised in two sections
dealing with a thick and thin film, respectively. Some of the results presented in this
chapter were published in Physical Review B, DOI: 10.1103/PhysRevB.101.054438 and
Journal of Physics: Condensed Matter, DOI: 10.1088/1361-648x/ab4d2f.
4.1 Introduction
In the last few decades, 3d-transition metal borides have gained certain interest due to
their highly customisable mechanical, electrical, thermal, and magnetic properties
compared to generic 3d-transition metals or alloys77–79. One of such 3d-transition metal
borides is CoFeB, which has received special attention from a fundamental research point
of view, but also in industrial applications. The increasing interest in CoFeB alloys relates
to their atypical properties, such as structurally smooth growth80, soft magnetic
properties81, high spin-polarisation82, and very low Gilbert damping83, which makes them
especially suitable for magnetic tunnel junction devices84. By exploiting the benefits
mentioned above, S. Ikeda et al. presented a milestone improvement in tunnel
magnetoresistance ratio in CoFeB based MTJ of 355% at RT85. The improved TMR ratio
was ascribed to the improvement of the texture of the MgO barrier and the CoFe upon
post-annealing due to amorphous growth of the CoFeB during deposition. In the same
year, D. D. Djayaprawira et al. reported that a 20% inclusion of B in the CoFe maintains
it amorphous upon the deposition, thereby preventing any lattice mismatch issues at the
interface with MgO80. This allows the thin MgO to grow with a well-defined (001) texture
and to serve as a template for CoFeB crystallisation during a post-deposition annealing
process. In 2008, a TMR ratio of ~600% at RT was reported in Co20Fe60B20 / MgO
/ Co20Fe60B20 annealed at 525°C86. On the other hand, it has also been observed that
annealing at higher temperatures, although preferential from a crystallisation point of view,
could induce interlayer diffusion in the MTJs (see chapter 6), resulting in a degradation of
The electrodes: 3d transition metal boride Chapter 4
47
the TMR ratio. Therefore, it is interesting to understand the influence of annealing
temperature and composition in the crystallisation of CoXFe(80-X)B20 alloys in detail.
Previous studies mainly focused on the crystallisation of CoFeB alloys using conventional
techniques like X-ray diffractometry (XRD)87, transmission electron microscopy (TEM)88
and also some indirect methods such as resistivity, magnetoresistance89 or magnetometry
measurements90. However, these techniques inherit some limitations regarding the
sample volume required for obtaining a reliable signal (e.g. XRD), are invasive (TEM),
require complex microfabrication processes for realising devices (magnetoresistance
measurements) or provide only an indirect indication of the crystallisation (electrical and
magnetic measurements).
In the first part of this chapter, the dielectric tensor (including the diagonal optical and off-
diagonal magneto-optical components) of Co50Fe50, Co40Fe40B20, and Co60Fe20B20 are
presented, highlighting the changes in the spectra with composition and annealing
temperature. For the ease of modelling the magneto-optical constants, at first thick films
are discussed in section 4.2. These films can be treated as bulk-like, considering the
penetration depth of light in the considered spectral range. The influence of B inclusion
on the dielectric function of Co50Fe50 was modelled, supporting that the major changes in
the optical and magneto-optical spectra of CoXFe(80-X)B20 with annealing temperature are
due to changes in the crystalline structure. These results were further supported by
investigating the crystallisation of the layers using XRD and sheet resistance
measurements. Ellipsometry has proven to be a very sensitive method for investigating
thin films91, also providing the possibility of probing changes in the crystalline structure,
as investigated for Si92, diblock copolymers93 or organic photovoltaic devices94. So far,
only a few studies have regarded the optical and magneto-optical properties of CoFe
alloys95–97 or CoFeB98,99.
4.2 Thick films
Magnetron sputtering was used to deposit 100 nm of CoXFe(80-X)B20 in two stoichiometries
(X = 40 and 60), as well as Co50Fe50 on silicon wafers with native silicon oxide. The layers
were capped with 5 nm of Pt to prevent oxidation of the CoFeB. The deposition was
performed at RT with a base pressure below 2×10-4 Pa and Ar working pressure of
0.35 Pa. The wafers were diced in (1 × 1) cm2 pieces, and each piece was annealed for
The electrodes: 3d transition metal boride Chapter 4
48
30 min at temperatures in the range of 300°C to 600°C in steps of 50°C in the macro-
MOKE system. The Co50Fe50 sample served as a standard for crystalline CoFe, and no
further annealing was performed.
The diffractometry measurements (XRD, grazing incidence XRD (GIXRD), and XRR)
were conducted using the SmartLab diffractometer from Rigaku, in order to probe the
crystallisation and thickness of the films. The surfaces of the samples were studied by
atomic force microscopy. The sheet resistance of all samples annealed at different
temperatures was measured with the four-point probe (4-PP) measurement technique.
In order to extract the dielectric function (εxx = ε1xx + i ε2xx) of CoFeB from the measured Ψ
and Δ spectra, an optical model analogous to the physical layer structure was constructed.
Thus, “Si / SiO2(1.8 nm) / CoFeB(tCoFeB) / Pt(tPt) / surface roughness” layered optical
model was built using the reported dielectric functions of Si100, SiO2100, and Pt101 layers.
Additionally, the layer thicknesses determined by XRR (see Table 2) and the surface
roughness determined by AFM were used in the optical model and were kept unchanged
throughout the analysis. The unknown dielectric function of CoFeB was addressed as a
parameterised dielectric layer composed of a Drude oscillator function to account for free-
charge-carrier driven transitions and two Lorentzian oscillators to describe the dispersion
arising from interband transitions. This model was further adjusted in terms of the Drude
and Lorentzian parameters to respond to the structural changes resulting from the
annealing.
Table 2. Nominal and XRR determined thicknesses of the CoXFe(80-X)B20, Co50Fe50, and Pt layer for as the as-deposited samples
Sample ID Nominal thickness XRR thickness
CoFeB Pt CoFeB Pt
Co50Fe50 70 nm 5 nm 61.5±5 nm 3.7±0.3 nm
Co40Fe40B20 100 nm 5 nm 97.4±5 nm 4.9±0.3 nm
Co60Fe20B20 100 nm 5 nm 104.6±5 nm 4.8±0.3 nm
The electrodes: 3d transition metal boride Chapter 4
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4.2.1 Structural properties
Figure 27 presents the XRD θ-2θ scans of the CoFeB samples annealed in vacuum at
different temperatures. The pronounced CoFe(110) reflex observed at 500°C and above
indicates the crystallisation of the films. In accordance with previous studies, the creation
of a crystalline alloy from the initial CoFeB compound occurs while boron diffuses out of
the lattice, resulting in CoFe crystalline grains surrounded by amorphous boron102,103. A
closer look at the diffractograms indicates that Co40Fe40B20 crystallises in a polycrystalline
and polytextured fashion, as the present (110) and (211) peaks correspond to different
crystallographic orientations of the body-centred cubic (bcc) CoFe. Co60Fe20B20, on the
other hand, reveals a strong (110) texture, with more intense (110) and (220) peaks
occurring. The strong (110) texture was also confirmed by additional rocking scan analysis
of the (110) out-of-plane crystallite orientation distribution (see Figure 30 and the related
discussion).
At 550°C and above, a shift and broadening of the Pt(111) peak are found for both
stoichiometries, most probably suggesting a degradation of the Pt layer, possibly due to
alloying or intermixing at the interface with CoFeB. Here, it is worth mentioning that three
stray reflexes at 61°, 43° and 97°, as indicated in the figures are from silicon (400) due to
the Cu-Kβ radiation, Ag(200), and Ag(400) from silver epoxy, respectively.
Figure 27. X-ray diffraction patterns recorded for Si / SiO2(1.8 nm) / CoFeB(100 nm) / Pt(5 nm) before and after annealing under UHV at the indicated temperature for Co40Fe40B20 (a.) and Co60Fe20B20(b.). Additionally, the scan of the as-deposited Si / SiO2(1.8 nm) / CoFe(100 nm) /Pt(5 nm) sample is presented in black in figure 27(a.).The respective reflexes of constituent materials are marked by dotted lines along with the miller indices [powder diffraction file of CoFe (00-049-1567), Pt (00-004-0802), Si (00-027-1402), and Ag (00-004-0783) from the International Centre for Diffraction Data (ICDD)].
The electrodes: 3d transition metal boride Chapter 4
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The observation of a distinct out-of-plane (110) texture formation for the Co60Fe20B20 film
is further confirmed by additional rocking scans (omega scans at fixed detector position
2θCoFe(110)), providing an estimate for the crystal misalignment, the so-called mosaicity.
Figure 28 shows the rocking scan profiles for the Co40Fe40B20 and Co60Fe20B20 film, which
exhibit a clear prefered out-of-plane (110) texture (or in general, crystalline phase)
formation for the latter composition, once annealed at 500°C or higher. The transition from
the amorphous to the (110) textured structure occurs very suddenly, as confirmed by the
dramatic shape change of the rocking scan profile from 450°C to 500°C. After an initial
FWHM of about 12° after annealing at 500°C, the (110) texture of Co60Fe20B20 improves
further to an FWHM of below 9° after annealing at 600°C. In contrast, the Co40Fe40B20
rocking scan reveals no peak and only an angular independent increase in the rocking
scan intensity at 600°C can be observed, confirming random, polycrystalline CoFe
crystallisation. Therefore the omega scans of the samples annealed at the temperatures
below 550°C were not measured as no peak is expected for these samples.
Figure 28. Rocking curve measured at the CoFe(110) reflex for the Co40Fe40B20 (a.) and Co60Fe20B20 (b.) films. The inset in (b.) shows the FWHM of particular Gaussian fits of the obtained peaks.
The vertical coherence lengths, corresponding to the crystallite size in the normal direction
to the sample surface, were calculated from the FWHM of the CoFe(110) peak using the
Scherrer formula (see section 3.3.1). For the investigated films, a maximum crystallite
size of around (25 ± 2) nm is obtained for the annealing temperature of 600°C, as shown
in Figure 29 a, which is consistent with previously reported studies on 100 nm thick CoFeB
films88. As detected by cross-section scanning electron microscopy (SEM) studies
performed in a trench made by focused ion beam (FIB) (see Figure 29 b), the CoFeB alloy
does not fully crystallise within the 30 min applied annealing steps. The crystallisation
The electrodes: 3d transition metal boride Chapter 4
51
starts from the top interface with Pt and expands for (25 – 30) nm, in agreement with the
vertical coherence length of the crystallites determined from XRD. No remarkable
differences in the SEM micrographs were observed for Co40Fe40B20 and Co60Fe20B20;
thus, exemplary images for Co40Fe40B20 only are presented.
Figure 29. The CoFe crystallite sizes (vertical coherence lengths) calculated using the Scherrer equation for Co40Fe40B20 and Co60Fe20B20 (a) determined from XRD shown in Figure 27. SEM micrograph collage of Si / SiO2(1.8 nm) / Co40Fe40B20(100 nm) / Pt(5 nm) before (above) and after annealing (below) recorded in the FIB trench at 36° stage tilt (b.).
To investigate near-surface changes in the CoFeB layers further and to avoid the intense
peak from the silicon substrate, GIXRD was performed at a fixed angle of incidence of
Ω = 1°, shown in Figure 30. All the aforementioned CoFe peaks plus the CoFe(200),
which was hidden by the Si substrate peak for symmetric θ-2θ scans, are present after
annealing at 500°C, 550°C and 600°C for Co40Fe40B20, as well as for Co50Fe50 in the as-
deposited state (see Figure 30 a), reasserting the polycrystalline nature of the CoFe alloy
in this composition. On the contrary, for Co60Fe20B20 shown in Figure 30 b, none of the
CoFe peaks are detected, as the Bragg condition is not fulfilled for any detector angle due
to the pronounced CoFe(110) texture and the fixed incident angle of Ω = 1° (cf. Figure 27
and Figure 28, as well as the corresponding discussion). The Pt passivation layer
deposited on the top of Co40Fe40B20 layer exhibits polycrystallinity, although annealing at
temperatures larger than 500°C induces obvious peak shifts in both, symmetric and
grazing incidence XRD scans. On the other hand, the Pt layer on top of Co60Fe20B20 reveal
pronounced (111) texture as seen from the symmetric θ-2θ scans shown in Figure 27 with
comparably large Pt(111) and even obvious Pt(222) peaks, whereas in grazing incidence
only (111) and (220) peaks occur, but no Pt(200). The peaks from diagonal lattice planes
occur due to the pronounced (111) texture of Pt, further considering a random in-plane
The electrodes: 3d transition metal boride Chapter 4
52
crystal arrangement (fibre texture), the grazing incident angle and particular mosaicity of
the crystallites. Comprehensively summarising all structural measurements introduced
before, it seems that the Pt texture serves as a template for the CoFe crystallisation taking
place at high-temperature annealing, as it was clearly shown that the CoFe crystallisation
starts from the top interface with Pt and the observed differences in the CoFe texture were
also found in Pt, whose crystal structure further shows distinct distortion for annealing
above 500°C with converging Pt and CoFe peak positions.
Figure 30. The grazing incidence diffractogram recorded for CoFeB thick films before and after annealing at various temperature for two stoichiometries: Co40Fe40B20 (a.) and Co60Fe20B20 (b.). The respective reflexes of the constituent materials are marked by dotted lines along with the Miller indices.
4.2.2 Optical properties
Spectroscopy ellipsometry
Figure 31 shows the evolution of Ψ and Δ spectra for the sample with Co40Fe40B20 and
Co60Fe20B20 annealed at different temperatures. For clarity, only a selection of the data is
shown. It can be seen in the Ψ spectra for both sets of samples that no significant change
occurs up to 400°C, suggesting that no structural changes have occurred so far. However,
upon annealing at 450°C, sharpening of spectral feature in Ψ spectra between 1.0 eV to
2.5 eV can be observed for both sets of samples, which becomes more prominent with a
further rise in annealing temperature. Similarly, a gradual increase in the amplitude of Δ
spectra can be observed for all the samples. In comparison to the XRD results, where the
first noticeable evidence of crystallisation is visible only at 500°C for the two sets of
samples, this suggests higher sensitivity of the SE with respect to structural changes.
Furthermore, the comparison of Ψ and Δ spectra measured for as-deposited Co50Fe50
and annealed Co40Fe40B20 and Co60Fe20B20 shows a great resemblance of the spectral
The electrodes: 3d transition metal boride Chapter 4
53
feature developing between 1.0 eV to 2.5 eV upon annealing, suggesting the
crystallisation of CoXFe(80-X)B20 into a CoFe phase.
Figure 31. The evolution of and (inset) spectra recorded for the Co40Fe40B20 (a.) and Co60Fe20B20 (b.)
thick films before and after annealing at various temperatures. For comparison, the and spectrum of Co50Fe50 is plotted along with the Co40Fe40B20 spectra in (a.).
Using the optical model discussed in section 4.2, the complex dielectric functions (ε1xx and
ε2xx) were determined for the two investigated CoFeB stoichiometries and Co50Fe50. For
the ease of discussion, the spectra can be divided into two main regions: (i) the near-
infrared (NIR) region below 1.0 eV (NIR), accounting for intraband transitions, and (ii) the
visible and ultraviolet (UV) region above 1.0 eV, related mainly to interband contributions.
The NIR region of the spectrum is described by a Drude type contribution, related to the
free-electron absorption in Co50Fe50, and will be discussed in more detail later in this
section.
In the case of Co50Fe50, the Drude contribution is followed by a broad structure centred at
around ~1.5 eV in the ε1xx spectra (corresponding feature at ~2.0 eV in ε2xx spectra), see
Figure 32. This feature has been previously ascribed to the hybridisation of p and d
orbitals, resulting in direct transitions of occupied d- and unoccupied p-states in CoFe
alloys with a bcc crystalline phase96,97.
In order to understand the influence of B inclusions on the optical properties of the
Co50Fe50, the complex dielectric function of (Co50Fe50) + B was simulated. For this
purpose, the Bruggemann effective medium approximation approach was used to
calculate the optical constants of the mixed material with the host matrix of Co50Fe50
containing B inclusions. In this approach, 15% of the (Co50Fe50) + B film volume is
assumed to be a spherical inclusion of B into the metallic Co50Fe50. However, it should be
The electrodes: 3d transition metal boride Chapter 4
54
noted that this is only a coarse approximation of the actual situation; since previous
studies suggested that B migrates to CoFe grain boundaries or to the neighbouring
layers102,103. The most obvious change induced to the dielectric function spectra of
Co50Fe50 by the B inclusion is visible in the ε2xx spectrum, namely a decrease of the
absolute values. Since B is a non-metallic material, its addition to the metallic Co50Fe50
increases the dielectric losses. Consistently, the values of ε1xx increase, indicating an
increase in the relative permittivity of CoFe-B. A slight broadening of the spectral features
is also observed, but rather negligible when compared with the changes in the features of
the dielectric functions of the CoFe-B alloys extracted from the experimental ellipsometry
spectra before and after annealing, see Figure 33. The good correspondence between
the simulated complex dielectric function of B incorporated in CoFe and the dielectric
function determined for the Co40Fe40B20 indicates that during the crystallisation process
CoFe crystallites are formed and B migrates outside the crystallites, i.e. to the grain
boundaries. This scenario is in line with the results of previous studies of the local
structure of CoFeB102 and of crystalline CoCrPt-B alloys used for recording media in hard
disk drives104.
Figure 32. The complex dielectric function (ε1xx & ε2xx) spectra of the Co50Fe50 (red), Co40Fe40B20 (blue) and B (grey) 105, together with the simulated ε1xx & ε2xx of (Co50Fe50)+B with 15 % B inclusion(yellow). More detailed information about the B inclusion is given in the text.
The dielectric functions of the as-deposited CoFeB alloys present only weak and very
broad spectral features, which gradually become more pronounced with an increase in
annealing temperature, as shown in Figure 33. The characteristic spectral feature of
Co50Fe50 at ~1.5 eV occurs in the ε1xx spectra for the samples annealed at 450°C. This
suggests that 450°C is the onset temperature for crystallisation. As the optical
spectroscopy has an information depth limited to a few 10 nm, the changes visible in the
The electrodes: 3d transition metal boride Chapter 4
55
spectra at 450°C indicate that the crystallisation takes place near the surface
(CoXFe(80- X)B20 / Pt interface), as supported by scanning electron microscopy (SEM)
images (cf. Figure 29 b) demonstrating nucleation at the Pt interface. Noticeably, the
pronounced CoFe reflex was observed in XRD scans starting at 500°C, indicating that the
optical spectroscopy allows probing the incipient phase of crystallisation with very small
crystallites. In fact, a remarkable resemblance of the dielectric function of the Co40Fe40B20
after annealing at 600°C and the as-deposited Co50Fe50 is found, which is consistent with
the similarities in the crystalline structure observed with XRD. This suggests that at 600°C
B diffuses completely out of the CoFe crystallites. The systematic decrease in ε2xx with
annealing temperature is furthermore consistent with a greater ordering within the films
due to crystallisation. The characteristic spectral feature in ε1xx spectra of Co60Fe20B20 is
red-shifted relative to Co40Fe40B20, probably due to the difference in the stoichiometric
composition. Additionally, comparing the amplitudes of the ε2xx spectra (mostly <1.0 eV)
of the two stoichiometries, it is evident that the Co40Fe40B20 has lower dielectric losses
due to the free electrons in comparison to Co60Fe20B20. This, in turn, implies that
increasing Co concentration increases the charge carrier concentration, which is
consistent with the empirical finding that the resistivity of Co is almost half of that of
Fe105,106.
Figure 33. The annealing temperature dependent evolution of ε1xx (a.) and ε2xx (b.) spectra for Co40Fe40B20 (solid line) and Co60Fe20B20 (dashed line), and CoFe (black).
The classical Drude equation defines the free charge carrier concentration contribution to
the dielectric function, which in its mathematical form is equivalent to a Lorentzian
oscillator positioned at 0 eV:
The electrodes: 3d transition metal boride Chapter 4
56
휀(𝐸) = −ℏ
휀0𝜌(𝜏𝑠𝐸2 + 𝑖ℏ𝐸) 4-1
where, ε0 is the vacuum dielectric constant, ℏ is the reduced Planck constant, τs is the
mean scattering time of the free carriers between successive collisions and ρ is the optical
resistivity. The analysis of the Drude contribution to the dielectric function allows deriving
the optical resistivity (ρ) and scattering time (τs) of the investigated films. The resistivity
and the scattering time ultimately relate to the ordering state of the films, according to the
Fuchs size-effect theory107.
These parameters are shown in Figure 34 for both CoFe-B stoichiometries. The resistivity
remains barely unchanged until 400°C, followed by a maximum at 450°C and a
subsequent decrease with increasing annealing temperature. It should be noted that this
evolution cannot be explained by the B diffusion since reference108 showed that the
migration of B starts already at 200°C (see section 6.3). The decrease in resistivity can
be ascribed to an increase in ordering and decrease in the number of defects, which, in
fact, is consistent with the increase in the crystallite size derived from the XRD
measurements. The presence of a maximum at 450°C relates very likely to a temperature
of nucleation of the crystallites, where the resistivity increases due to the formation of
grain boundaries and defects, originating from the low level of ordering of the crystal. The
poly-textured phase in Co40Fe40B20 (in contrast to the well-oriented phase in Co60Fe20B20)
could arguably also explain the difference in the resistivity and scattering times between
the two alloys, since more mismatched grain boundaries and defects would lead to higher
resistivity due to the shorter mean free path, i.e. longer carrier scattering time.
Sheet resistance (R) measurements were conducted on all the samples in order to
investigate the influence of annealing on the electrical properties of the layers. The
change in sheet resistance of the CoFeB samples with annealing temperature is shown
in Figure 34 b. Up to 400°C, no significant change in the sheet resistance is found. Above
this temperature, a monotonous decrease with increasing temperature is observed,
consistent with the changes observed for the optical resistivity parameter calculated from
the Drude model (cf. Figure 34 a). The maximum at around 450°C is not very obvious as
the entire samples (including the cap layer and 100 nm thick mostly amorphous CoFeB)
are probed electrically, although crystalline CoFe occurs in the upper part only and optical
The electrodes: 3d transition metal boride Chapter 4
57
spectroscopy, as a surface-sensitive measurement technique, probes that crystalline
volume exclusively. Given the increase in crystallite size and ordering within the films with
the annealing temperature revealed by XRD, a decrease in the charge carrier (electrons)
scattering due to defects and grain boundaries is expected, which results in a decrease
of the sheet resistance. The trend in R for both stoichiometries is noticeably similar to the
change in ρ obtained from the SE measurements.
Figure 34. Drude parameters resistivity (ρ) in black and scattering time (τs) in red colour as a function of annealing temperature for Co40Fe40B20 (solid symbol) and Co60Fe20B20 (empty symbol). The lines in the figure are drawn to guide the eye (a.). Sheet resistance of the Co40Fe40B20 (filled circles in red) and Co60Fe20B20 (unfilled circles in blue) layers passivated with a Pt thin film as a function of the annealing temperature (b.).
Magneto-optical spectroscopy
Figure 35 a shows the measured θK and ηK MOKE spectra of Co50Fe50, in comparison to
θK reported by D. Weller et al. for Co48Fe52109. Even though the amplitude of θK is slightly
lower than previously reported109, the line shape of both experiments resembles closely.
In fact, the present data is closer to the first-principle calculations performed by Maurer et
al. for this CoFe composition110. The off-diagonal dielectric function of Co50Fe50 was
calculated using the procedure discussed in section 3.3.2 and is shown in Figure 35 b as
εxy (h)2, in order to highlight the spectral features111. It is well established by theoretical
studies that the spin polarised density of states of 3d-transition metals and their alloys are
fairly similar, resulting in similar electronic transitions in magneto-optical spectra96. These
spectral features noticed in the optical region of the spectrum can be explained based on
the theoretical predictions by Kwang Joo Kim et al. for Fe3Co and Co3Fe using the tight-
binding linear-muffin-tin orbitals (TBLMTO) method with the local spin density
approximation (LSDA)97,110. They assign the transition at 2.0 eV as originating mainly from
transitions from the occupied minority-spin d triplet states at lower energy into the
The electrodes: 3d transition metal boride Chapter 4
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unoccupied minority-spin p states. These d→p transitions in the minority-spin bands
become possible through p-d hybridisation.
Figure 35. Measured polar Kerr effect rotation (θK) and ellipticity (ηK) spectra of Co50Fe50 in comparison
to the previously published reference109,110 (a.) and calculated xy (h)2 as the function of photon energy (b.).
Figure 36 shows the evolution in θK and ƞK spectra for the CoFeB samples annealed at
different temperatures for the two investigated stoichiometries. Similar to the SE spectra,
no significant changes in θK and ηK spectra were observed up to 400°C. Upon annealing
at 450°C, the characteristic line shape of the θK spectrum starts resembling that of CoFe.
Annealing at higher temperatures results in the enhancement of spectral features at
~2.0 eV and ~4.7 eV. The improvement in the features upon annealing at 450°C or above
is consistent with considerable ordering in the lattice, and again, this development in the
spectra can directly be correlated with the increase in crystalline ordering of CoFe. It can
also be observed that these features are slightly red-shifted for Co60Fe20B20 compared to
Co40Fe40B20, contrary to previous theoretical calculations, where no significant differences
were found on the MOKE spectra due to different CoFe content96. In fact, this shift may
as well be related to the differences found with XRD in the crystalline structure of both
compounds. In this context, we note that besides composition, also the crystalline
environment influences significantly the magneto-optical properties of the material112. In
fact, the larger amplitude of the spectral features of Co40Fe40B20 is furthermore an
indication of higher spin-polarisation (magnetisation) for the lower Co content, which is
consistent with the calculated Slater–Pauling curve113.
The electrodes: 3d transition metal boride Chapter 4
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Figure 36. The polar Kerr effect measured rotation (θK) and ellipticity (ηK) spectra for the 100 nm thick CoFeB films before and after annealing at the indicated temperatures for two stoichiometry Co40Fe40B20 (a. & c.) and Co60Fe20B20 (b. & d.).
The calculated εxy (h)2 as a function of the photon energy for the two investigated
compositions annealed at 450°C and 600°C is shown in Figure 37. Similar to the off-
diagonal dielectric function of CoFe shown in Figure 35 b, the imaginary part of ε2xy (h)2
shows two main features in the measured spectral range, at ~2.0 eV and ~4.5 eV. As
discussed earlier, the features in the εxy spectra are ascribed to transitions from the filled
part of the 3d band to empty hybridised pz states near Fermi energy (spin-down states).
The relative shifts in the positions of the spectral features of Co60Fe20B20 to lower energies
with respect to Co40Fe40B20 suggest a dependence of the DOS on the stoichiometry. Such
a shift can be explained based on the theoretical calculations of Liu et al. related to the
electronic structure of CoFe alloys with different stoichiometry 114. Although the DOS near
the Fermi energy is similar for both Co and Fe, due to the higher electronegativity and a
smaller exchange splitting of Co, the minority-spin orbitals of Co are situated at lower
energies than the Fe orbitals. A higher percentage of Co in the alloy will, accordingly, lead
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to a shift of the electronic states of the alloy states to lower energies. The change in
amplitude of the spectral features relates, as mentioned previously, to the degree of spin-
polarisation, which is increasing in the CoFe alloys case with increasing Fe content 113.
Figure 37. The calculated xy(h)2 as a function of the photon energy for two stoichiometries, Co40Fe40B20 and Co60Fe20B20, for Si / SiO2(1.8 nm) / CoXFe(80-X)B20(100 nm) / Pt(5 nm) layer stacks annealed at 450°C and 600°C.
4.3 Thin films
The results from the thick films suggest that optical spectroscopy techniques are very
sensitive to structural changes. In this sub-chapter, we will focus on the structural and
optical properties of thin films on the example of Co60Fe20B20.
The Co60Fe20B20 thin films with nominal thicknesses ranging from 10 nm to 20 nm covered
by a 3 nm gold passivation layer were deposited by DC magnetron sputtering on thermally
oxidised silicon substrates. Similar to the thick film samples, the thickness of the
Co60Fe20B20 / Au bilayers was verified using XRR measurements on the as-deposited
samples (see Table 3). The depositions were performed at RT with a base pressure below
2 x 10-4 Pa and a working pressure of 0.35 Pa, while using Ar as process gas. The
samples were annealed for 30 min in vacuum at 350°C and subsequently for 30 min at
400°C. The first annealing step was performed at 350°C as the transformation of the thin
films from amorphous into a crystalline bcc CoFe phase is expected to occur above
325°C115, in contrast to the thick films studied before, where temperature as high as 450°C
was required for initial crystallisation. The characterisation of the samples with the same
methods as used in section 4.2, was performed in the as-deposited state and repeated
after each of the subsequent annealing steps.
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4.3.1 Structural properties
The X-ray diffractograms recorded using the 3000PTS diffractometer before and after
annealing at 350°C and 400°C are shown in Figure 38. For the as-deposited state of the
sample, only the Au(111) and Si (substrate) peaks at 38.2° and 69.1°, respectively, were
observed, indicating that the Co60Fe20B20 film is amorphous. After annealing at 350°C, no
peak corresponding to the CoFe crystalline phase was detected, suggesting that the
Co60Fe20B20 layer is still in an amorphous phase or the crystallites are not large enough
to be detected by the XRD equipment. A further annealing step at 400°C resulted in the
occurrence of a well distinguishable 2θ peak around 45°, characteristic for the CoFe(110)
crystalline phase. For this peak, a crystallite size of (8.9 ± 2) nm was calculated using the
Scherrer formula.
Figure 38. X-ray diffractogram recorded for Si / SiO2(100 nm) / Co60Fe20B20(20 nm) / Au(3 nm), in the as-deposited state (squares) and after annealing at 350°C (circles) and 400°C (triangles).
Table 3. Nominal and XRR determined thicknesses of the Co60Fe20B20 and Au layer for as the as-deposited samples
Sample ID Nominal thickness XRR thickness
Co60Fe20B20 Au CoFeB Au
1 10 nm 3 nm 10.5±0.2 nm 2.6±0.1 nm
2 15 nm 3 nm 14.8±0.2 nm 3.2±0.1 nm
3 20 nm 3 nm 18.9±0.2 nm 3.6±0.1 nm
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4.3.2 Optical properties
Spectroscopy ellipsometry
For the SE, an optical model similar to that of thick films was used to derive the optical
constants for each sample. A multi-sample analysis based on a parametric model
consisting of one Drude free electron model and a series of five Gauss-Lorentz oscillators
was employed in order to represent the fine structure of the optical dispersion in the
Co60Fe20B20 in the entire investigated spectral range. In the multi-sample analysis, the
dielectric function of the samples was considered to be identical for the considered range
of thicknesses. This model was then further adjusted to respond to the structural changes
induced by the annealing steps.
Figure 39 shows the evolution of the Ψ and Δ spectra recorded at an angle of incidence
of 65° for Co60Fe20B20(20 nm) / Au(3 nm) before and after annealing at 350°C and 400°C.
Similar to the thick films, noticeable changes can be observed with the increase in
annealing temperature. Significant changes occur already after the 350°C annealing step
in both Ψ and Δ spectra, whereas no reflex corresponding to CoFe(B) was observed in
XRD (cf. Figure 38). Here it should be noticed that the spectral features observed in the
Ψ and Δ spectra for the thin film samples cannot be compared with the ones measured
for the thick film due to interference between reflected light from the surface and interfaces
in the thin film layer stack (CoFeB-Au, CoFeB-SiO2, and Si-SiO2 ).
Figure 39. Experimental SE ψ and ∆ spectra recorded at 65° AOI for Si / SiO2(100 nm) / Co60Fe20B20(20 nm) / Au(3 nm) before and after vacuum oven annealing at 350°C and 400°C.
Similar changes can also be observed in the real (ε1xx) and imaginary (ε2xx) parts of the
dielectric function (diagonal elements of the dielectric tensor) of Co60Fe20B20 obtained
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from the SE measurements, before and after annealing, as shown in Figure 40 a. For the
as-deposited films, the dielectric function relates well to the typical line shape of the optical
dielectric function reported for amorphous Co20Fe60B20 films with thicknesses between
10 nm and 40 nm grown on Si / SiO2(1000 nm), i.e. with very broad features98. It should
be noted that thicker films of Co60Fe20B20 (100 nm) discussed previously in section 4.2,
exhibit even lower spectral features (see Figure 40 b in comparison to 20 nm thick films).
While for the modelling of thicker films one Drude oscillator and two Gauss oscillators
located at 0.5 eV and 2.2 eV were employed, in the case of the multi-sample analysis
performed for the films thicknesses in the range from 10 nm to 20 nm, additional
oscillators located at 1.4 eV, 3.3 eV, 3.7 eV and 4.6 eV had to be employed in order
achieve lower MSE values of the simulation. This apparent discrepancy between the thick
and thin film optical properties might have several reasons: possible oxidation of the top
layer of Co60Fe20B20 would have a larger contribution to the optical response of the thin
compared to the thicker films. In addition, a granular structure of the film, which is more
pronounced in the case of the thinner films or the related plasmonic effects in the thin Au
overlayer116 or in the Co60Fe20B20 might affect the thin film dielectric function.
Figure 40.The ε1xx (black) and ε2xx (red) components of the complex of the dielectric function of Co60Fe20B20 (10 nm to 20 nm) before (continuous line) and after annealing at 350°C (dashed line) and 400°C (dotted line) (a.). Comparison of the complex dielectric function of as-deposited 20 nm and 100 nm thick Co60Fe20B20 films (b.).
Upon annealing the thin films, two effects become visible: (i) a change in the low energy
slope of the ε2xx and (ii) a sharpening of the spectral features. The low energy slope of ε2xx
initially increases from the as-deposited state to the sample annealed at 350°C and then
it decreases again in the samples annealed at 400°C. The decrease in the slope of ε2xx
with annealing temperature corresponds to the characteristic behaviour of metals upon
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crystallisation and the corresponding reduction in the dielectric losses117. Liang et al.
reported merely no changes of the Co20Fe60B20 optical constants of a 40 nm film upon
annealing at 300°C for 2 h, whereby 300°C is below the crystallisation temperature of
CoFeB98.
The most significant changes upon annealing of the thin films occur at 2.6 eV and 3.8 eV
in ε1xx. The lineshape of the spectra obtained after annealing resembles the reported one
for Co20Fe60B20 films (2 nm)98. Liang et al. explained such a lineshape by the optical
anomaly occurring when isolated metal islands percolate into a continuous metallic layer.
Extrapolating to the case of the investigated system, such a structure-rich lineshape of
both ε1xx and ε2xx might be consistent with the increasing roughness of the film after
annealing. Slight oxidation of the CoFeB and/or the diffusion of the B towards the film
surface upon annealing (see section 6.3) are factors that might also influence the dielectric
function and can hardly be taken into account by modelling. Nevertheless, the occurrence
of a clear CoFe(110) peak in the X-ray diffractogram of the 20 nm film annealed at 400°C
(see Figure 38) indicates that the observed changes in the dielectric function (Figure 40)
are more likely to be related to the crystallisation of CoFe(B).
The electronic origin of the observed spectral features is discussed in the following. The
broad pronounced feature in ε2xx at ~2.6 eV was previously observed for FeCo alloys in
the bcc phase 97,115 and ascribed to direct interband transitions in the minority-spin bands
between occupied d- and unoccupied p-states118. It is worth mentioning that the optical
response may also be influenced by interband transitions from gold typically found around
2.5 eV116 or by plasmonic effects in Au or even in CoFeB. While the interband transitions
are taken into account by the optical constants for gold119 used in the discussed model,
the plasmonic effects were not considered. Since the complete removal of the Au
signature at 2.5 eV by the used optical model cannot be warranted, any comparison
should be taken with care. The energy range above 3.0 eV is dominated by the typical
interband transitions of the Fe or Co sites, including d-p as well as d-d transitions. In the
context of pure bcc-Fe, the peaks at 2.8 eV and 4.0 eV were ascribed to transitions along
with high symmetry points of the crystal lattice120. It can be assumed that this is also the
case in a CoFeB alloy, suggesting that the observed transitions are likely to be consistent
with the presence of a CoFe bcc crystalline phase (where the boron diffused out of the
crystallites), in agreement with the presence of the CoFe(110) peak in the diffractogram
of the 400°C annealed sample.
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Magneto-optical spectroscopy
The changes in the Kerr spectra before and after annealing are compared in Figure 41. A
broad feature centred at about 2.5 eV is in this case also very pronounced, becoming
clearly narrower upon annealing. The spectra compare well with previous studies on CoFe
alloys96, suggesting that the B content in the alloy may have a minimal influence on the
magneto-optical response in the studied energy range. The amplitude of the signal
improves with the annealing, similarly to the previously reported case of FePt alloys121.
Following the interpretation proposed by Cebollada et al., the enhancement of the
magneto-optical Kerr effect of an alloy upon annealing can be regarded as a sign of
improvement in the magnetic ordering of the films. Due to the strong optical interference
oscillations in the MOKE spectra of the thin films, the off-diagonal dielectric function for
these samples could not be calculated with any reasonable assumptions in the optical
model.
Figure 41. Comparison of Kerr rotation θK (black) and ellipticity ηK (red) for Si / SiO2(100 nm) / Co60Fe20B20(20 nm) / Au(3 nm) before (continuous line) and after annealing at 350°C (dashed line) and 400°C (dotted line).
4.4 Conclusion
In this chapter, the evolution of the structural, optical, and magneto-optical properties of
optically thick (100 nm) films of CoXFe(80-X)B20 (X = 40 and 60) and optically thin (from
10 nm to 20 nm) Co60Fe20B20 films were discussed.
The 100 nm films passivated with a 5 nm Pt cap layer were annealed between 300°C and
600°C. The structural and electrical properties of the films were assessed by XRD and
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electrical four-point probe measurements, respectively. The comparison of the
CoXFe(80- X)B20 dielectric function extracted from spectroscopic ellipsometry with that of
Co50Fe50 allowed to identify CoFe specific spectral features and to analyse the impact of
B on the optical properties of the CoFeB alloys. The (magneto-) optical spectroscopic
techniques are proven to be extremely sensitive to structural changes. The analysis of the
Drude component of the dielectric function of CoFeB allowed extracting information
regarding the resistivity and charge carrier scattering time, which is closely related to the
crystalline order in the films. Corroborating the results of spectroscopic ellipsometry, SEM,
and XRD, we can conclude that the nucleation of the crystallisation starts at the interface
between the CoXFe(80-X)B20 and the crystalline Pt capping layer.
The magneto-optical off-diagonal component of the dielectric function of CoXFe(80-X)B20
extracted from the MOKE spectra shows significant changes with the composition of the
alloy as well as with the structural evolution from the amorphous to the crystalline phase.
Similar to the thick film samples, the evolution in the optical and magneto-optical spectra
presented a strong correlation to the crystallisation of CoFeB. The dielectric function of
Co60Fe20B20 thin films showed more spectral features which are absent in 100 nm
samples. These features are the possible outcome of surface oxidation that may have a
greater influence on optical properties in thin films than thick films or because of the
plasmonic enhancement due to the granular gold overlayer at the surface.
This study underlines the utility of spectroscopic ellipsometry and MOKE spectroscopy for
material optimisation in the field of metallic alloys for spintronic applications. It also opens
the possibilities of such (magneto-) optical spectroscopic techniques for non-invasive and
in-situ characterisation method, which is even compatible with contemporary CMOS
microfabrication technology.
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Chapter 5: Setting exchange bias using laser vs oven
annealing techniques
This chapter presents the proof-of-concept that focused laser irradiation can be used as
a method to set a defined pinned magnetisation direction through the exchange bias effect
in MTJs. The discussion is based on the comparison of the conventional vacuum
annealing to the laser annealing for varying radiation intensity and scanning speed. The
aim is to achieve the optimal laser annealing parameters so as to set a maximum
exchange bias field strength in the given stack. Further, the efficacy of the laser annealing
as a direct-writing technique was tested by writing a text on a simplified FM/AFM layer
stack and visualised using MOKE microscopy imaging to read the text. Some of the results
presented in this chapter were published in IEEE Transactions on Magnetics, DOI:
10.1109/TMAG.2018.2873428.
5.1 Introduction
A tremendous effort has been put into the integration of spintronic devices to the prevailing
CMOS technology. The miniaturisation of the devices has pushed the researchers to
study new materials and find solutions for technological challenges. One among such
challenges is the annealing of the spintronic devices, which is intended not only to improve
the performance of the device but also to assist in tailoring the properties of the device.
For instance, in the case of MTJ sensor devices, annealing and subsequent cooling in the
presence of a magnetic field allows inducing exchange bias in the device, which can be
exploited to tune the sensing range, the directional sensitivity etc. A method that allows
to locally anneal the micron and sub-micron devices can be highly beneficial. Recently,
Albisetti et al. demonstrated an innovative method of using thermally assisted magnetic
scanning probe microscopy (SPM) to pattern and set the exchange bias in an
IrMn / CoFeB bilayer system122. However, this method suffers some severe drawbacks,
including the fact that it is extremely time-consuming due to the use of SPM, which hinders
its industrial application. The laser-based annealing method is a possible method to
generate the heat locally in the exposed region, and with the help of suitable optics, a
wide device area can be covered. This method gives the advantage of setting the
exchange bias locally and in a short period of time, even for large scale wafer
processing51,123.
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For tailoring the exchange bias in the FM and AFM layers and optimise the magnetic field
cooling parameters for setting a maximum exchange bias, a series of experiments were
conducted. The sample deposition was performed by magnetron sputtering by the
industrial partner “Singulus Technologies AG”. The layer sequence for the sample
investigated is “Si / SiO2(100 nm) / Ta(5) / CuN(30 nm) / Ta(3 nm) / Ni81Fe19(2 nm) /
IrMn(8 nm) / Co40Fe40B20(2 nm) / MgO(1.8 nm) / Co40Fe40B20(2.2 nm) / Ta(10 nm)”. The
samples were either heat-treated within the SQUID-VSM system (conventional annealing)
or locally with focused laser irradiation. The samples were characterised by SQUID or
MOKE magnetometry or, alternatively, by FORC. Details regarding the measurement
techniques are presented in chapter 3. A pictorial representation of the temperature profile
generated due to laser annealing and the selective alignment of magnetisation in
CoFeB / IrMn in a typical MTJ layer stack is shown in Figure 42.
An additional study was conducted with the aim to understand to which extent the seed
layer, in addition to the laser annealing parameters, influences the exchange bias
between CoFeB (pinned layer) and IrMn in the given stack. For this MTJs were prepared
by magnetron sputtering at Singulus technologies AG with the layer stack: Si /
SiO2(100 nm) / Ta(5 nm) / Seed layer / IrMn(8 nm) / Co40Fe40B20(2.3 nm) / MgO(1.8 nm)
/ Co40Fe40B20(2.3 nm) / Ta(5 nm), where the seed layer was chosen to be Ru(5 nm) or
Ni81Fe19(2 nm). The samples were exposed to laser irradiation by scanning square areas
of (0.5 × 0.5) mm2 in a continuous wave (CW) mode at scan speeds varying from
0.5 mm·s-1 to 5000 mm·s-1 and at laser peak intensities in the range from 60 kW·cm- 2 to
800 kW·cm-2. The exposure was later repeated for larger sample areas of (6 × 6) mm2
with scan speeds of 50 mm·s-1, 500 mm·s-1 and 5000 mm·s-1, at 600 kW·cm-2 to
1000 kW·cm-2 laser intensities. Furthermore, the influence of the number of scan
repetitions on the exchange bias field was analysed by repeating the laser annealing at
5000 mm·s-1 scanning speed and 480 kW·cm-2. For comparison, an analogous study with
conventional annealing technique was performed in the SQUID-VSM system for
temperatures ranging from 123°C to 400°C and cooled in the presence of a 120 mT in-
plane magnetic field, the same as used for all laser annealing processes. The coercive
and exchange bias fields for the pinned layer are calculated from hysteresis loops
measured using the NanoMOKETM2 setup.
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Figure 42. A schematic diagram showing a typical layer stack of the MTJ sensor and the laser-induced temperature profile for setting the exchange bias in the FM/AFM layers. The CoFeB (FM) layer next to IrMn (AFM) is the pinned layer and the upper CoFeB layer is the free layer, both sandwich the MgO tunnel barrier. The Ni81Fe19 and CuN serve as the seed and buffer layer, respectively. The Ta layers work as the passivation (on the top) and adhesion layer (at the bottom).
5.2 Magnetisation reversal of complex MTJs layer stack
The M(H) hysteresis of the “Si / SiO2(100 nm) / Ta(5 nm) / CuN(30 nm) / Ta(3 nm) /
Ni81Fe19(2 nm) / IrMn(8 nm) / Co40Fe40B20(2 nm) / MgO(1.8 nm) / Co40Fe40B20(2.2 nm) /
Ta(10 nm) layer stack is shown in Figure 43. Assuming that the magnetisation of the free
layer (FL) and the exchange biased layer (PL) are aligned in-plane at zero applied
magnetic field, the hysteresis loop of the MTJ stack can be explained qualitatively as
follows. Prior to the annealing (in Figure 43 a), for large magnetic fields applied in the
positive direction (H > 0.2 T) both FL and PL are saturated and aligned in the direction of
the external magnetic field (point a). When the field is reduced close to zero
(0.2 T > H > 0.05 T), the PL magnetisation rotates gradually to some random orientation
with no net magnetisation due to arbitrary coupling to the IrMn layer, resulting in increased
coercivity (between the points a and b), while the FL remains aligned in the direction of
the applied magnetic field. When further decreasing the magnetic field to zero and then
changing the field direction, the FL rotates (across b to c), yet PL remains randomly
oriented. The increment of the magnetic field in the negative direction also rotates the PL
magnetisation in the direction of the field so that both layers align along the magnetic field
(through c to d). A similar explanation can be drawn for sweeping the magnetic field from
negative to positive saturation, where the resulting coercivity of the pinned layer is
enhanced due to the exchange bias effect, whereas the soft-magnetic free layer reveals
small coercivity as seen from the figure.
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Figure 43. Magnetisation as a function of the in-plane applied external magnetic field on an as-grown sample (a.) and measured along the easy-axis of the bottom pinned layer after 120 mT field cooling from 227°C (b.). The arrow in green and red represent the magnetisation of the free layer and pinned layer, respectively.
Similar to the as-deposited sample, the M(H) hysteresis loop was recorded for the sample
annealed at 227°C and cooled to RT in the presence of a 120 mT in-plane applied
magnetic field, see Figure 43 b. This causes the exchange bias effect at the FM/AFM
interface with the FM magnetisation to be aligned according to the applied magnetic field.
After the annealing the M(H) hysteresis loop can be described as follows: for a large
positive magnetic field, both the CoFeB layers are saturated, and their magnetisation lies
in the direction of the magnetic field (point a). Lowering the field through zero leads to the
rotation of the FL, with a very sharp reversal (through point a and b) and making the FL
and PL anti-parallel to each other. Further increase field toward negative values rotates
the PL creating a parallel arrangement of PL and FL (at point c) again. This reversal is
only possible when the applied field overcomes the exchange bias field induced in the PL
by the annealing. Moreover, the exchange bias reduces the two symmetric hysteresis
present in I and III quadrants in one loop, that is shifted by HEB with the increased coercive
field (HC). This can be attributed to the gradual reversal of PL as a single magnetic layer
due to exchange biasing. Nevertheless, as the magnetisation reversal of the FL and the
PL can be clearly separated, the FL can rotate according to a low external field
independently with much softer magnetic properties. Thus being the basis of high
sensitive magnetoresistive measurement principles on the basis of the GMR and TMR
effect. This loop allows us to determine HEB and HC for the PL alone by measuring the
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fields at which it switches: as is conventional, HEB is the average of these two fields, whilst
HC is half the difference between them.
5.2.1 Setting the exchange bias by conventional oven annealing
In the first step, the most favourable temperature for setting the exchange bias in the MTJ
stack described in Figure 42 was investigated by thermal treatment in the SQUID-VSM
system. For this, the samples were annealed at different temperatures for 5 min and then
cooled in the presence of 120 mT magnetic fields at a constant cooling rate of 15 K·min- 1.
Figure 44 shows the magnetisation loop measured at room temperature for a pristine
sample and after each annealing process. The saturation magnetisation (Ms) was
observed to increase with annealing, which is the signature of improvement in crystallinity
in the CoFeB layers. As anticipated, with increasing annealing temperature, a prominent
exchange bias effect between the CoFeB and IrMn layer appears, resulting in a stronger
exchange bias field strength and increased coercivity of the PL. A maximum exchange
bias field of HEB = (110 ± 5) mT and coercivity of HC = (24 ± 1.5) mT was noted for the
sample annealed at 227°C and was observed to be unaffected with an increase in
temperature to 327°C. Further, a very sharp reversal centred around H = 0 mT of the FL
can be seen, which verifies that the FL is not influenced by any demagnetising field from
the PL.
Figure 44. SQUID-VSM measured M(H) hysteresis loops recorded at RT for an as-deposited sample and after consecutive field cooling process from different temperatures as indicated in the legend.
This study was further extended to determine the blocking temperature (i.e. the
temperature at which the biasing vanishes, TB), which in turn defines the working
temperature range of the final microfabricated magnetic field sensors. Initially, the sample
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was annealed to set the exchange bias. M(H) loops were then recorded at RT and further
iteratively after increasing the temperature in steps up to 327°C. With elevated
temperature and in the absence of magnetic field, the thermal energy overcomes the
exchanges anisotropy energy, resulting in a gradual reduction in exchange bias and
coercivity finally going to zero at TB. From Figure 45, it is evident that the HEB linearly
decreases, vanishing completely at TB = (267 ± 20)°C. A similar trend was observed for
HC. The measured TB is significantly lower than the TN of bulk IrMn due to the fact that for
the thin films, TB is the resultant of finite-size scaling124,125.
Figure 45. Temperature dependence of the coercivity HC (a.)and the exchange bias field HEB for the determination of the TB. The sample was initially cooled in 120 mT field in-plane applied magnetic from 327°C to RT before the measurements at increasing temperature.
5.2.2 Setting the exchange bias by laser annealing
Continuous-wave laser annealing
The results from the conventional oven annealing presented in the previous sub-section
have established a baseline for understanding the exchange bias in the MTJ layer stack.
Further experiments with localised annealing on a microscale using laser irradiation were
targeted to achieve similar results as with conventional oven annealing. Samples having
the same layers (Si / SiO2(100 nm) / Ta(5 nm) / CuN(30 nm) / Ta(3 nm) / Ni81Fe19(2 nm)
/ IrMn(8 nm) / Co40Fe40B20(2 nm) / MgO(1.8 nm) / Co40Fe40B20(2.2 nm) / Ta(10 nm)) were
irradiated with a laser beam using the experimental setup described in section 3.2.3. For
these experiments, square areas-of-interest of (0.5 × 0.5) mm2 were heated by raster
scanning with continuous laser irradiation (CW) in the presence of an external magnetic
field of 120 mT, aligned in-plane. The optical images of the laser annealed samples with
test series about the laser scanning speed and different laser intensities are shown in
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Figure 46. The magnetic characterisation for the irradiated areas was conducted with
MOKE magnetometry measurements in longitudinal geometry.
a.
b.
c.
d.
Figure 46. Optical micrograph of an exemplary sample annealed with CW irradiation with different scanning speed, 50 mm·s-1 (a.), 500 mm·s-1 (b.), 2000 mm·s-1 (c.), and 4000 mm·s-1 (d.). In each square, the different laser power was used as indicated, the patterns, as well as the labels, were created by laser ablation.
Figure 47 exhibits the coercive field and exchange bias field of the PL determined from
the MOKE hysteresis loops as a function of the CW laser intensity for four different scan
speeds (from 50 mm·s-1 to 4000 mm·s-1). In terms of the HEB evolution with the laser
intensity, a general trend can be deduced for all speeds investigated and can be described
in three different regimes. In the first regime (e.g. from 67 kW·cm-2 to 190 kW·cm-2 for the
50 mm·s-1 scanning speed), the total energy donated to the layer system is not sufficient
to reach the necessary temperatures, and thus, not all ferromagnetic domains can be
pinned in the direction of the cooling field. The second region resembles a plateau and
refers to a regime where the donated energy is sufficient to set the HEB completely. This
indicates the range of laser parameters where the Néel temperature of the
antiferromagnet is reached at the IrMn / CoFeB interface, and therefore the area-of-
interest for applications. The third region is characterised by a decay of the exchange bias
field strength, which indicates that the additional energy provided to the system is possibly
introducing degradation of the layers due to damage at the interfaces through distinct
diffusion (discussed in chapter 6) and alloying, or ultimately melt (note the colour change
in Figure 46). This is detrimental to the exchange coupling of the CoFeB / IrMn system
and even more the TMR effect of the whole MTJ. It is not possible to estimate the
temperature induced in the MTJ stack by the different laser intensities at different
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scanning speeds currently, but the cooperation partners from LHM have started to
simulate the heat profile in the samples.
Figure 47. The coercive (a.) and exchange bias field (b.) of the PL as a function of the applied laser intensity with scanning speed of 50 mm·s-1, 500 mm·s-1, 2000 mm·s-1, and 4000 mm·s-1 for CW laser annealing. For comparison, the ranges of the maximum values of HC and HEB in the case of oven annealing are shown as hashed areas, including the error bars.
From these investigations, the optimal laser intensity parameter window for setting the
exchange bias field is concluded in Table 4. For the slower scanning velocity (50 mm·s-1)
the optimal window for laser intensities is very limited since a slower scanning velocity
implies that the area-of-interest was exposed to radiation for a longer time interval. This
results in a larger heat transfer to the sample and hence higher temperatures.
The maximum HEB values achieved with CW laser annealing are very close to those
achieved with conventional annealing but yet slightly lower. Since the laser annealing is
performed by scanning a focused CW-laser at 5 μm line spacing, the additional heat of
previous scans and the horizontal heat propagation with a distinct broadening of the heat-
affected zone need to be taken into account. Assuming that the heat dissipation through
the substrate is constant, the local thin film temperature profile may vary drastically with
multiple transitions of the Néel temperature, especially for slower scanning.
Table 4. The optimal laser intensity parameter for the various laser scanning speed.
Scanning speed Optimal laser intensity HEB HC
50 mm·s-1 (200 - 300) kW·cm-2 (97 ± 3) mT (28 ± 2) mT
500 mm·s-1 (250 - 450) kW·cm-2 (103 ± 3) mT (26 ± 3) mT
2000 mm·s-1 (300 - 500) kW·cm-2 (98 ± 4) mT (27 ± 2) mT
4000 mm·s-1 (300 - 500) kW·cm-2 (98 ± 4) mT (26 ± 3) mT
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Pulsed laser annealing
Similar to the CW laser annealing, in this experiment, the sample with the same layer
stack was irradiated with PW laser with varying intensity. In this case, a step distance
between consecutive pulses of 2.5 µm was used, using pulses of 100 ns duration. The
coercive and the exchange bias field strength obtained in the pulsed laser annealed
samples are shown in Figure 48. This plot can also be categorised into three regions,
similar to the results presented in Figure 47. Laser intensities in the range from
500 kW·cm-2 to 900 kW·cm-2 were observed to induce a maximum average exchange
bias field of HEB ≈ (93 ± 5) mT. In this intensity range, the exchange bias field values do
not significantly depend on the laser intensity. In comparison with the CW laser annealed
sample, the exchange bias field values are significantly lower, which is consistent with the
previous studies reposted by Berthold et al.123. This suggests that the PW laser annealing
method has a much higher cooling rate, which in turn results in a lower heat transfer to
the MTJ stack as compared to the CW laser annealing and thus in lower exchange bias
field values.
Figure 48. Coercive field (black symbols, left axis) and exchange bias field strength (red symbols, right axis) as a function of the laser intensity for the samples annealed using PW laser. For comparison, the ranges of the maximum values of HC and HEB in the case of oven annealing (cf. Figure 47) are shown as hashed areas with included error bars.
The coercivity of the pinned layer increases with an increase in intensities up to a value
of HC = (29 ± 4) mT (Figure 48, black symbols.). The HC does not change significantly
(within the error bars) up to a laser intensity of 900 kW·cm-2. In the intensity range from
900 kW·cm-2 to 1750 kW·cm-2, the HC values decrease. Interestingly in comparison to the
CW laser annealing, above 1500 kW·cm-2, the HC increases significantly again. This
increase of the coercivity could be due to alloying effects caused by interdiffusion in the
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magnetic thin films because of the intense heat generated by the short laser pulses. This
is well supported by the observed decrease in exchange bias field strength.
Further investigation of the structural integrity of these samples with XPS depth profiling
showed signs of significant diffusion of Cu from the CuN layer into the uppermost layers
of the stack, even at the laser intensities for which maximum exchange bias field was
observed (Dr. V. Dzhagan, private communication). For the investigations presented in
the following sections, the CuN layer is skipped from the layer stack design.
5.2.3 Influence of the seed layer on the laser annealing
Figure 49 exhibits the exchange bias field strength obtained at different scanning speeds
as a function of the CW laser intensity for the sample with the Ru and Ni81Fe19 seed layer
in the Si / SiO2(100 nm) / Ta(5 nm) / Seed layer / IrMn(8 nm) / Co40Fe40B20(2.3 nm) /
MgO(1.8 nm) / Co40Fe40B20(2.3 nm) / Ta(5 nm) layer stack. When comparing with a
reference value obtained from the same sample but annealed by vacuum oven annealing
at 280°C, it is evident that comparable or even larger exchange bias fields can be
achieved by laser annealing at broadly all scan speeds by choosing the appropriate range
of laser intensities. The evolution in the exchange bias field with the increasing laser
intensity is similar to the results discussed in section 5.2.2 and can be divided into three
different regimes, namely the incremental, plateau and a decremental regime based on
the trends observed in HEB (cf. section 5.2.2).
Figure 49. The exchange bias field strength as a function of the CW laser intensity at different scanning speeds applied on a (0.5 × 0.5) mm2 sample areas for Ru (a.) and Ni81Fe19 (b.) seed layer. The horizontal dashed bar indicates the maximum exchange bias field strength (including the error bars of the exchange bias value) achieved for vacuum oven annealing at 280°C.
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The difference in the obtained maximum exchange bias field strength (average value of
the exchange bias field overall laser intensities comprised in the plateau region) between
samples with a Ru or Ni81Fe19 seed layer are shown in Figure 50.
Figure 50. The difference in the exchange bias field strength determined from the measurements shown in Figure 49. In addition, results of oven annealing at different temperatures are shown in the inset. The samples with a Ru seed layer present a larger exchange bias, both for the oven and laser annealing, with more pronounced differences arising in the latter.
As depicted in the inset, the larger exchange bias field strength for the samples with Ru
seed layer is found for both laser and oven annealed samples, with more pronounced
differences after laser annealing, where a maximum of ΔHEB = 18 mT is found (compared
to ΔHEB ≤ 7 mT). However, for a laser scanning speed of 5000 mm·s-1, no difference in
the pinning field could be observed, which is attributed to difficulties in data evaluation as
the MOKE hysteresis loops for those samples show a pronounced non-linear background.
No other reasonable and physical correct explanation could be found.
The reason for the larger exchange bias field strength with the Ru seed layer is an
improved (111)-texture of the IrMn layer126,127 as proved by XRD (see Figure 51). The
IrMn(111) peak intensity is almost doubled compared to the samples with a Ni81Fe19 seed
layer, which is the case for all laser annealed samples as well (not shown). Note that
Ru(002) also contributes to the detected IrMn(111) peak, as the peak centre of gravity
shifts slightly to larger 2θ values, but the peak at 2θ = 41.36° matches perfectly the
literature value for IrMn(111) [powder diffraction file of IrMn (00-029-0687) from the
International Centre for Diffraction Data (ICDD)] in case of the growth on top of a Ni81Fe19
seed layer. The Ru peak is not expected to contribute significantly to the enhancement
on the IrMn(111) peak intensity, especially taking into account the small Ru layer
thickness. Furthermore, the crystallite size determined from the pronounced Kiessig
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fringes is larger for Ru as given in each figure, although the nominal IrMn layer thickness
is constant for all the samples, supporting the expected epitaxial growth of IrMn on Ru128.
Figure 51. XRD θ–2θ scans of the samples with (a.) Ru and (b.) NiFe seed layer as-deposited and annealed in an oven at 320°C and 400°C. The IrMn(111) orientation (2θ = 41.36°) is represented by the dashed lines, highlighting a small peak shift in (a.) due to the presence of the (002)-textured Ru seed layer (2θ = 42.15°). The given crystallite size was determined from the pronounced Kiessig fringes.
On the other hand, and although the thin film texture does not suffer by annealing at
temperatures up to 400°C, significant changes in the magnetic properties occur for oven
annealing above 320°C. The difference in exchange bias field strength for Ru and NiFe
samples vanishes and the subsequent degradation of the exchange bias (see Figure 50
inset) at these temperatures can be attributed to a degradation of the FM/AFM interface
responsible for the exchange bias effect, as B and even more Mn diffusion (see chapter 6
and 129,130) is likely to occur. Hence, the exchange bias field strength of the studied layer
stacks in general starts to decrease for annealing temperatures larger than about 280°C,
no matter what seed layer material is used.
5.2.4 Influence of the pattern size and number of scan repetitions for laser
annealing
This section investigates systematically the influence of the size of the laser-treated area
on the resulting value of the exchange bias field. For this purpose, the results discussed
in the previous section from the test patterns of (0.5 × 0.5) mm2 on a sample piece are
now compared to the laser annealing on a single sample piece of (6 × 6) mm2 each with
the same layer stacks. Similar to the studies discussed in section 5.2.4, CW laser
annealing was performed at different scanning speeds as well as laser intensities and the
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magnetisation reversal was measured using NanoMOKETM2, concentrating on the PL
reversal. The change in the laser intensity ΔI = I(0.5 × 0.5) mm2 - I(6 × 6) mm
2 required to induce
the same exchange bias field strength in (0.5 × 0.5) mm2 square patterns compared to
the (6 × 6) mm2 sample pieces is plotted in Figure 52 as a function of the laser scanning
speed.
The fact that the same exchange bias field strength is achieved for lower laser intensities
in the large laser-treated patterns is a consequence of the scanning speed-dependent
local heat input in competition with the heat dissipation to the ambience and the substrate.
Whereas for a scan speed of 5 mm·s-1 an equal exchange bias field strength is achieved
for approximately the same laser intensities, the required intensity decreases with
increasing scanning speed in the larger pattern. As the area increases, more energy is
provided to the sample which is accumulated throughout the whole laser annealing
process, thus, increasing the overall sample temperature, especially for a heat input
happening much faster than the heat dissipation at increasing laser scanning speeds.
Figure 52. Change in the laser intensity I = I(0.5 × 0.5) mm
2 - I(6 × 6) mm2 needed to induce the same exchange
bias field strength in (0.5 × 0.5) mm2 square patterns (Figure 55) and (6 × 6) mm2 square patterns as a function of the laser scanning speed.
This is further supported by the results of an experiment in which the laser scanning was
repeated several times on the same trace (see Figure 53). With an increasing number of
scan repetitions at an exemplary scanning speed of 5000 mm·s-1 and fixed laser intensity
of 480 kW·cm-2, a clear deterioration of the exchange bias field strength is observed,
already after five repetitions.
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Figure 53. The exchange bias field strength as the function of the number of CW laser annealing
repetitions at 5000 mm·s-1 scan speed and 240 kW·cm-2 intensity.
For comparison, vacuum oven annealing was performed in SQUID–VSM, where similar
dependencies of exchange bias field strength and the temperature (cf. Figure 49) can be
observed with a maximum exchange bias field strength obtained after field cooling from
250°C (heating and cooling rate of 15 K·s- 1). Higher annealing temperatures lead to a
reduction of the exchange bias (see Figure 54), as it is also the case for higher laser
powers (compare Figure 49), concomitant with deteriorating diffusion processes as
discussed earlier. However, in-field vacuum annealing at 300°C repeated for eight times
does not lead to a relevant change of the exchange bias field strength thereafter, as
shown in Figure 54.
Figure 54. MH hysteresis loops of the sample with layer stack Si / SiO2(100 nm) / Ta(5 nm) / Ru(5 nm) / IrMn(8 nm) / Co40Fe40B20(2.3 nm) / MgO(1.8 nm) / Co40Fe40B20(2.3 nm) / Ta(5 nm) as-deposited (black), after one annealing cycle at 250°C (red), followed by nine cycles of annealing at 300°C (showing only iterations #2 and #9, since no significant changes were found)
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These findings indicate that the degradation of the exchange bias upon repeated
annealing by laser result from intrinsic structural or chemical changes of the IrMn / CoFeB
system, which are not caused by the subsequent reheating, but rather caused by a steady
increase of the sample temperature with each heating cycle. The laser annealing is
accompanied by very large temperature gradients and by time effects due to the scanning
of the laser beam. These effects may induce not only element diffusion and alloying, but
also large stress, probably influencing the magnetic properties of the films. In addition,
note that the laser annealing is performed under ambient conditions, where
contaminations cannot be excluded. However, no reasonable increase in surface
roughness was observed by AFM in any laser annealed sample at those moderate laser
powers with constant exchange bias.
5.3 Application of FORC analysis: from single CoFeB layer to MTJ layer
stack
The in-plane magnetic hysteresis loops of the MTJ stacks measured in the as-deposited
sample and after setting exchange bias (cf. Figure 44.) shows a very small switching field
distribution for the free layer and a gradual switching of the pinned layer due to exchange
biasing. To further understand the switching field behaviour of the individual layers, FORC
analysis was conducted. This sub-chapter summarises the results obtained for a single
CoFeB layer, for an exchange biased CoFeB layer and for the complete MTJ stack, finally.
Single CoFeB layer
Figure 55 shows the magnetisation loops and the FORC distribution of a 100 nm thick
Co60Fe20B20 film passivated with 5 nm Pt in the as-deposited state. This layer mimics the
free layer in MTJ. A single narrow symmetric distribution around the HI axis can be seen
in the FORC plot, indicating the absence of any type of exchange bias or interaction
between the magnetic domains of CoFeB. The FORC distribution is centred at HC = 0 mT.
Such a behaviour is often described as the reversible part of the FORC distibution131 and
ascribed to a superparamagnetic (SP) behaviour131. In this case, it can be explained by
the fact the as-deposited (amorphous) CoFeB layer has very small grains which have a
magnetisation relaxation time similar to or lower than the average measurement time132,
hence resulting in HC→0 mT. The fact that the FORC distribution centred at zero is not
sharp but exhibits a small tail suggests that the film also contains grains which show
typical single domain (SD) switching133. Such grains might consist of CoFe.
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Figure 55. Magnetisation loop (a.) and contour plot presenting the FORC distribution (b.) of the sample with Si / SiO2(1.8 nm) / Co60Fe20B20(100 nm) / Pt(5 nm) in the as-deposited state. Both measurements were performed by SQUID-VSM.
Exchange biased CoFeB
The FORC distribution of the as-deposited Si/SiO2(100 nm) / Ta(3 nm) / Ru(3 nm) /
IrMn(8 nm) / Co40Fe40B20(5 nm) / Ta(3 nm) layer stack is shown in Figure 56 a. The most
obvious influence of the addition of the AFM layer into the stack on the FORC distribution
is the occurrence of the relocation of the maxima on both HC and HI axis. The fact that the
maxima are not symmetric around the HI = 0 is due to random pinning sites134,135 at the
Co40Fe40B20 / IrMn interface. The main contributions to the FORC distribution are located
along +HI, suggesting that most of the pinned domains have preferred magnetisation
reversal direction along positively applied fields. Similar to the free layer FORC
distribution, there is a finite magnetisation at HC = 0 mT (for HI between 0 mT and 20 mT),
which indicates that the amorphous Co40Fe40B20 grains are small but have strong
interaction due to the exchange bias field.
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Figure 56. FORC distribution of a Si / SiO2(100 nm) / Ta(3 nm) / Ru(3 nm) / IrMn(8 nm) / CoFeB(5 nm) / Ta(3 nm) layer stack in the as-deposited state (a.) and after setting the exchange bias by cooling from 200°C in the presence of 120 mT magnetic field(b.).
Figure 56 b shows the FORC diagram of the same layer stack after annealing at 200°C
for 30 min, followed by cooling in the presence of an external magnetic field of 120 mT.
In this case, one pronounced maximum with a well defined FORC contour (HI ≈ 40 mT,
HC ≈ 10 mT) can be observed. This clearly shows the influence of exchange bias in
Co40Fe40B20 films134. A second, smaller contour occurs at HI ≈ 30 mT and HC ≈ 0 T,
indicating that although the applied heat treatment and magnetic field were sufficient to
set the exchange bias, the Co40Fe40B20 film still contains some SP domains. This means
that the Co40Fe40B20 is not completely crystalline after 30 min of annealing at 200°C.
MTJ layer stack
Figure 57 shows the magnetic hysteresis loop and the FORC distribution contour of the
MTJ layer stack consisting of Si / SiO2(100 nm) / Ta(5 nm) / CuN(30 nm) / Ta(3 nm) /
Ni81Fe19(2 nm) / IrMn(8 nm) / Co40Fe40B20(2 nm) / MgO(1.8 nm) / Co40Fe40B20(2.2 nm) /
Ta(10 nm) after oven annealing at 227°C for 30 min and subsequent cooling in the
presence of an applied magnetic field (120 mT). As expected, the FORC countors of the
MTJ stack are the superposition of the free and pinned layers reversal. Due to the large
difference between the switching field distribution of the free and pinned layer, the FORC
distribution contour for each layer appears feeble in the unified FORC plot.
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Figure 57. MH loop (a.) and FORC distribution (b.) recorded at RT for a Si / SiO2(100 nm) / Ta(5 nm) / CuN(30 nm) / Ta(3 nm) / Ni81Fe19(2 nm) / IrMn(8 nm) / Co40Fe40B20(2 nm) / MgO(1.8 nm) / Co40Fe40B20(2.2 nm) / Ta(10 nm)” MTJ layer stack after setting the exchange bias.
The closer look at the individual layer reveals that the FORC distribution of the pinned
layer is centred around HI = -110 mT suggesting the pinning of the CoFeB adjacent to
IrMn with the coercive field distribution concentrated around HC = 10 mT. This FORC
distribution is elongated along the HC direction, indicating a distribution of switching fields
due to the interaction of several exchanged coupled single domains of CoFeB. The FORC
contribution from the free layer is concentrated around HI = 0 mT(at HC ≈ 2 mT) axis,
indicating that the free layer is independent of any coupling due to the exchange-coupled
CoFeB layer. Further, it can be observed that the FORC distribution of the free layer is
not symmetrical with respect to HI = 0 mT. This may be an artefact due to a too large
measurement step compared to the very narrow switching field reversal of the free layer.
5.4 Potential of direct-write laser annealing technique
With the knowledge gained about the laser annealing method, an experiment was devised
to image the magnetic domains in FM/AFM heterostructures on the laser annealed and
unannealed areas. For this, a simpler layer stack consisting of only the pinned layer was
used “Si / SiO2(100 nm) / Ta(3 nm) / Ru(3 nm) / IrMn(8 nm) / CoFeB(5 nm) / Ta(3 nm)”.
The text “HLPH”, abbreviation for “Halbleiterphysik” was written with a CW laser beam of
380 kW·cm-2 intensity at 50 mm·s-1 scanning speed while an applied in-plane magnetic
field of 120 mT under ambient atmosphere. MOKE images were recorded using the
MOKE microscopy at the Leibniz-Institut für Festkörper- und Werkstoffforschung (IFW),
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Dresden by Dr. Ivan Soldatov with a home-built MOKE microscope136. The hysteresis
loops were derived from the grey contrast of two areas-of-interest recorded with the
sweeping magnetic field. An exemplary MOKE microscopy image captured under an
applied magnetic field of H = 25.5 mT, along with the hysteresis loop measured at the
exposed and unexposed region is shown in Figure 58.
It is clearly visible from the hysteresis loop measured at two different regions that the
application of laser annealing in the presence of a magnetic field leads to a shift of the
hysteresis loop toward positive magnetic fields due to the induced exchange anisotropy,
see Figure 58 b. In order to understand the grey contrast in the MOKE images, the
Meiklejohn and Bean model can be used. A magnetic field < -60 mT is applied parallel to
the sample surface (in-plane), the CoFeB layer is saturated in the direction of the applied
magnetic field, resulting in monochromatic grey images. When the magnetic field is swept
towards positive values (opposite to initial direction) through H = 0 mT, the magnetic
domains of CoFeB follow the direction of the applied magnetic field. However, the
domains which are pinned due to exchange bias at the interface between CoFeB and
IrMn hinder this rotation. This results in a different grey contrast for the laser-treated and
untreated regions, which makes the text visible. To rotate the pinned CoFeB
magnetisation in the laser treated regions, a higher magnetic field has to be applied to
overcome the CoFeB / IrMn exchange interaction, a fact which is observed as a shift of
the hysteresis loop along the positive direction of the applied magnetic field. Finally, when
the applied magnetic field reaches the positive values required to induce the saturation
magnetisation, the MOKE microscopy image shows a monochrome grey tone with no
visible text, similar to the situation when the sample was saturated in negative applied
magnetic fields. When the field is swept from the positive saturation field towards negative
magnetic fields, the magnetisation of the pinned CoFeB regions rotates at the smaller
fields, because in this case, the exchange bias at the FM/AFM interface aids the reversal.
Hence, in the full cycle of magnetic field sweep, the text is visible only under a positive
applied magnetic field. On the other hand, in the negative field, the magnetisation for both
pinned and unpinned CoFeB remains saturated; therefore, no text is visible. Here, it is
worth mentioning that the contrast visible in the MOKE micrographs is due to the
difference in the reversal of laser exposed as compared to the unexposed region. The
magnetic domains of CoFeB are very small due to the pinning at the AFM/FM interface
and are beyond the optical resolution limits.
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Figure 58. MOKE microscopy image of the text “HLPH” (“Halbleiterphysik”) inscribed with a CW laser beam on the Si /SiO2(100 nm) / Ta(3 nm) / Ru(3 nm) / IrMn(8 nm) / CoFeB(5 nm) / Ta(3 nm) layer stack. The grey contrasts in the image have the pure magneto-optical origin and correspond to different magnetisation orientations in the FM layer. The hysteresis loop measured on the laser exposed (red) and unexposed region (green); the corresponding area-of-interest is marked in the MOKE image.
5.5 Conclusion
In summary, the results presented in this chapter successfully demonstrate that the laser
annealing can be used for the selective alignment of the magnetisation of the pinned layer
in a typical MTJ. The exchange bias was induced at the AFM/FM interface by annealing
using two methods: large area-annealing in a vacuum and localised heating using laser
irradiation, followed by subsequent cooling in the presence of an external magnetic field.
For this, the temperature on one hand, and the laser intensity, scanning speed, and mode
(PW & CW) on the other hand, were systematically varied to set the maximum exchange
bias. The benchmark values for the exchange bias and the coercive field were obtained
by homogenous heating of the layer stack within the SQUID-VSM magnetometer. The
maximum exchange bias field obtained with CW laser irradiation was marginally lower
than the homogenously heated sample (SQUID-VSM) and the exchange bias set with
pulsed laser annealing was ~18% lower than that achieved with CW laser annealing.
It was shown that in-field laser annealing results in comparable exchange bias field
strengths with larger process windows for increasing scanning speed, being applicable
for different layer stacks. For laser intensities within this process window, no deterioration
of the texture, as well as surface roughness, was observed. Significant changes in HEB
and HC were observed when increasing the laser processed area from (0.5 x 0.5) mm2 to
(6 x 6) mm2. This indicates that there is a delicate dependence of the results of the laser
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annealing process on the heat propagation through the thin film system into the substrate
or heat dissipation to the ambience.
A systematic study of a CoFeB, CoFeB/AFM, and MTJ layer stack using first order
reversal curve measurements present the insight about the irreversible and reversible
magnetisation of a free layer and pinned layer in an MTJ. FORC diagram of the single
CoFeB layer was observed to be concentrated only at one region around HC = 0 mT
(HI = 0mT) indicating that the layer consists of very fine grains of CoFe which show a
superparamagnetic behaviour. For the CoFeB / IrMn layer stack, FORC plots
demonstrate both main distributions due to exchange bias and satellite distribution
associated with CoFeB layer and inhomogeneity at the CoFeB / IrMn interfaces. The
FORC distributions of the MTJ revealed that the coercivities field distribution has two
distinct regions after setting the exchange bias. The first region around HI = 0 mT
produced due to the top CoFeB (FL) whose coercive field distribution is similar to that of
observed for the single CoFeB layer. The distribution further suggests that the FL have
no magnetic interaction to the other constituent layers. The second region is due to the
CoFeB adjacent to the IrMn where the spread of the coercive field is significantly wider
than that of FL and centred at HI = - 110 mT indicating the strong exchange bias at
CoFeB / IrMn interface.
Using MOKE microscopy, the reversal mechanisms of the laser annealed and the
unannealed regions were investigated. Although the magnetic domains of the pinned
CoFeB layer were not visible due to the optical resolution limits, the qualitative information
from the contrast of the recorded images was sufficient to demonstrate the functionality
of laser annealing as a direct-write technique.
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Chapter 6: Exchange bias and diffusion processes in
laser annealed CoFeB/IrMn thin films
In the previous chapter, the results suggested that using Ru as a seed layer improves the
exchange bias due to the improved crystalline texture of IrMn. However, the performance
of an exchange biased system strongly relies on the overall integrity of the constituent
layers in the stack. This chapter presents the influence of laser-induced localised
annealing on the magnetic properties and the diffusion processes occurring in the Si /
SiO2(100 nm) / Ta(3 nm) / Ru(3 nm) / IrMn(8 nm) / CoFeB(5 nm) / Ta(3 nm) layer stack,
using X-ray photoemission spectroscopy depth profiling technique. The results are
compared to those obtained by standard vacuum oven annealing and correlated to the
magnetic properties investigated by magneto-optical Kerr effect magnetometry. The
following chapter was previously published in the Journal of Magnetism and Magnetic
Materials, DOI: 10.1016/j.jmmm.2019.165390.
6.1 Introduction
The large TMR yields rely on thermal treatment to crystalise the CoFeB/MgO layers137,138
together with an appropriate boron migration from the CoFeB103,139. In sensor
applications, this procedure is furthermore crucial to set a reference magnetisation
through the exchange bias effect, appearing as a uniaxial magnetic anisotropy in strongly
coupled AFM/FM thin film systems. Although the migration of boron was reported to be
beneficial for achieving a large TMR ratio, the diffusion mechanisms underlying the
thermal treatment of these devices remain yet controversial. For instance, B migrated
from CoFeB during crystallisation was suggested to be incorporated in the MgO barrier,
providing coherent tunnelling for spin-polarised electrons and hence, improving TMR140–
142. On the contrary, other studies reported that the diffusion of B distorts the tunnel barrier
adversely, resulting in the degradation of TMR143,144. Large TMR ratios were also ascribed
to the migration of B into the Ta capping layer with the formation of Ta-B86,140,145,146.
Furthermore, amorphous CoFeB may work as a suitable diffusion barrier for Mn from the
IrMn antiferromagnet, whereas certain crystalline phases of CoFe facilitate the diffusion
of Mn through vacancies in grain boundaries. Previous studies were so far limited to
conventional oven annealing and rapid thermal annealing processes147.
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Chapter 6
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6.2 Magnetometry investigations
The MOKE magnetometry results obtained from CW laser-treated samples with the layer
stack Si / SiO2(100 nm) / Ta(3 nm) / Ru(3 nm) / IrMn(8 nm) / CoFeB(5 nm) / Ta(3 nm) for
various scanning speeds and irradiation intensities are shown in Figure 59 a. As already
discussed in the previous chapter, the values of the exchange bias field as a function of
the laser intensity can be segmented into three regions for each scanning speed. Despite
the similar trend, the absolute values of HEB and HC differ considerably (by a factor of 2)
due to the different FM layer thicknesses. The choice of the increased CoFeB thickness,
in this case, relates to the requirements for the investigations by XPS depth profiling.
Figure 59 b shows HEB and HC as a function of the annealing temperature for a sample
that was thermally treated in the macro-MOKE chamber.
Figure 59. Exchange bias (red) and coercive field (black) strength determined by MOKE magnetometry as a function of laser intensity and laser scanning speed for Si / SiO2(100 nm) / Ta(3 nm) / Ru(3 nm) / IrMn(8 nm) / CoFeB(5 nm) / Ta(3 nm) sample (a.). The hashed areas in red and black represent the values of exchange bias and coercive field (with the error bars) obtained for a sample annealed in vacuum at 200°C. The coercivity (HC) and exchange bias field (HEB) temperature dependence for the oven annealed sample is shown in (b.). The lines in the figure are guide for the eye.
HEB increases as the annealing temperature increases up to a maximum value of
HEB = (63 ± 2) mT at HC = (11.5 ± 0.5) mT for the sample annealed at 200°C. A further
increase in annealing temperature results in a gradual decrease in HEB. The HC values
compared to the as-deposited sample were observed to decrease with increasing
annealing temperature till 400°C and show an increment once again upon annealing at
500°C. The trend in HEB and HC as a function of temperature is similar to that observed
for laser annealed samples as a function of laser intensity. A maximum of
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HEB = (58 ± 2) mT and HC = (11.4 ± 0.5) mT was recorded for the sample laser annealed
with 380 kW·cm-2 at a scanning speed of 50 mm·s-1.
6.3 XPS-depth profiling
6.3.1 Laser annealed samples
The XPS depth profiling was performed according to the procedure described in section
3.3.5. For the XPS measurements, a distribution profile of all constituent elements is
compared to the nominal layer stack of the as-deposited sample in Figure 60. Even though
oxygen adsorbed at the surface in the form of tantalum oxide is observed in the first three
Ar+ milling steps, its concentration reduces significantly within the Ta layer, inferring that
the layer stack underneath is preserved and not oxidised due to passivating Ta-O. A small
fraction of B can also be observed in the IrMn layer due to the high energy Ar ion sputtering
induced migration of B atoms. In conclusion, all the elements show a well-defined
distribution in accordance with the nominal layer stack, implying that the layers have sharp
interfaces in the as-deposited state.
Figure 60. XPS depth profile of an as-deposited layer Si / SiO2(100 nm) / Ta(3 nm) / Ru(3 nm) / IrMn(8 nm) / CoFeB(5 nm) / Ta(3 nm) sample. The thickness values provided in the layer stack and the upper part of the graph are the nominal values used for the deposition process.
The XPS depth profile of the laser annealed samples at laser intensities of 120 kW·cm-2,
380 kW·cm-2, and 900 kW·cm-2 for all the scanning speeds is shown in Figure 61. To ease
the presentation of results, only Co, Mn, Ru, and O normalised atomic distributions are
plotted. At 120 kW·cm-2 (Figure 61 a) the distribution profiles of all considered elements
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are similar to those of the as-deposited samples, and no change is detectable with the
scanning speed, as expected due to the small induced rise in temperature, being well
below TN. Up to 380 kW·cm-2 laser intensity (Figure 61 b), no diffusion can be identified,
proving that the layers remain intact. The lower HEB after laser annealing compared to
oven annealing, as well as the decrease of HEB with increasing scanning speed shown in
Figure 59 should therefore not be ascribed to significant diffusion of Mn. However, this
cannot be fully clarified with the experimental study performed here, as the depth
resolution of the method does not allow to study changes exactly at the AFM/FM interface
that cause the exchange bias effect. Furthermore, in the particular case of the scanning
speed of 50 mm·s-1, it can be argued that closer study in the range of 250 kW·cm-2 and
380 kW·cm-2 of laser intensities could reveal larger HEB values. At higher laser intensities,
dramatic changes are observed (Figure 61 c).
Figure 61. Distribution profiles of Co, Mn, Ru, and O for the layer stack Si / SiO2(100 nm) / Ta(3 nm) / Ru(3 nm) / IrMn(8 nm) / CoFeB(5 nm) / Ta(3 nm) laser annealed at different laser scanning speeds as indicated in the legend for (a.) 120 kW·cm-2, (b.) 380 kW·cm-2, and (c.) 900 kW·cm-2 laser intensities.
Namely, substantial diffusion of Mn towards the surface for all scanning speeds, as well
as diffusion of Co towards the substrate for 50 mm·s-1 and 500 mm·s-1. For the lowest
laser scanning speed, Ru diffusion occurs too, as well as an increase in average surface
Exchange bias and diffusion processes in laser annealed CoFeB/IrMn thin films
Chapter 6
92
roughness (σavg) from 1.6 nm to 13 nm (RMS surface roughness increases from 0.5 nm
to 1.9 nm) as determined by atomic force microscopy. These results confirm that for laser
intensities ≥900 kW·cm-2, the degradation in HEB and HC is effectively due to the migration
of Mn from the IrMn alloy to CoFeB and finally, intermixing of all the layers. At slower laser
scanning speed, a larger heat load and consequently a larger temperature increase
occurs within the layers, ultimately leading to the diffusion of Mn and Co. The oxygen
distribution profile remains almost unchanged with no considerable diffusion toward the
CoFeB / IrMn / Ru layers for all intensities, thus excluding any oxidation of the layer stack
upon laser irradiation.
Figure 62. Normalised B distribution profile, for all laser annealed samples of the Si / SiO2(100 nm) / Ta(3 nm) / Ru(3 nm) / IrMn(8 nm) / CoFeB(5 nm) / Ta(3 nm) layer stack, indicating a strong migration of B towards the Ta capping layer and the Ru / Ta seed layer.
A closer inspection of the B1s XPS spectra reveals no signature of boron oxide being
present next to elemental boron for laser intensities up to 380 kW·cm-2 for none of the
three scanning speeds. At 900 kW·cm-2 and 50 mm·s-1 scanning speed, a significant peak
of boron oxide at 192.5 eV148 was recorded, see Figure 63 a, which vanishes after a
sputtering time of 100 s (Figure 63 b), meaning that the presence of oxygen in this sample
is limited to the surface. For 500 mm·s-1 a small (two less pronounced peaks at 187 eV
and 192 eV) and for 5000 mm·s-1 no response from boron oxide was observed. Due to
the strong intermixing of all layers and the low intrinsic photoemission cross-section of
B1s core levels, it has to be noted that it was challenging to deconvolute the B oxide peak
and the B metal peak from the background. The lack of B-O peak for 500 mm·s- 1 and
5000 mm·s-1, indicates that the Ta layer works as an effective passivation layer to prevent
Exchange bias and diffusion processes in laser annealed CoFeB/IrMn thin films
Chapter 6
93
the oxygen diffusion in the underlying layer. Hence suggesting vacuum is not mandatory
for the laser annealing.
Figure 63. XPS B spectra recorded after a sputtering time of the 20s for 900 kW·cm-2 laser annealed samples with Si / SiO2(100 nm) / Ta(3 nm) / Ru(3 nm) / IrMn(8 nm) / Co40Fe40B20(5 nm) / Ta(3 nm) layer stack (a.). A clear boron oxide peak was detected at 192.4 eV (dashed line) for the lowest scanning speed only. The B-O peak vanishes after a sputtering time of 100s (b.).
6.3.2 Vacuum oven annealed samples
Pieces of size (6 × 6) mm2 from the same wafer (“Si / SiO2(100 nm) / Ta(3 nm) / Ru(3 nm)
/ IrMn(8 nm) / Co40Fe40B20(5 nm) / Ta(3 nm)”) were vacuum annealed in the macro-
MOKE chamber (P = 10-7 mbar) with an applied magnetic field of 120 mT at 100°C,
200°C, 300°C, 400°C, and 500°C for 30 min. Similar to the laser annealed samples, the
samples annealed in the oven were studied using the XPS depth profile technique, using
identical equipment in the HLPH-TUC group. In this case, binding energy spectra from
0 eV to 1.3 keV were surveyed. For the Ar+ sputtering ion energy of 200 eV was used to
mill (2 × 2) mm2 craters iteratively in 36 cycles with a sputtering time of 90 s. XPS spectra
were then obtained after each sputtering cycle from the centre of the crater with a
monochromatic Al Kα X-ray source and a spot size of 300 µm.
The results for the samples annealed in the presence of H = 120 mT magnetic field at
200°C for 30 min (where maximum HEB is observed see Figure 59 b) are shown in Figure
64 a. Similar to the laser annealed samples, only Co, Mn, and Ru normalised atomic
distributions are plotted. No noticeable diffusion of Mn in Co and Ru for the sample
annealed at 200°C can be observed, indicating the AFM and FM layers remain intact. The
depth profile of B in the layer stack is shown in Figure 64 b. The expected migration of B
Exchange bias and diffusion processes in laser annealed CoFeB/IrMn thin films
Chapter 6
94
towards Ta was observed for temperatures as low as 200°C, whereas its concentration in
the CoFeB layer only starts decreasing significantly above 300°C. Above this
temperature, B was found not only in the Ta cap layer but also in the Ru seed layer,
indicating that the Ta and Ru have a higher affinity to absorb B compared to other
materials in the layer stack.
Figure 64. Depth profile of the atomic concentration of Co, Mn, and Ru calculated from XPS measurements of Si / Si(100 nm) / Ta(3 nm) / Ru(3 nm) / IrMn(8 nm) / CoFeB(5 nm) / Ta(3 nm) after vacuum field annealing at 200°C (a.). The B diffusion profile for as-deposited samples and samples annealed at 200°C, and 500°C (b.).
6.4 Structural analysis
Figure 65 shows the diffraction patterns recorded for the as-deposited state and after CW
laser annealed samples for 380 kW·cm-2 and 900 kW·cm-2 intensities for all three
investigated scanning speeds. The θ-2θ scans (recorded with the Rigaku SmartLab
equipment) show a pronounced IrMn(111) peak around 2θ = 41.5° with a Ru(002)
shoulder peak at higher angles. For none of the measured samples, any crystalline peak
of CoFe was observed, which is consistent with the XPS depth profile study. From the
XPS study it is known that even though the migration of B is visible for all the annealing
parameters, a major concentration of B can still be observed in the CoFeB layer, which
could explain why no CoFe crystalline phase was found in the XRD measurements.
Additionally, as was discussed in the previous section 6.2, the samples irradiated with
900 kW·cm-2 laser intensity at 50 mm·s-1 and 500 mm·s-1 scanning speed showed
intermixing of layers. A similar observation can be made from the XRD scans. The IrMn
peaks are shifted towards higher angles, indicating the disruption of the IrMn crystal
structure.
Exchange bias and diffusion processes in laser annealed CoFeB/IrMn thin films
Chapter 6
95
Figure 65. θ-2θ scans for an as-deposited (yellow) and laser annealed with 380 kW·cm-2 (red), and 900 kW·cm-2 (blue) Si / SiO2(100 nm) / Ta(3 nm) / Ru(3 nm) / IrMn(8 nm) / CoFeB(5 nm) / Ta(3 nm) layer stack for scanning speeds of 50 mm·s-1, 500 mm·s-1, and 5000 mm·s-1. Expected positions for oriented crystalline structures are marked by vertical dashed lines.
6.5 Topographic characterisation
It is often considered that the topography of the layer stack could be used as an indicator
to reflect the FM/AFM interface quality. Hence, the average surface roughness (σavg) can
be used as a qualitative parameter to address the FM/AFM interface roughness. The σavg
as a function of the scanning speed for selected laser annealed samples is shown in
Figure 66. It can be seen that for the investigated scanning speeds, the samples irradiated
with 900 kW·cm-2 have higher roughness as compared to samples exposed to
380 kW·cm-2. For the samples irradiated with 900 kW·cm-2 laser intensity the σavg
decreases with increasing scanning speed. A maximum of σavg = (13 ± 1) nm was
observed for the sample treated at a scanning speed of 50 mm·s-1, where the XPS profile
showed a significant intermixing, i.e. the disruption of the FM/AFM interface. On the other
hand, the samples exhibiting the highest exchange bias (laser-treated at 380 kW·cm-2 and
oven annealed at 200°C, see Figure 59) showed a non-deteriorated surface roughness
σavg comparable to the as-deposited state. It is, thus, possible to conclude that in the
samples with high exchange bias values, the interface between FM/AFM is very smooth.
These results complement the XPS depth profile results very well.
Exchange bias and diffusion processes in laser annealed CoFeB/IrMn thin films
Chapter 6
96
Figure 66. Average surface roughness (σavg) of the Si / SiO2(100 nm) / Ta(3 nm) / Ru(3 nm) / IrMn(8 nm) / CoFeB(5 nm) / Ta(3 nm) layer stack irradiated with 380 kW·cm-2 (red) and 900 kW·cm-2 (yellow) laser intensity for the investigated scanning speeds. The average surface roughness for the 200°C oven annealed samples are shown as dashed blue line. Insets are the 3D topography images of the selected samples (in green: laser annealed samples; in grey: oven annealed samples).
6.6 Conclusion
In this chapter, the effect of different laser annealing parameters on exchange-coupled
FM/AFM layer systems was investigated, using a Si / SiO2 / Ta / Ru / IrMn / CoFeB / Ta
layer stack. The evolution of the magnetic properties determined from MOKE
magnetometry can be well explained by the diffusion process induced by the laser (or
oven) annealing and revealed by XPS depth profiling. The diffusion of B was already
observed below TN for laser as well as oven annealed samples. Increasing laser intensity
and a decreasing laser scanning speed results in the diffusion of Mn towards the surface,
which is the primary reason for the decrease in HEB. At even higher laser powers or slower
scanning speeds, also other materials start to migrate. In comparison to oven annealing,
it was found that laser annealing, despite the shorter time-scale with ultra-fast and
spatially very strong confined temperature gradients, induces comparable diffusion
processes and dependencies, mainly linked to the studied material diffusions (B, Mn, Co,
Ru) and the exchange bias field strength. Finally, and even though the B distribution by
XPS depth profile measurements is difficult to be determined because of its small intrinsic
photoemission cross-section, its profile is a direct proof that exchange bias is independent
on the B migration.
Summary Chapter 7
97
Chapter 7: Summary and outlook
This thesis focuses on investigating the most important prerequisites for an efficient and
sensitive CoFeB based magnetic tunnel junction device, i.e., crystallisation of CoFeB,
setting the exchange bias in FM/AFM layer at the microscopic level using the laser
annealing (direct-write technique), as well as the selective alignment of magnetisation in
the reference layer, otherwise referred as the direct-write technique for the directional
sensitivity.
The crystallisation of CoXFe(80-X)B20 alloys triggered by vacuum oven annealing was
investigated using X-ray diffraction and scanning electron microscopy, as well as optical
and magneto-optical Kerr effect spectroscopy for annealing temperatures ranging from
300°C to 600°C. The transformation of ~100 nm thick CoXFe(80-X)B20 films from amorphous
CoFeB to polycrystalline CoFe was revealed by the sharpening of spectral features
observed in optical and magneto-optical dielectric functions spectra. The influence of B
on the dielectric function was assessed both experimentally and by optical modelling. By
analysing the Drude component of the optical dielectric function, a consistent trend
between the charge carrier scattering time/resistivity and the annealing temperature was
observed, in agreement with the results of electrical investigations by means of the four-
point-probe method. In the case of Co60Fe20B20 thin films, interestingly, more spectral
features were observed in the dielectric spectra after annealing in comparison to the
~100 nm thick films. Though the origin of these additional features in the dielectric spectra
can not be pinpointed, the evolution observed in the spectra with annealing certainly
proves that the (magneto-) optical spectroscopies were very sensitive to the crystallisation
of the investigated layers. Thus this thesis presents a non-destructive, high precision,
swift, and highly sensitive approach to probe the crystallisation of ultra-thin CoXFe(80-X)B20
films based on (magneto-) optical spectroscopy and highlights the advantages of such
spectroscopy techniques for the process and material optimisation in the field of
spintronics. The knowledge gained in this work will be expanded to probe the laser
annealing induced crystallisation of CoFeB in magnetic tunnel junctions, which otherwise
is not possible using the state-of-the-art techniques (XRD) due to the ultrathin films and
multiple layers in the layer stack. It has also promoted the interest in investigating the
dynamic change in the optical properties of the CoFeB during the laser annealing in
collaboration with LHM, Mittweida.
Summary Chapter 7
98
This work gives first insight into the applicability of laser annealing for setting the exchange
bias in magnetic tunnel junctions. It treats some of the challenges concerning the optimal
layer stack and optimal laser annealing parameters for the magnetic tunnel junctions. At
first, the influence of laser intensity, scanning speed, and laser mode (CW & PW) on the
magnetic properties of a typical magnetic tunnel junction layer stack was investigated.
The observed changes in the magnetic properties were compared with similar samples
annealed using a standard vacuum oven, as the benchmark process. It was observed that
with the CW mode laser annealing an exchange bias comparable to the conventional oven
annealing method could be achieved. The optimal laser intensity window for setting the
maximum exchange bias increases with the increase in scanning speed in the
investigated layer stacks with no obvious degradation in structural and magnetic
properties. During the course of this work, the first order reversal curve technique was
applied to investigate the complex magnetisation reversal of a magnetic tunnel junction.
This promising approach will be extended for measuring magnetoresistive first order
reversal curves in collaboration with Fraunhofer ENAS. The combined information from
the magnetisation and magnetoresistance first order reversal curve measurements will
shed new lights on magneto-transport characteristics of the microfabricated magnetic
tunnel junctions.
Finally, the diffusion processes occurring at the interface of CoFeB / IrMn due to laser
annealing were probed using X-ray photoemission spectroscopy depth profiling
technique. These results were compared to those obtained by standard vacuum oven
annealing and correlated to the magnetic properties investigated by magneto-optical Kerr
effect magnetometry. It can be concluded from the depth profiling that the degradation in
exchange bias depends solely on the dealloying AFM material (IrMn) and is independent
of the migration of boron in the layer stack. The (magneto-) optical spectroscopic studies
in combination with the results obtained from depth profiling will help in developing non-
destructive means to probe interlayer diffusion in the microfabricated sensors.
The results of this work have laid the foundation for the future development of the 3D
magnetic field sensor based on the magnetic tunnel junction devices. For such sensors,
the laser annealing has been found to be a promising technique for achieving local and
selective realignment of the magnetisation for directional sensitivity. An important
progress has been made by the combined assessment of the structural, magnetic
Summary Chapter 7
99
properties and by proving the efficacy of optical spectroscopic techniques for optimising
the spintronic device fabrications process.
Appendix
100
Appendix A
Units for magnetic properties
Quantity Symbol Gaussian & cgs
emu
Conversion
factor, C
SI &
rationalised
mks
Magnetic flux density,
magnetic induction B Gauss (G) 10-4 Tesla (T), Wb·m-2
Magnetic flux Φ Maxwell (Mx), G·cm2 10-8 Weber (Wb),
Volt second (V·s)
Magnetic potential difference,
magnetomotive force U, F Gilbert (Gb) 10/4π Ampere (A)
Magnetic field strength,
magnetising force H
Oersted (Oe); Gb·cm-
1 103/4π A·m-1
(Volume) magnetisation g M emu·cm3 h 103 A·m-1
(Volume) magnetisation 4πM G 103/4π A·m-1
Magnetic polarisation,
intensity of magnetisation J, I emu·cm3 4π·10-4 T, Wb·m-2
(Mass) magnetisation σ, M emu·g 1
4π·10-7
A·m2·kg-1
Wb·m·kg-1
Magnetic moment m emu, erg·G 10-3 A·m2, Joule per
Tesla (J·T-1)
Magnetic dipole moment j emu, erg·G 4π·10-10 Wb·m
(Volume) susceptibility Χ, κ dimensionless,
emu·cm3
4π
(4π)2·10-7
Dimensionless
Henry per meter
(H/m), Wb/(A·m)
(Mass) susceptibility Χρ, κρ cm3·g, emu·g 4π·10-3
(4π)2·10-10
m3·kg
H·m2·kg-1
(Molar) susceptibility Χmole,
κmole cm3·mol, emu·mol
4π·10-6
(4π)2·10-13
m3·mol-1
H·m2·mol-1
Permeability μ dimensionless (4π)2·10-7 H·m-1, Wb·(A·m)-1
Relative permeability i μr dimensionless dimensionless
(Volume) energy density,
energy product k W erg·cm3 10-1 J·m-3
Demagnetisation factor D,N dimensionless 1/4π dimensionless
Multiply a number in Gaussian units by “C” to convert it to SI.
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101
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List of figures
111
List of figures
Figure 1. The milestones in the evolution of magnetoresistive devices. ...........................2
Figure 2. The classification of magnetic materials [Image adapted from14]. .....................7
Figure 3. The Stoner model of ferromagnetic metals illustrated for the 3d shell, and
nomenclature used for the band description of magnetism (a.). Occupied electron states
below the Fermi energy EF are shaded, unoccupied states above EF are shown unshaded.
Hext, m and Δ denote the external magnetic field, magnetic moment and exchange
splitting, respectively. The calculated density of states of Fe, Co, Ni, and Cu (b.) [Images
are taken from15]. .............................................................................................................9
Figure 4. The magnetic lattice structure of IrMn superimposed over the crystal lattice (a.).
Magnetic sublattice showing inbound (A-yellow) and outbound magnetisation (B-green)
with zero net magnetisation (b.) [images adapted from24]. ............................................. 11
Figure 5. Exemplary M(H) hysteresis loop of the ferromagnetic material Co60Fe20B20 after
30 min annealing at 600°C: Mr is the remanent magnetisation at H=0; Ms denotes the
saturation magnetisation and HC the coercivity. ............................................................. 12
Figure 6. M(H) hysteresis loop of CoFeB / IrMn bilayer at room temperature (RT). The
green line represents the M(H) response when the sample is cooled from 200°C to RT in
the presence of a magnetic field (120 mT), and the red line is the response of the as-
deposited, demagnetised state. ..................................................................................... 14
Figure 7. Uniaxial magnetic anisotropy induced through the exchange bias effect in
CoFeB / IrMn after cooling from 200°C in the presence of 120 mT magnetic field. The
coercive fields (a.) and the exchange bias fields (b.) are plotted as determined by angle-
resolved longitudinal-MOKE magnetometry at room temperature. ................................. 15
Figure 8. The development road map of the magnetoresistive devices [image was taken
from34]. ........................................................................................................................... 16
Figure 9. Schematic representation of the tunnel magnetoresistance in the case of two
identical ferromagnetic layers separated by a non-magnetic insulating barrier such as
MgO. The tunnelling process conserves the spin. When the electronic states on each side
of the barrier are spin-polarised, the electrons will more easily find free states to tunnel
through the barrier if the magnetisations are parallel (a.) than if they are antiparallel (b.)
to each other due to the ratio of the density of states of both electrons (spin-up, spin-
down) at the Fermi level. The arrows in red and grey show the higher and lower tunnelling
probability of spin-polarised electron through a tunnel barrier, respectively. The yellow
balls represent the electrons with their intrinsic spin orientation and direction of rotation in
grey and blue arrows.[image redrawn from36]. ............................................................... 17
List of figures
112
Figure 10. Various polarisation states of light occurring as a result of the various phase
difference between the two components of the electric field (along x- and y- axes) of equal
amplitude (𝐸𝑦 = 𝐸𝑥) [image taken from37]. ..................................................................... 19
Figure 11. The schematic of spectroscopic ellipsometry [image redrawn from40 ............ 23
Figure 12. The schematic of an optical model for an ambient / thin film / substrate
structure, showing the reflected and refracted light at each interface. Using the Fresnel
coefficients, the contribution of reflections from each interface can be calculated [image
was taken from40]. .......................................................................................................... 24
Figure 13. The three geometrical configurations for MOKE, namely polar (a.), longitudinal
(b.) and transversal (c.). ................................................................................................. 26
Figure 14. The schematic energy state diagram of a 3d-ferromagnet, showing optical
transitions induced by left (blue) and right (red) polarised photons for a system where only
spin-orbit coupling is present (left diagram) and for a system where spin-orbit coupling
and exchange interaction is present (middle diagram). The notation in the | ⟩ brackets
contains the orbital number (𝑙), magnetic number (𝑚), and spin orientation (↑ or ↓). The
right side diagram shows the corresponding absorption spectra of left and right circular
polarised light. [adapted from44] ..................................................................................... 28
Figure 15. Schematic representation of the magnetron sputtering. The argon ions (in red)
are responsible for the target etching. The ejected particles (in grey) are sputtered towards
the substrate. The direction of the magnetic field used for confining the electrons and the
ions close to the target is illustrated as blue arrows. ...................................................... 32
Figure 16. Oven sample mounting platform with grading to place the sample on oven
heater stick [image was taken from48]. ........................................................................... 33
Figure 17. CAD modelled isometric view (a.) and transverse view (b.) of the heater
assembly developed for the macro-MOKE system......................................................... 34
Figure 18. Schematics of the experimental setup for Nd: YAG laser [image prepared for
the own publication49]. .................................................................................................... 35
Figure 19. Drawing of the diffractometer showing the general scheme of various
goniometers and measurement axes [image taken from52]. ........................................... 36
Figure 20. A basic principle of atomic force microscopy [image was taken from54]. ....... 38
Figure 21. The spectroscopic ellipsometry setup M2000 for measuring the complex ratio
(𝜌) of the Fresnel reflection coefficients. ........................................................................ 39
Figure 22. Schematic diagram of the magneto-optical Kerr effect spectrometer used in
this thesis for measuring Kerr rotation (θK) and ellipticity (ηK). ....................................... 40
List of figures
113
Figure 23. A cross-sectional view of the utilised MPMS SQUID - VSM MPMS-3 setup
[image taken from60]. ...................................................................................................... 41
Figure 24. Top view of the Nano-MOKETM2 magnetometer. The yellow arrows point out
the optical components in the optical path of the laser shown with the red dashed line. 42
Figure 25. An exemplary first order reversal curve (FORC) recorded for Si /
SiO2(100 nm) / Ni81Fe19(20 nm). The reversal magnetic field (Ha) and regular magnetic
field (Hb) are shown in red and black solid dots, respectively (a.). The major hysteresis
loop (MHL) is shown in the red dashed line. (b.) The FORC distribution calculated from
the measured FORC loops. The contour denotes the maxima of the distribution at the
“irreversible” located at about HC = 2.2 mT, HI = 0 mT. .................................................. 44
Figure 26. Process flow diagram of XPS-depth profiling. ............................................... 45
Figure 27. X-ray diffraction patterns recorded for Si / SiO2(1.8 nm) /
CoFeB(100 nm) / Pt(5 nm) before and after annealing under UHV at the indicated
temperature for Co40Fe40B20 (a.) and Co60Fe20B20(b.). Additionally, the scan of the as-
deposited Si / SiO2(1.8 nm) / CoFe(100 nm) /Pt(5 nm) sample is presented in black in
figure 27(a.).The respective reflexes of constituent materials are marked by dotted lines
along with the miller indices [powder diffraction file of CoFe (00-049-1567), Pt (00-004-
0802), Si (00-027-1402), and Ag (00-004-0783) from the International Centre for
Diffraction Data (ICDD)]. ................................................................................................ 49
Figure 28. Rocking curve measured at the CoFe(110) reflex for the Co40Fe40B20 (a.) and
Co60Fe20B20 (b.) films. The inset in (b.) shows the FWHM of particular Gaussian fits of the
obtained peaks. .............................................................................................................. 50
Figure 29. The CoFe crystallite sizes (vertical coherence lengths) calculated using the
Scherrer equation for Co40Fe40B20 and Co60Fe20B20 (a) determined from XRD shown in
Figure 27. SEM micrograph collage of Si / SiO2(1.8 nm) / Co40Fe40B20(100 nm) / Pt(5 nm)
before (above) and after annealing (below) recorded in the FIB trench at 36° stage tilt (b.).
....................................................................................................................................... 51
Figure 30. The grazing incidence diffractogram recorded for CoFeB thick films before and
after annealing at various temperature for two stoichiometries: Co40Fe40B20 (a.) and
Co60Fe20B20 (b.). The respective reflexes of the constituent materials are marked by dotted
lines along with the Miller indices. .................................................................................. 52
Figure 31. The evolution of and (inset) spectra recorded for the Co40Fe40B20 (a.) and
Co60Fe20B20 (b.) thick films before and after annealing at various temperatures. For
comparison, the and spectrum of Co50Fe50 is plotted along with the Co40Fe40B20
spectra in (a.). ................................................................................................................ 53
List of figures
114
Figure 32. The complex dielectric function (ε1xx & ε2xx) spectra of the Co50Fe50 (red),
Co40Fe40B20 (blue) and B (grey) 103, together with the simulated ε1xx & ε2xx of (Co50Fe50)+B
with 15 % B inclusion(yellow). More detailed information about the B inclusion is given in
the text. .......................................................................................................................... 54
Figure 33. The annealing temperature dependent evolution of ε1xx (a.) and ε2xx (b.) spectra
for Co40Fe40B20 (solid line) and Co60Fe20B20 (dashed line), and CoFe (black). ............... 55
Figure 34. Drude parameters resistivity (ρ) in black and scattering time (τs) in red colour
as a function of annealing temperature for Co40Fe40B20 (solid symbol) and Co60Fe20B20
(empty symbol). The lines in the figure are drawn to guide the eye (a.). Sheet resistance
of the Co40Fe40B20 (filled circles in red) and Co60Fe20B20 (unfilled circles in blue) layers
passivated with a Pt thin film as a function of the annealing temperature (b.). ............... 57
Figure 35. Measured polar Kerr effect rotation (θK) and ellipticity (ηK) spectra of Co50Fe50
in comparison to the previously published reference107,108 (a.) and calculated (h)2×xy as
the function of photon energy (b.). ................................................................................. 58
Figure 36. The polar Kerr effect measured rotation (θK) and ellipticity (ηK) spectra for the
100 nm thick CoFeB films before and after annealing at the indicated temperatures for
two stoichiometry Co40Fe40B20 (a. & c.) and Co60Fe20B20 (b. & d.). ................................ 59
Figure 37. The calculated xy(h)2 as a function of the photon energy for two
stoichiometries, Co40Fe40B20 and Co60Fe20B20, for Si / SiO2(1.8 nm) / CoXFe(80-
X)B20(100 nm) / Pt(5 nm) layer stacks annealed at 450°C and 600°C. ........................... 60
Figure 38. X-ray diffractogram recorded for Si / SiO2(100 nm) / Co60Fe20B20(20 nm) /
Au(3 nm), in the as-deposited state (squares) and after annealing at 350°C (circles) and
400°C (triangles). ........................................................................................................... 61
Figure 39. Experimental SE ψ and ∆ spectra recorded at 65° AOI for Si / SiO2(100 nm) /
Co60Fe20B20(20 nm) / Au(3 nm) before and after vacuum oven annealing at 350°C and
400°C. ............................................................................................................................ 62
Figure 40.The ε1xx (black) and ε2xx (red) components of the complex of the dielectric
function of Co60Fe20B20 (10 nm to 20 nm) before (continuous line) and after annealing at
350°C (dashed line) and 400°C (dotted line)(a.). Comparison of the complex dielectric
function of as-deposited 20 nm and 100 nm thick Co60Fe20B20 films (b.). ...................... 63
Figure 41. Comparison of Kerr rotation θK (black) and ellipticity ηK (red) for Si /
SiO2(100 nm) / Co60Fe20B20(20 nm) / Au(3 nm) before (continuous line) and after
annealing at 350°C (dashed line) and 400°C (dotted line). ............................................ 65
List of figures
115
Figure 42. A schematic diagram showing a typical layer stack of the MTJ sensor and the
laser-induced temperature profile for setting the exchange bias in the FM/AFM layers. The
CoFeB (FM) layer next to IrMn (AFM) is the pinned layer and the upper CoFeB layer is
the free layer, both sandwich the MgO tunnel barrier. The Ni81Fe19 and CuN serve as the
seed and buffer layer, respectively. The Ta layers work as the passivation (on the top)
and adhesion layer (at the bottom). ................................................................................ 69
Figure 43. Magnetisation as a function of the in-plane applied external magnetic field on
an as-grown sample (a.) and measured along the easy-axis of the bottom pinned layer
after 120 mT field cooling from 227°C (b.). The arrow in green and red represent the
magnetisation of the free layer and pinned layer, respectively. ...................................... 70
Figure 44. SQUID-VSM measured M(H) hysteresis loops recorded at RT for an as-
deposited sample and after consecutive field cooling process from different temperatures
as indicated in the legend. .............................................................................................. 71
Figure 45. Temperature dependence of the coercivity HC (a.)and the exchange bias field
HEB for the determination of the TB. The sample was initially cooled in 120 mT field in-
plane applied magnetic from 327°C to RT before the measurements at the increasing
temperature. ................................................................................................................... 72
Figure 46. Optical micrograph of an exemplary sample annealed with CW irradiation with
different scanning speed, 50 mm·s-1 (a.), 500 mm·s-1 (b.), 2000 mm·s-1 (c.), and
4000 mm·s-1 (d.). In each square, the different laser power was used as indicated, the
patterns, as well as the labels, were created by laser ablation. ...................................... 73
Figure 47. The coercive (a.) and exchange bias field (b.) of the PL as a function of the
applied laser intensity with scanning speed of 50 mm·s-1, 500 mm·s-1, 2000 mm·s-1, and
4000 mm·s-1 for CW laser annealing. For comparison, the ranges of the maximum values
of HC and HEB in the case of oven annealing are shown as hashed areas, including the
error bars. ....................................................................................................................... 74
Figure 48. Coercive field (black symbols, left axis) and exchange bias field strength (red
symbols, right axis) as a function of the laser intensity for the samples annealed using PW
laser. For comparison, the ranges of the maximum values of HC and HEB in the case of
oven annealing (cf.Figure 47) are shown as hashed areas with included error bars. ..... 75
Figure 49. The exchange bias field strength as a function of the CW laser intensity at
different scanning speeds applied on a (0.5 × 0.5) mm2 sample areas for Ru (a.) and
Ni81Fe19 (b.) seed layer. The horizontal dashed bar indicates the maximum exchange bias
field strength (including the error bars of the exchange bias value) achieved for vacuum
oven annealing at 280°C. ............................................................................................... 76
List of figures
116
Figure 50. The difference in the exchange bias field strength determined from the
measurements shown in Figure 49. In addition, results of oven annealing at different
temperatures are shown in the inset. The samples with a Ru seed layer present a larger
exchange bias, both for the oven and laser annealing, with more pronounced differences
arising in the latter. ......................................................................................................... 77
Figure 51. XRD θ–2θ scans of the samples with (a.) Ru and (b.) NiFe seed layer as-
deposited and annealed in an oven at 320°C and 400°C. The IrMn(111) orientation (2θ =
41.36°) is represented by the dashed lines, highlighting a small peak shift in (a.) due to
the presence of the (002)-textured Ru seed layer (2θ = 42.15°). The given crystallite size
was determined from the pronounced Kiessig fringes. ................................................... 78
Figure 52. Change in the laser intensity I = I(0.5 × 0.5) mm2 - I(6 × 6) mm
2 needed to induce the
same exchange bias field strength in (0.5 × 0.5) mm2 square patterns (Figure 55) and
(6 × 6) mm2 square patterns as a function of the laser scanning speed. ........................ 79
Figure 53. The exchange bias field strength as the function of the number of CW laser
annealing repetitions at 5000 mm·s-1 scan speed and 240 kW·cm-2 intensity. ............... 80
Figure 54. MH hysteresis loops of the sample with layer stack Si / SiO2(100 nm) / Ta(5 nm)
/ Ru(5 nm) / IrMn(8 nm) / Co40Fe40B20(2.3 nm) / MgO(1.8 nm) / Co40Fe40B20(2.3 nm) /
Ta(5 nm) as-deposited (black), after one annealing cycle at 250°C (red), followed by nine
cycles of annealing at 300°C (showing only iterations #2 and #9, since no significant
changes were found) ...................................................................................................... 80
Figure 55. Magnetisation loop (a.) and contour plot presenting the FORC distribution (b.)
of the sample with Si / SiO2(1.8 nm) / Co60Fe20B20(100 nm) / Pt(5 nm) in the as-deposited
state. Both measurements were performed by SQUID-VSM. ........................................ 82
Figure 56. FORC distribution of a Si / SiO2(100 nm) / Ta(3 nm) / Ru(3 nm) / IrMn(8 nm) /
CoFeB(5 nm) / Ta(3 nm) layer stack in the as-deposited state (a.) and after setting the
exchange bias by cooling from 200°C in the presence of 120 mT magnetic field(b.). .... 83
Figure 57. MH loop (a.) and FORC distribution (b.) recorded at RT for a Si /
SiO2(100 nm) / Ta(5 nm) / CuN(30 nm) / Ta(3 nm) / Ni81Fe19(2 nm) / IrMn(8 nm) /
Co40Fe40B20(2 nm) / MgO(1.8 nm) / Co40Fe40B20(2.2 nm) / Ta(10 nm)” MTJ layer stack
after setting the exchange bias. ..................................................................................... 84
Figure 58. MOKE microscopy image of the text “HLPH” (“Halbleiterphysik”) inscribed with
a CW laser beam on the Si /SiO2(100 nm) / Ta(3 nm) / Ru(3 nm) / IrMn(8 nm) /
CoFeB(5 nm) / Ta(3 nm) layer stack. The grey contrasts in the image have the pure
magneto-optical origin and correspond to different magnetisation orientations in the FM
layer. The hysteresis loop measured on the laser exposed (red) and unexposed region
(green); the corresponding area-of-interest is marked in the MOKE image. .................. 86
List of figures
117
Figure 59. Exchange bias (red) and coercive field (black) strength determined by MOKE
magnetometry as a function of laser intensity and laser scanning speed for
Si / SiO2(100 nm) / Ta(3 nm) / Ru(3 nm) / IrMn(8 nm) / CoFeB(5 nm) / Ta(3 nm) sample
(a.). The hashed areas in red and black represent the values of exchange bias and
coercive field (with the error bars) obtained for a sample annealed in vacuum at 200°C.
The coercivity (HC) and exchange bias field (HEB) temperature dependence for the oven
annealed sample is shown in (b.). .................................................................................. 89
Figure 60. XPS depth profile of an as-deposited layer Si / SiO2(100 nm) / Ta(3 nm) /
Ru(3 nm) / IrMn(8 nm) / CoFeB(5 nm) / Ta(3 nm) sample. The thickness values provided
in the layer stack and the upper part of the graph are the nominal values used for the
deposition process. ........................................................................................................ 90
Figure 61. Distribution profiles of Co, Mn, Ru, and O for the layer stack Si / SiO2(100 nm) /
Ta(3 nm) / Ru(3 nm) / IrMn(8 nm) / CoFeB(5 nm) / Ta(3 nm) laser annealed at different
laser scanning speeds as indicated in the legend for (a.) 120 kW·cm-2, (b.) 380 kW·cm-2,
and (c) 900 kW·cm-2 laser intensities. ............................................................................. 91
Figure 62. Normalised B distribution profile, for all laser annealed samples of the Si /
SiO2(100 nm) / Ta(3 nm) / Ru(3 nm) / IrMn(8 nm) / CoFeB(5 nm) / Ta(3 nm) layer stack,
indicating a strong migration of B towards the Ta capping layer and the Ru / Ta seed layer.
....................................................................................................................................... 92
Figure 63. XPS B spectra recorded after a sputtering time of the 20s for 900 kW·cm-2 laser
annealed samples with Si / SiO2(100 nm) / Ta(3 nm) / Ru(3 nm) / IrMn(8 nm) /
Co40Fe40B20(5 nm) / Ta(3 nm) layer stack (a.). A clear boron oxide peak was detected at
192.4 eV (dashed line) for the lowest scanning speed only. The B-O peak vanishes after
a sputtering time of 100s (b.). ......................................................................................... 93
Figure 64. Depth profile of the atomic concentration of Co, Mn, and Ru calculated from
XPS measurements of Si / Si(100 nm) / Ta(3 nm) / Ru(3 nm) / IrMn(8 nm) / CoFeB(5 nm)
/ Ta(3 nm) after vacuum field annealing at 200°C (a.). The B diffusion profile for as-
deposited samples and samples annealed at 200°C, and 500°C (b.). ........................... 94
Figure 65. θ-2θ scans for an as-deposited (yellow) and laser annealed with 380 kW·cm-2
(red), and 900 kW·cm-2 (blue) Si / SiO2(100 nm) / Ta(3 nm) / Ru(3 nm) / IrMn(8 nm) /
CoFeB(5 nm) / Ta(3 nm) layer stack for scanning speeds of 50 mm·s-1, 500 mm·s-1, and
5000 mm·s-1. Expected positions for oriented crystalline structures are marked by vertical
dashed lines. .................................................................................................................. 95
Figure 66. Average surface roughness (σavg) of the Si / SiO2(100 nm) / Ta(3 nm) /
Ru(3 nm) / IrMn(8 nm) / CoFeB(5 nm) / Ta(3 nm) layer stack irradiated with 380 kW·cm-2
(red) and 900 kW·cm-2 (yellow) laser intensity for the investigated scanning speeds. The
average surface roughness for the 200°C oven annealed samples are shown as dashed
blue line. Insets are the 3D topography images of the selected samples (in green: laser
annealed samples; in grey: oven annealed samples)..................................................... 96
List of tables
118
List of tables
Table 1. The selection rules of (magneto-) optical transitions ........................................22
Table 2. Nominal and XRR determined thicknesses of the CoXFe(80-X)B20, Co50Fe50, and
Pt layer for as the as-deposited samples .......................................................................48
Table 3. Nominal and XRR determined thicknesses of the Co60Fe20B20 and Au layer for
as the as-deposited samples ..........................................................................................61
Table 4. The optimal laser intensity parameter for the various laser scanning speed. ...74
Abbreviations
119
Abbreviations
AFM Antiferromagnet
FL Free Layer
FM Ferromagnetic
GMR Giant Magnetoresistance
HV High Vacuum (0.1 𝑚𝑃𝑎)
IR Infrared
LASER laser
Light Amplification by Stimulated Emission of Radiation
MTJ Magnetic Tunnel Junction
MO Magneto-optical
MOKE Magneto-optical Kerr Effect
MR Magnetoresistance
MTJ Magnetic tunnel junction
NM Non-magnetic material
PL Pinned Layer
RT Room Temperature
SE Spectroscopic Ellipsometry
SEM Scanning Electron Microscope
SQUID-VSM Superconducting Quantum Interference Device – Vibrating Sample Magnetometry
TEM Transmission Electron Microscope
TMR Tunnel Magnetoresistance
UHV Ultra-High Vacuum (0.1 𝑛𝑃𝑎)
UV Ultraviolet
XPS X-ray Photon Spectroscopy
XRD X-ray Diffraction
XRR X-ray Refractometry
Erklärung
120
Erklärung
Ich erkläre, dass ich die vorliegende Arbeit selbständig und nur unter Verwendung der
angegebenen Literatur und Hilfsmittel angefertigt habe.
Januar, 2020
Apoorva Sharma
curriculum vitae
121
curriculum vitae
Apoorva Sharma
Date of birth 02 May 1987
Place of birth Bikaner, India
Nationality Indian
Gender Male
Marital status Married
Foreign language English, German
Academic
2005-2009
Bachelor’s Electronics and Communication Engineering, University of Rajasthan, Jaipur India.
“Microcontroller based automation system”.
First-class
honours
2010-2012
CSIR-Central Electronics Engineering Research Institute, Pilani-333031, India.
“Magnetic Nanocomposite materials and their application in V-Groove inductor”.
2013-2016
Master of Science at Faculty of Electrical Engineering and Information Technology, University of Technology Chemnitz, Chemnitz Germany.
“Magnetic Characterisation of [Co/Ni] multilayers for magnetoresistive applications”.
Sehr gut
Scientific contributions
122
Scientific contributions
List of publications
A. Sharma, M.A. Hoffmann, P. Matthes, S. Busse, O. Selyshchev, P. Mack, H. Exner, A. Horn, S.E. Schulz, D.R.T. Zahn, and G. Salvan.
Exchange bias and diffusion processes in laser annealed CoFeB/IrMn thin films.
J. Magn. Magn. Mater. 489, 165390 (2019).
A. Sharma, M.A. Hoffmann, P. Matthes, N. Kohler, S. Busse, M. Muller, H. Exner, S.E. Schulz, D.R.T. Zahn, and G. Salvan.
Magnetic tunnel junctions: Laser annealing versus oven annealing.
IEEE Trans. Magn. 55, 1 (2019).
A. Sharma, M.A. Hoffmann, P. Matthes, O. Hellwig, C. Kowol, D. R.T Zahn, S. E. Schulz, and G. Salvan.
Crystallisation of optically thick films of CoxFe(80-x)B20: evolution of the (magneto-) optical and structural properties
Phys. Rev. B 101, 054438 (2020)
A. Sharma, P. Matthes, I. Soldatov, S. S. P. K. Arekapudi, B Böhm, M. Lindner, O. Selyshchev, N.T.N. Ha, M. Mehring, C. Tegenkamp, S. E Schulz, Dietrich R T Zahn, Y. Paltiel, O. Hellwig and G. Salvan.
Control of magneto-optical properties of cobalt-layers by adsorption of a-helical polyalanine self-assembled monolayers.
J. Mater. Chem. C, 2050-7526 (2020).
M.A. Hoffmann, A. Sharma, P. Matthes, S. Okano, O. Hellwig, R. Ecke, D.R.T. Zahn, G. Salvan, and S.E. Schulz.
Spectroscopic ellipsometry and magneto-optical Kerr effect spectroscopy study of thermally treated Co60Fe20B20 thin films.
J. Phys.: Condens. Matter (2019) (in press).
Scientific contributions
123
S. Nikam, A. Sharma, M. Rahman, A.Teli, S. Mujawar, D.R.T. Zahn, G. Salvan, P. S. Patil, S.C. Sahoo, and P.B. Patil.
Pulsed Laser Deposited CoFe2O4 Supercapacitor Electrode
RSC Adv., 10, 19353-19359, (2020).
N.T.N. Ha, A. Sharma, D. Slawig, S. Yochelis, Y. Paltiel, D.R.T. Zahn, G. Salvan, and C. Tegenkamp.
Charge-ordered -helical polypeptide monolayers on Au(111)
J. Phys. Chem. C 124, 10, 5734–5739, (2020).
S. Okano, A. Sharma, A. Nishimura, C. Günther, O. D. Gordan, K. Ikushima, G. Salvan, V. Dzhagan and D.R.T. Zahn.
Voltage controlled dielectric function of bilayer graphene.
Adv. Optical Mater., 8, 2000861, (2020).
C. Saengruengrita, A. Sharma, D. Solonenkob, P. Thamyongkita, S. Wacharasindhua, S. Sattayapornc, G. Salvanb, D. R. T. Zahn, and N. Insin.
Iron oxide nanospheres and nanocubes modified with carboxyphenyl porphyrin and their magnetic, optical properties and photocatalytic activities in room temperature amide synthesis.
J. Magn. Magn. Mater. (in press)
C. Saengruengrit, P. Ritprajak, S. Wanichwecharungruang, A. Sharma, G. Salvan, D.R.T. Zahn, and N. Insin.
The combined magnetic field and iron oxide-PLGA composite particles: Effective protein antigen delivery and immune stimulation in dendritic cells J. Colloid Interface Sci. 520, 101 (2018).
J.D. John, S. Okano, A. Sharma, O. Selyshchev, M. Rahaman, N. Miyachi, K. Enomoto, J. Ochiai, I. Saito, T. Masuzawa, T. Yamada, D.H.C. Chua, D.R.T. Zahn, and K. Okano
Transport properties of Se/As2Se3 nanolayer superlattice fabricated using rotational evaporation
Adv. Funct. Mater. 29, 1904758 (2019).
Scientific contributions
124
R.P. Dhavale, P.P. Waifalkar, A. Sharma, R.P. Dhavale, S.C. Sahoo,
P. Kollu, A.D. Chougale, D.R.T. Zahn, G. Salvan, P.S. Patil, and
P.B. Patil.
Monolayer grafting of aminosilane on magnetic nanoparticles: An efficient approach for targeted drug delivery system
J. Colloid Interface Sci. 529, 415 (2018).
J. D. John, S. Okano, A. Sharma, M. Rahaman, O. Selyshchev, N. Miyachi, K. Enomoto, J. Ochiai, I. Saito, G. Salvan, T. Masuzawa, T. Yamada, D. H. C. Chua, D. R. T. Zahn, and K. Okano.
Observation of two-level defect system in amorphous Se superlattices.
Appl. Phys. Lett., 116, 19, (2020).
M. Rahaman, O. Selyshchev, Y. Pan, I. Milekhin, A. Sharma, G. Salvan, S. Gemming, and D. R. T. Zahn.
Radiative decay of dark exciton related emission in a sandwiched monolayer WSe2 revealed by room temperature micro and nano photoluminescence.
Submitted.
R. Patra, H. Stöcker, A. Sharma, M. Monecke, G. Salvan, R. Mattheis,
S. Pofahl, R. Schäfer, O.G. Schmidt, and H. Schmidt.
Magneto-optical response of multilayer structures with ferromagnetic NiFe, CoFe, or CoFeB thin films.
Submitted.
V.C. Karade, A. Sharma, R.P. Dhavale, R.P. Dhavale, P.S. Patil, J. H. Kim, D.R.T. Zahn, A.D. Chougale, G. Salvan, and P.B. Patil.
APTES monolayer grafting on magnetic nanospheres for controlled release of anticancer drug Nintedanib.
Submitted.
Scientific contributions
125
List of conferences and workshops
Magnetic Properties of Porphyrin Magnetite Nanocomposites
March 2017, DPG-Frühjahrstagung, Dresden, Germany (Poster).
Magneto-optical Spectroscopy and Spectroscopic Ellipsometry of Co60Fe20B20
April 2017, IEEE Intermag 2017, Dublin, Ireland (Poster).
Imaging Ellipsometer Nanofilm-EP4
June 2017, Accurion summer school 2017, Göttingen, Germany.
Condensed Matter Magnetism : Bulk Meets Nano
October 2017, European School on Magnetism, Halbleiterphysik, Cargèse, France.
Spintronics (won the third prize)
November 2017, Science Slam, Zentrum für den wissenschaftlichen Nachwuchs-TU
Chemnitz, Chemnitz, Germany. (Talk).
Magneto-optical Spectroscopy and Spectroscopic Ellipsometry of Co60Fe20B20
Thin Films
March 2018, 10th Workshop Ellipsometry (WSE 10), Chemnitz, Germany (Talk).
Novel Method of Setting Exchange Bias in Tunnel Magnetoresistance Devices
with Laser Annealing
March 2018, DPG-Frühjahrstagung, Berlin, Germany (Poster).
Spectroscopic Ellipsometry of Diffusion in Magnetic Multilayer Stacks
September 2018, Joint European Magnetic Symposia (JEMS), Mainz, Germany
(Poster).
Magnetite Nanoparticles and Their Biomedical Applications (won first prize)
April 2019, Nanobio19, online poster presentation, (Poster).
126
Acknowledgements
This research work is the result of teamwork, and nothing would have been possible
without the collaboration and the patience of many people. In this context, I thank all the
project partners involved in this work for a very fruitful scientific collaboration.
I would like to thank Prof. Dr. Georgeta Salvan and Prof. Dr. Dr. Dietrich R.T. Zahn for
giving me this opportunity to take up the challenges of this project. I further extend my
gratitude to Prof. Dr. Georgeta Salvan for being a constant source of motivation, for her
patience and for her deep human understanding. She is simply the best forewoman to
work under.
All my thanks to my supervisors Dr. Patrick Matthes and Ms Maria Almeida, for their
continuous support. In particular, I want to thank Dr. Patrick Matthes for the never-ending
discussions in which I always discovered that physics is just simple solutions for difficult
problems. I am really grateful to Ing. Axel Fechner (“The engineer that can fix anything
except broken hearts”) for his technical support and coffee corner discussions. My special
thanks belong to Mrs Sybille Raschke and Mrs Jane Eisentraut, who always took care of
all my administrative work and really pushed me to speak German with them. I want to
thank my fellow group members for the support for providing a very enjoyable scientific
ambience. During my PhD, I met some wonderful people at various conferences who
either directly or indirectly contributed to this thesis. To this, a special thanks to Dr. Ivan
Soldatov and Dr. Prashant Patil for efficient and productive collaborations. I also thank Mr
Fabian Ganss for maintaining the SQUID-VSM facility at its best throughout this work. I
am grateful to Dr. Volodymyr Dzhagan and Dr. Paul Mack for the XPS depth profile
measurements and the valuable discussion for data analysis.
Finally, no thanks will ever be enough for my family, my mother Mrs Rashmi Sharma, my
father Mr Mahesh Kumar Sharma,and my sister Mrs Pallvi Sharma because their trust
and pride in me never rest. My friends Himani, Shun, Ilya, Sasha (Oleksandr) and Josh,
who never fails to remind me that there is more to life than physics alone.
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