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Chapter II
Literature Review
CHAPTER- II
LITERATURE REVIEW
2.1 Magnetism of materials
The origin of magnetism of materials comes from the orbital and spin motions of
electrons, in which the electrons interact with one another. The best way to understand
different types of magnetism is to describe how materials respond to magnetic fields. The
main different is that the some materials there is no collective interaction of atomic
magnetic moments, while in other materials there is a very strong interaction between
atomic moments [7]. Magnetism arises due to the moving charges, such as an electric
current in a coil of wire. Even when there is no current present in a material, still there are
magnetic interactions. The processes which create magnetic field in an atom are:
i) Nuclear spin, some nuclei such as in hydrogen atom, have a net spin which creates a
magnetic field. ii) Electron spin, an electron has two types of spin states, which are
called spin up and spin down. iii) Electron orbital motion. There is a magnetic field due
to the orbital motion of the electron. Each of these magnetic fields interacts with one
another and with external magnetic fields. Some of these interactions are strong; others
are negligible weak. In most case, magnetic interaction with nuclear spin is a very minor
effect. But interactions between the intrinsic spin of one electron and the intrinsic spin of
another electron are strongest for very heavy elements such as the actinides. This is called
spin-spin coupling. For these elements this coupling can shift the electron orbital energy
levels. The interaction between on electron’s intrinsic spin and its orbital motion is called
spin-orbit coupling has a significant effect on the energy levels of the orbital in many
inorganic compounds.
2.2 Magnetic Hysteresis
The lag or delay of a magnetic material known commonly as Magnetic
Hysteresis, relates to the magnetisation properties of a material by which it firstly
becomes magnetised and then de-magnetised. We know that the magnetic flux generated
by an electromagnetic coil is the amount of magnetic field or lines of force produced
within a given area and that it is more commonly called "Flux Density” given the symbol
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B with the unit of flux density being the Tesla, T. It is known that the magnetic strength
of an electromagnet depends upon the number of turns of the coil, the current flowing
through the coil or the type of core material being used and if we increase either the
current or the number of turns we can increase the magnetic field strength, symbol H.
The relative permeability, symbol μr is the product of the absolute permeability μ
and the permeability of free space μo (a vacuum) and this was given as a constant.
However, the relationship between the flux density, B and magnetic field strength, H can
be defined by the fact that the relative permeability, μr is not a constant but a function of
the magnetic field intensity thereby giving magnetic flux density as: B = μ H. Then the
magnetic flux density in the material will be increased by a larger factor as a result of its
relative permeability for the material compared to the magnetic flux density in vacuum,
μoH and for an air-cored coil this relationship is given as:
B

A
and
B
 0
H
For ferromagnetic materials the ratio of flux density to field strength (B/H) is not constant
but varies with flux density. However, for air cored coils or any non-magnetic medium
core such as woods or plastics, this ratio can be considered as a constant and this constant
is known as μo, the permeability of free space, (μo = 4.π.10-7 H/m). By plotting values of
flux density, (B) against the field strength, (H) we can produce a set of curves called
Magnetisation Curves, Magnetic Hysteresis Curves or more commonly B-H Curves
for each type of core material used as shown below.
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2.2.1 Magnetic Hysteresis Loop
Figure 2.1 Magnetic Hysteresis of a ferromagnetic core.
The Magnetic Hysteresis loop above (Fig 2.1) shows the behavior of a
ferromagnetic core graphically as the relationship between B and H is non-linear. Starting
with an unmagnetised core both B and H will be at zero, point 0 on the magnetisation
curve. If the magnetisation current, i is increased in a positive direction to some value the
magnetic field strength H increases linearly with i and the flux density B will also
increase as shown by the curve from point 0 to point
a as it heads towards saturation.
Now if the magnetising current in the coil is reduced to zero the magnetic field around
the core reduces to zero but the magnetic flux does not reach zero due to the residual
magnetism present within the core and this is shown on the curve from point
a to point
b. To reduce the flux density at point b to zero we need to reverse the current flowing
through the coil. The magnetising force which must be applied to null the residual flux
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density is called a Coercive Force. This coercive force reverses the magnetic field rearranging the molecular magnets until the core becomes unmagnetised at point
c. An
increase in the reverse current causes the core to be magnetised in the opposite direction
and increasing this magnetisation current will cause the core to reach saturation but in the
opposite direction, point
d on the cure which is symmetrical to point b. If the
magnetising current is reduced again to zero the residual magnetism present in the core
will be equal to the previous value but in reverse at point e.
Again reversing the magnetising current flowing through the coil this time into a
positive direction will cause the magnetic flux to reach zero, point
f on the curve and as
before increasing the magnetisation current further in a positive direction will cause the
core to reach saturation at point
a. Then the B-H curve follows the path of a-b-c-d-
e-f-a as the magnetising current flowing through the coil alternates between a positive
and negative value such as the cycle of an AC voltage. This path is called a Magnetic
Hysteresis Loop.
The effect of magnetic hysteresis shows that the magnetisation process of a
ferromagnetic core and therefore the flux density depends on which part of the curve the
ferromagnetic core is magnetised on as this depends upon the circuits past history giving
the core a form of "memory". Then ferromagnetic materials have memory because they
remain magnetised after the external magnetic field has been removed. However, soft
ferromagnetic materials such as iron or silicon steel have very narrow magnetic hysteresis
loops resulting in very small amounts of residual magnetism making them ideal for use in
relays, solenoids and transformers as they can be easily magnetised and demagnetised.
Since a coercive force must be applied to overcome this residual magnetism, work
must be done in closing the hysteresis loop with the energy being used, being dissipated
as heat in the magnetic material. This heat is known as hysteresis loss, the amount of loss
depends on the material's value of coercive force. By adding additive's to the iron metal
such as silicon, materials with a very small coercive force can be made that have a very
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narrow hysteresis loop. Materials with narrow hysteresis loops are easily magnetised and
demagnetised and known as soft magnetic materials.
2.2.2 Magnetic Hysteresis Loops for Soft and Hard Materials
Figure 2.2 Magnetic Hysteresis Loops for soft and Hard Materials
In Fig 2.2, Magnetic Hysteresis results in the dissipation of wasted energy in the form of
heat with the energy wasted being in proportion to the area of the magnetic hysteresis
loop. Hysteresis losses will always be a problem in AC transformers where the current is
constantly changing direction and thus the magnetic poles in the core will cause losses
because they constantly reverse direction. Rotating coils in DC machines will also incur
hysteresis losses as they are alternately passing north the south magnetic poles. As said
previously, the shape of the hysteresis loop depends upon the nature of the iron or steel
used and in the case of iron which is subjected to massive reversals of magnetism, for
example transformer cores, it is important that the B-H hysteresis loop is as small as
possible.
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2.3 The origin of Atomic moments, spin and orbital states of Electrons
The Schrödinger equation has led to information on the energy levels that can be
occupied by the electrons [8]. The states that electrons can have are characterized by four
quantum numbers.
i.
The principal quantum number (n) with values 1, 2, 3,…. determines the size of
the orbits and describes its energy. Electrons in orbits with n=1, 2, 3,… are
referred to as occupying the K, L, M,…. shells respectively.
ii.
The orbital angular momentum quantum number ( l ) describes the angular
momentum of the orbital motion.
The number ‘ l ’ can take one of the integral
values 0, 1, 2,….,n-1 depending on the shape of orbit. The electros with a certain
quantum number ‘ l ’ are referred as s, p, d… Electrons.
iii.
The magnetic quantum number ml describes the component of the orbital angular
momentum l along a particular direction. In most cases, this direction is chosen
along the applied magnetic field. For a given l , ml  l , l  1,...,0,..., l  1, l.
iv.
The spin quantum number ms describes the electron’s spin s along a particular
direction, usually along he applied magnetic field. The allowed values of ms are
+ ½ and – ½
According to Pauli’s principle, it is not possible for two electrons to occupy the same
state, that is the states of two electrons are characterized by different sets of the quantum
numbers n, l , ms and ml . The maximum number of electrons occupying a given shell is
given by,
n 1
2 (2l  1)  2n2
(2.6)
l 0
Basically the moving electron can be considered as a current flowing in a wire that
coincides with the electron orbit.
The corresponding magnetic effects can then be
derived by considering the equivalent magnetic shell. An electron with an orbital angular
momentum l has an associated magnetic movement
l  
e
2m
l  B l ,
Where  B  Bohr magnetron.
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(2.7)
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The absolute value of the magnetic moment is given by
l  B l (l  1),
(2.8)
and its projection along the direction of the applied field is
lz  ml B,
(2.9)
The situation is different for the spin angular moment is
s   ge
e
2m
s   g e  B s,
(2.10)
Where ge (=2.002290) is the spectroscopic splitting factor or g-factor for the free
electron.
The component in the field direction is
sz   ge ms B,
(2.11)
The energy of the magnetic moment  in a magnetic field H is given by the
Hamiltonian
H  o .H  .B
(2.12)
Where B is the flux density or magnetic induction.
o  4 107 TMA1 is the vacuum permeability.
The lowest energy Eo , the ground –state energy, is reached for  and H parallel using
Eq. (2.11) and ms   1 , one finds for one single electron
2
1
Eo   o sz H   ge ms o  B H   ge o  B H
2
(2.13)
1
For an electron with spin quantum number, ms   1 , the energy equals  ge o  B H .
2
2
This corresponds to an antiparallel alignment of the magnetic spin moment with respect
to the field. In the absence of a magnetic field, the two states characterized by ms   1
2
are degenerate, that is, they have the same energy. However, under an applied field this
degeneracy is destroyed and the splitting in energy levels can be observed. This is
showed in Fig 2.3
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Figure 2.3 Effect of a magnetic field on the energy levels of the two electron states with
ms   1 and ms   1 .
2
2
However, when describing the atomic origin of magnetism, one has to consider orbital
and spin motion of the electros and the interaction between them. The total orbital (L)
angular momentum of a given atom is defined as
L   li
(2.14)
i
where the summation symbol includes all electrons. The same argument applies to the
total spin angular momentum, defined as
S   si
(2.15)
i
Here, one has to bear in mind that the summation over a complete shell is zero, the only
contribution coming from incomplete shells.
The resultants S and
L thus formed are loosely coupled through the spin-orbit
interaction to the resultant total angular momentum J :
J  LS
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(2.16)
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Figure 2.4 Resultant total angular momentum J
This type of coupling is called as Russell-Saunders coupling. This coupling is applicable
to most magnetic atoms. Most of the magnetic properties of different types of materials
depend on the applied magnetic field and temperature [8]. At zero temperature, the
situation is simple because for any of the number of atoms (N) only the lowest level will
be occupied. In this case, one obtains for the magnetization of the system
M   Ng J mB  Ng J J B
(2.17)
However, at finite temperatures, higher lying levels will become occupied.
The extent to which this happens depends on the temperature but also on the energy
separation between the ground-state level and the excited levels, that is, on the field
strength.
The magnetization M of the system can then be found from the statistical average   z 
of the magnetic moment z   g J mB .
This statistical average is obtained by weighing the magnetic moment µz of each state by
the probability that this state is occupied and summing over all states:
M  N  z 
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(2.18)
Chapter II
By substituting x 
Literature Review
 g J  B o H
kT
into Eq. (2.18) and using the relation d ln x  x 1dx
and demx  memx dx, after carrying out the differentiation, one finds
M  Ng B JBJ ( y)
(2.19)
Where BJ ( y) is the so-called Brillouin function, given by
BJ ( y) 
2J 1
(2 J  1) y 1
y
coth

coth
,
2J
2J
2J
2J
y
With
gJ  B o H
kT
(2.20)
(2.21)
Keep in mind that in this expression ‘H’ is the field responsible for splitting the energy
levels [8].
2.4 Classification of magnetic materials
In many atoms the electrons are completely paired, that is, for each electron
spinning in one direction, there is an electron spinning in the opposite direction. The
same situation exists in regard to the orbital motion of the electrons. Thus, the net current
circulating about any given axis is zero and there are no Amperian current in such
substances, consequently the resultant magnetic movement of the atom is also zero. Such
substances show very “weak magnetic effects” that they are called “non-magnetic”.
However, vacuum is the only true non-magnetic medium [9].
All materials show some magnetic effects with the exception of ferromagnetic
group these effects are weak. It has been established both theoretically and
experimentally that all mater may be classified into three groups, according to its
fundamentally different behavior under the action of a magnetic field. These are
classified as i) diamagnetic ii) paramagnetic and iii) ferromagnetic.
The diamagnets and paramagnets consist of atoms with completely occupied
electrons shells. In the absence of an external magnetic field, they do not possess a finite
magnetization. The ferromagnets have spontaneous magnetic order. These are usually
transition metal (3d) atoms with partially filled d and f electron shells.
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Figure 2.5 Types of magnetism: (A) paramagnetism (B) ferromagnetism (C)
antiferromagnetism (D) ferrimagnetisms (E) enforced ferromagnetism
As mentioned already, the magnetic moment of a free atom has three principal
sources: i) the electron spin ii) the electron orbital angular movement and iii) the change
in the orbital momentum induced by an applied magnetic field. The first two effects give
rise to paramagnetic contributions to the magnetization, while the third gives a
contribution to the diamagnetism.
The magnetization M is defined as the magnetic dipole moment per unit volume.
The magnetic susceptibility per unit volume is defined as

M
B
Where
M – Magnetization
B-Magnetic field intensity
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(2.22)
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It is a dimensionless quantity. Substance with a positive susceptibility is called
paramagnets, while materials with negatives susceptibility are diamagnets. The
ferromagnetic materials have a spontaneous magnetic moment, that is, a magnetic
moment exists even when the applied magnetic field is zero.
In ferromagnetic crystals, the spontaneous moments suggests that electron spins
and magnetic moments are arranged in a regular manner. From the thermodynamic point
of view, the entropy of any magnetic materials increases with temperature and decrease
with applied magnetic field. Therefore the spontaneous magnetization must decrease with
temperature. There is a transition temperature Tc above which the spontaneous
magnetization vanishes. This particular transition temperature is called the curieTemperature. Above the Curie temperature, the ferromagnetic materials become
paramagnetic. The paramagnetic susceptibility above Tc is given as [10]

C
T  TC
(2.23)
Where C is curie constant.
The above Eq. (2.23) is called curie-Weiss law. In antiferromagnetic material
spins are ordered in an anti-parallel configuration with zero net moment at temperatures
below the temperature TN, which is called the Neel temperature.
2.5 Semiconductor materials
Now a day’s semiconductor materials have been playing an important role in the
electronic industry. Because of their application in devices such computers and mobile
phones, the search for new semi conductor materials and the improvement of existing
materials has been continuous. The key property of a semi conductor material is that it
can be doped with impurities that alter its electronic property in a controllable way. The
semi conductor material most commonly used is crystalline inorganic solids. These
materials are classified according to the periodic table groups from which their
constituent atoms come. Table.1 shows a list of some most common semiconductors. In
this table it is important to note that oxide semi conductors, such as SnO2, ZnO or TiO2
are being widely used as DMS material.
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Table.1 Some common semi conducting materials
Semi conductor
Group
Silicon (Si)
IV
Germanium (Ge)
IV
Gallium arsenide (GaAs)
III-V
Indium nitride (InN)
III-V
Cadmium telluride (CdTe)
II-VI
Zinc Selenide (ZnSe)
II-VI
SnO2
Oxide semiconductor
TiO2
Oxide semiconductor
ZnO
Oxide semiconductor
2.5.1 Transport phenomena in semi conductors
In the second half of the 1940s, Bardeen, Brattain and Shockley discover the
transistor effect. After the discovery of the transistor effect, a break in the behavior of
doped Germanium conductivity at low temperature was observed. In 1956, Conwell [2]
and Mott [3] suggested a model for a “new” process of conduction in which charge
carriers conducts the electric current by thermally activated tunneling from an occupied
site to an empty site. This process has been known as “phonon assisted hopping” and was
the starting point of a number of transport theories, such as the model of miller and
Abrahams [4]. This model became the most widely accepted theory of conduction
between localized states and is the source of the variable range hopping (VRH) theory of
Mott[5].
According to this model, electrons are “intelligent” since in the conduction
process they hop from an initial state “i” to another “j” with energy as low as possible.
For such an energy the site “j” is statistically located for from “i”, involving a distance
rij=R, which generally is much larger than the decay length of the wave function.
According to Mott, the conductivity ‘σ’ depends on two factors (i) Boltzmann
 W
factor exp    and (ii) overlap of the wave functions exp  2 R  , where
 k 
K - Boltzmann constant
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T - Temperature
W - Hopping Energy
 1 - Wave function decay length
R - Distance separating Sites “i” and “j”
W 

Mott consider that the conductivity “  ” behaves as the factor exp  2 R 
 .
kT 

So it may be written as


   o exp  2 R 
W 

kT 
(2.1)
Where σo is a factor that is weakly temperature dependent. The idea of Mott was to
maximize the hopping probability.
So the express for the hopping energy W may be written in the form
W
3
4 R N ( EF )
(2.2)
3
This gives the minimum value of “R” as
R3 D 
3
1
2
8 N ( EF )k  4 T
1
1
(2.3)
4
where N (EF) is the density of states at the Fermi level. It then appears that the distance of
a hop increases as T decrease. Substituting Eq. (2.2) and Eq.(1.3) into Eq. l (2.1) we get
1
the T 4  Mott conductivity for non-crystalline semiconductor as


 3D   o
3D
  T 3D  14 
exp    o  
  T  


(2.4)
This mechanism is a powerful tool to explain and characterize electrical transport in
semiconductors. Unfortunately not all semiconductors can be fitted under this model.
The nearest neighbor hopping or Granular metal mechanism are just a few examples of
various mechanism that have been introduced as an alternative to variable range hopping
theory. In many cases the transport properties measured show a composite of different
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transport mechanism. Thus different models can be applied at different temperatures,
which prove that the transport in a semiconductor is a very complicated matter.
2.6 Dilute magnetic semi conductors
Room temperature ferromagnetism in semiconductor materials has been studied
for few years and has received renewed attention, to some extent because of possible uses
in spintronic application [11]. The research in this field is primarily focused on transition
metal doped II-VI, IV-VI, II-V and III-V compound semiconductors. Strong
ferromagnetic interactions between localized spin have been observed in Mn-doped II-VI
compounds. However, recently the group of materials researched has been extended to
various oxide semiconductors as well. Among these oxides the strong favorite is SnO2.
Recently room temperature ferromagnetism has been reported for SnO2 doped with
transition metals, showing the potential for achieving room temperature spintronic
technologies [12].
2.6.1 Ferromagnetism in oxide dilute magnetic semiconductors
In dilute solid solutions of a magnetic ion in a non magnetic metal crystal (such as
Mn in Sn) the exchange coupling between the ion and the conduction electron has
important consequences. The conduction electron gas is magnetized in the vicinity of the
magnetic ion, with the spatial dependence. This magnetization causes an indirect
exchange interaction between two magnetic ions. The interaction is known as the Fried
or Ruder man, Kittle, Kasuya and Yosida (RKKY) interaction.
A consequence of the magnetic ion-conduction electron interaction is the Kondo effect.
The occurrence of a resistance minimum is connected with existence of localized
magnetic moments on the impurity atoms. Where a resistance minimum is found, there is
a inevitably a local moment.
The central result is that the spin-dependent contribution to the resistivity is

 Spin  C  M 1 


3zJ
ln T   C o  C 1 ln T
EF

where J - exchange energy
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(2.23)
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Z -number of nearest neighbors
C -Concentration
 M -Measure of the strength of the exchange scattering
We see that the spin resistivity increases toward low temperature if J is negative.
If the phonon contribution to the electrical resistivity goes as T5 in the region of interest
and if the resistivities are additive, then the total resistivity has the form
  aT 5  C o  C 1 ln T
(2.24)
d
C
 5aT 4  1  0
T
dT
(2.25)
 C 
 1
 5a 
(2.26)
With a minimum at
Hence
15
Tmin
The temperature at which the resistivity is a minimum varies as the one-fifth
power of the concentration of the magnetic impurity, in agreement with experiment at
least for Fe in Cu.
Recently S.J. Liu et al [13] have discussed a situation where ferromagnetism in
magnetic semiconductors originates from Itinerant-Electron magnetism. This model is
suitable for semiconductors that are rich in charge carriers. The localized moments arises
when the magnetic atoms form part of an intermetallic compound. In these cases, the
unpaired electrons responsible for the magnetic moment are no longer localized and
accommodated in energy levels belonging exclusively to a given magnetic atom. Instead,
the unpaired electrons are delocalized; the original atomic energy levels having
broadened into narrow energy bands. Extend of this broadening depends on the
interatomic separation between the atoms. According to a calculation made by Heine
(1967), the following relation applies between the width of the energy bands W and the
interatomic separations (r).
W α r-5
(2.27)
The most prominent examples of itinerant-electron system are metallic system based on
3d transition elements, with the 3d electrons responsible for the magnetic properties. For
a discussion of the magnetism of the 3d electron bands, we will make the simplifying
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assumption that these 3d bands are rectangular. This means that the density of electron
states N(E) remains constant over the whole energy range spanned by the band width W.
A maximum of ten 3d electrons per atom, that is, five electrons of either spin direction,
can be accommodated in the 3d band. In the case of cu metal, because each cu atom
provides ten 3d electrons, the 3d band will be completely filled.
However, in the case of other 3d metals, less 3d electrons are available per atom
so that the 3d band will be partially empty. Such a situation is shown in Fig 2.6.a.
In Fig.2.6.a, we have indicated that there is no discrimination between electrons of spinup and spin-down direction with respect to band filling. Both types of electrons will
therefore be present in equal amount, meaning that there is no magnetic moment
associated with the 3d band in this case. However, this situation is not always a stable
one, as will be discussed below.
It is possible to define effective exchange energy U eff per pair of 3d electrons.
This can be regarded as the energy gained when switching from antiparallel to parallel
spin. In order to realize such gain in energy, electrons have to be transferred, say, from
the spin-down sub band into the spin-up sub band. This can be seen in Fig.2.6.b, this
implies an increase in Kinetic energy, which counteracts this electron transfer. However,
it will be shown below that such transfer is likely to occur if U eff is large and density of
states at the Fermi level EF is high. After the transfer, there will be more spin-up electrons
than spin-down electrons, and the magnetic moment, which has arisen, will be equal
to   (n1  n2 )B .
First we will derive a simple band model, which accounts for the existence of
ferromagnetism. The interaction Hamiltonian, can be written as
H  U eff n1n2
(2.28)
Where n1 and n2 represents the number of electrons per atom for each spin state, and
where the total number of 3d electrons per atom equals n  n1  n2 . Because Ueff is a
positive quantity, Eq. (2.28) will lead to the lowest energy if the product n1n2 is as small
as possible.
For equally populated sub bands, this product has its maximum value and hence the
highest energy, consequently, electron transfer is always favorable for the lowering of the
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exchange energy and this electron transfer will come to an end only if one of the two spin
sub bands is empty or has become completely filled up.
It is defined that N(E) as the density of states per spin sub band, and P as the fraction of
electrons that has moved from the spin-down band to the spin-up band.
This means
2 pn  n1  n2
(2.29)
Let us assume that the interaction Hamiltonian (Eq.2.28) leads to an increase in
the number of spin-up electrons at the cost of the number of spin -down electrons. The
corresponding gain in magnetic energy is then [8].
EM  U eff n1n2  U eff
1 2
n ,
4
EM  U eff n2 p 2
(2.30)
This energy gain is accompanied by an energy loss in the form of the amount of energy
EC needed to fill the states of higher kinetic energy in the band. For a small
displacement  E  EF  E2  E1  EF (see Fig. 2.6.b) this kinetic energy loss can be
written
1
as EC   E (n1  n2 )   Enp .
2
The
total
energy
variation
is
1
then EC  EM  U eff n2 p 2   Enp . Since, N ( EF ) E  (n1  n2 ), one may write,
2
E  EC  EM 
n2 p 2
1  U eff N ( EF ) 
N ( EF ) 
(2.31)
If 1  U eff N (EF )   0 , the state of lowest energy corresponds to p=0 and the
system is non-magnetic. However, if 1  U eff N ( EF )   0 , the 3d band is exchange split
(p>0), which corresponds to ferromagnetism, this is the stoner criterion for
ferromagnetism, which is frequently stated in the more familiar form (stoner, 1946)
If U eff N (EF )  1 , by means of this model, it can be understood that 3d magnetism
leads to non-integral moment values if expressed in Bohr magnetrons per 3d atoms,
  2 pnB . The conditions favoring 3d moments in metallic system are obviously; a
large value for U eff , but also a large value for N (EF). The density of states of the S- and
P-electron bands is considerably smaller than that of the d band, which explains why
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band magnetism is restricted to elements that have a partially empty d band. However,
not all of the d-transition elements give rise to d-band moments. For example, in the 4d
metal Pd, the stoner criterion is not met, although it comes very close to it.
1→E1,
2→EF ,
3→E2
Figure 2.6 Schematic representation of a partially depleted 3d band: (a) Paramagnetism,
(b) Weak ferromagnetism, (c) Strong ferromagnetism with n  5 , (d) Strong
ferromagnetism with n  5 .
2.6.2 3d elements doped Oxide semiconductor: SnO2
Recently SnO2 has been most widely researched oxide semiconductor for DMS
materials. SnO2 is a direct wide band-gap semiconductor with applications in
photovoltaic devices, gas detectors and transparent electronic devices such as transparent
widow heaters. The band gap of SnO2 is ~3.7eV at room temperature.
Another
advantage of SnO2 is that this band gap can be easily controlled by introducing interstitial
ions, such as CO-, Fe-, Ni-, V- and W- by varying Sn composition, the direct energy band
gap can be adjusted from 3.7 to 4.15 eV [13]. A controllable band gap is very important
for spin life time and spin transport in semiconductor based spintronic devices [14].
SnO2 has a Rutile (tetragonal), tp6 type crystal structure (2, 4) with lattice parameter
a= 4.737 Ao and C= 3.186 Ao. The mineral form of SnO2 is called cassiterite and this is
the main ore of tin. This colourless, diamagnetic Solid is amphoteric SnO2 is usually
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regarded as an oxygen-deficient n-type semiconductor. “Stannic acid” refers to hydrated
Tin dioxide, SnO2, which is also “stannic hydroxide”.
Figure 2.7 crystal structure of SnO2
For spintronic, theoretical prediction [15] suggests that room temperature ferromagnetism
is possible in Mn doped SnO2. Experimental results on 3d metal doped SnO2 are very
controversial. In fact, Mn-doped SnO2 films showed room temperature ferromagnetism
in some studies [17], while other observed Para magnetism [16].
Mn is just one of many transition metal ions that have been used in the synthesis
of magnetic SnO2. The samples were synthesized in the research for the best candidate in
the family of materials. So far, Ni-doped [18, 23, 29, 31], V-doped [20, 21, 22], Codoped [27, 25], W-doped [35] SnO2 have been reported. They were reported to be both
ferromagnetic and non-ferromagnetic even for the same sample or compositions.
Soft chemical technique or sol-gel is the most effective method that can help to
avoid the formation of unneeded impurities [36]. Such synthesis can provide better
control over the composition of materials than is obtained with high temperature vacuum
deposition or solid state synthesis technique. On the other hand, it is a simple, lower-cost
method, and easier. The solution made by this chemical technique could provide a
homogeneous distribution of dopant ions throughout the sample. Furthermore, the whole
synthesis process was performed in the absence of vacuum, which could prevent the
formation of metal nanoparticles. In the present research work, the 3d-metal doped and
codoped SnO2 nanocrystalline DMS were synthesized by soft chemical techniques or solgel method.
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The recent research progress on oxide dilute magnetic semiconductors has been
summarized in this section. It is evident that, many research groups have reported
observation of room temperature ferromagnetism in oxide semiconductors. However,
reports on the absence of ferromagnetism or impurity mediated ferromagnetism in such
materials are also numerous. It is of very great importance to synthesis and analyzes
materials that would be pure intrinsic magnetic semiconductors. From all of the oxide
semiconductors reported, SnO2 is probably the best candidates for spintronic applications.
2.6.3 Mn-doped SnO2
K. Gopinadhan et al. (2009) have investigated that it is possible to induce room
temperature ferromagnetism, in tin oxide thin films by introducing manganese in a SnO2
lattice; observed temperature dependence of the magnetization predicts a curie
temperature exceeding 550 k. The generation of additional free electrons by F doping in
Mn-doped SnO2 does not cause any increase in the magnetic moment per Mn ion,
suggesting no significant role of electron in bringing about the magnetic ordering [15].
Davinder Kaur et al. (2009) synthesized Mn doped SnO2 thin films. Magnetic studies
reveal RTFM, a systematic change in magnetic behavior from ferromagnetic to
paramagnetic was observed with increase in substrate temperature from 500 to 700o C for
Mn doped SnO2 film. The presence of room-temperature ferromagnetism could be
obtained by controlling the substrate temperature and Mn doping concentration [16].
Susmita kunda et al. [32] has reported that room temperature ferromagnetism in highly
transparent Mn (II) - doped ITO films deposited on soda lime silica glass and pure silica
glass. The Mn-O bond length and the Mn-Mn pair exchange interaction is found to be
depend on the substrate smaller bond length and stronger exchange interaction is found in
the case of films deposited on soda lime glass substrate and larger bong length and
weaker exchange interaction in the case of films deposited on pure silica glass substrate.
They achieved both magnetic ions doping in nanocrystals of In2O3 as well as strong room
temperature ferromagnetism without any degradation of optical transparency which could
be useful for many opto-spintronics based devices [32]. Z.M.Tian et al. (2008) prepared
Mn doped SnO2 nanoparticles by chemical co-precipitation method. They systematically
investigated the microstructure and magnetic properties of the prepared samples. The X-
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ray diffraction reveals that all samples are pure rutile type tetragonal phase. They found
that the cell parameters a and c decrease monotonously with the increase of Mn content
which indicated that Mn ions substitute into the lattice of SnO2. The XPS studies indicate
the mixture state of Mn3+ and Mn4+ in the samples. The magnetic investigations
demonstrate that the magnetic properties strongly depend on both the sintering
temperature and doping content. They found that no ferromagnetism was observed for
samples sintered at high temperature irrespective of the doping content. Also the samples
sintered at low temperature posses’ room temperature ferromagnetism when the doping
content is below 5% [34]. Yanxue chen et al. (2009) has investigated to the influence of
the growth temperature on the electrical conductance, the crystal structure, the surface
morphology and the optical property of Mn doped SnO2. X-ray diffraction pattern show
that the thin films grown at lower substrate temperature were amorphous, while those
grown at higher temperatures had only a single orientation. Atomic force microscopy
results indicate that the films grown at low temperature were extremely flat and they
became rougher as substrate temperature increased. Optical transmission reveals that the
absorption edge shifts towards the blue side and become steeper at increasing substrate
temperature. High- resolution transmission electron microscopy analysis clearly indicates
that the samples have a well-arranged tetragonal structure without any detectable
secondary phase. Inhomogeneous incorporation of Mn atoms into the SnO2 lattice was
observed [37]. B. Sathyaseelan et al. (2010) synthesized pure and transition metal ion
doped SnO2 nanoparticles using a simple chemical precipitation method. (Mn and Co) of
1, 3, and 5 mol% were doped in order to study the influence of structural and optical
properties. The structural, chemical and optical properties of the samples were analyzed
using x-ray diffraction, scanning electron microscopy, energy dispersive x-ray analysis,
and UV-Vis spectroscopy techniques. The SnO2 crystallites were found to exhibit
tetragonal rutile structure with lattice parameters, revealing that the metal ions get
substituted in the SnO2 lattice. It is observed that the peak position (110) get shifted to
lower angles by increasing the dopant concentration. Hence the lattice parameters and
the cell volume get decreased. The SEM photograph shows that the grain size of pure
SnO2 is larger that of the metal-doped SnO2 indicating grain grown upon doping. TEM
photograph revels that by increasing the content of metal ions, average grain size
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decreased. A significant red shift in the UV absorbing band edge was observed with the
increase in the amount of Mn and Co contents [38]. H. Kimura et al. (2001) fabricate Mn
doped SnO2 Epitaxial films by pulsed laser deposition method. Thin films with (101)
and (100) orientations were grown on r-(102) and c (001) sapphire substrates,
respectively. They found that Mn ions are soluble into SnO2 films up to 30 mol%.
Transmission spectra show d-d transition absorption in mid –gap region due to presence
of the Mn ion. Additional doping of Sb induces an n-type conduction with a carrier
concentration as high as 6.9x1019cm-3 at 300k. The resistivity rapidly increases with
decreasing the temperature below 50k, where considerable increase of the resistivity is
observed in a magnetic field below 20k [39]. Feng Gu et al. (2003) studies the room
temperature photo luminescent properties of Mn- doped tin oxide nanoparticles. The
samples are crystalline with a tetragonal rutile structure of tin oxide. The FTIR and the
UV-Vis absorption spectra of the samples have also been investigated. The photo
luminescence spectra are measured at room temperatures as a function of different
sintering temperatures and the Mn2+ concentration, respectively. The luminescence
processes are associated with oxygen vacancies in the host and related with the
recombination of electrons in single occupied oxygen vacancies with photo excited holes
in the valence band [40]. A.J Nadolny et al. (2002) fabricated the thin layers of semi
magnetic Sn1-xMnxTe semiconductor with Mn content x≤0.04 and layer thickness 0.22µm were grown by molecular beam epitaxy(MBE) on BaF2(111) substrates with SnTe
and Mn fluxes, they employed an additional Te flux to provide the efficient way of
controlling the deviation from stoichiometry in the layers. By manipulating the MBE
conditions they obtained SnMnTe crystalline layer with the conducting hole
concentration in the range from P=5x1019 cm-3 to 2x1021 cm-3. In the layers grown under
extra Te flux condition they observe ferromagnetic transition with the Curie temperature
Tc ≤ 6k.
Layers with the same Mn content but grown with no extra Te flux are
paramagnetic. This dramatic change of magnetic properties, referred to as a carrier
concentration induced ferromagnetic transition, can be understood in the frame work of a
model attributing the ferromagnetic interactions in SnMnTe to the mechanism of indirect
exchange via conducting holes [41]. Y.W.Heo et al. (2004) studies the effects of high
dose implantation of Ni, Fe, Co and Mn ions into SnO2. The x-ray diffraction showed no
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evidence of secondary phase formation in the SnO2. The Mn –implanted SnO2 remained
paramagnetic, as also reported for samples doped during thin film growth, but the Fe, Co,
and Ni–implanted SnO2 showed evidence of hysteresis with approximate curie
temperature of 300k. The carrier density in the implanted region appears to be too low to
support carrier-mediated origin of the ferromagnetism and formation of bound magnetic
polarons may be one explanation for the observed magnetic properties [42].
K.Gopinadhan et al. (2007) induce room temperature ferromagnetism, exhibiting high
transition temperature dependence of the magnetization in a SnO2 lattice. The
temperature dependence of the magnetization study predicts a curie temperature
exceeding 550 k in these films.
X-ray and electron diffraction investigations, in
combination with magnetization studies, reveal that the RTFM is not due to any
ferromagnetic impurities. X-ray structure factor calculations suggest that Mn has been
incorporated in the SnO2 matrix. Also, the optical absorption study predicts a random
alloy formation for Sn1-xMnxO2-δ systems in conformity with the predictions of the other
observed properties. Different levels of F doping in Sn0.09Mn0.10O2-δ have been carried
out, and the results indicate that the RTFM is not carrier mediated [50].
2.6.4 Ni-doped SnO2
Nguyen Hoa hong et al. (2005) had developed Ni-doped SnO2 thin films on
different substrates. All the sample shows room-temperature ferromagnetism and films
grown on LAO (LaAlO3) substrates have a large magnetic moment, about one order
larger than that of films grown under the same conditions on STO (SrTiO3) and Al2O3
substrates [17]. Junying zhang et al. (2010) reported that magnetic moment per Ni ion
decreases with the increase of Ni doping because antiferromagnetic Super– exchange
interaction takes places in the nearest neighbour Ni2+ ions for the samples with high Ni
content and also the annealing at oxidizing and reducing atmosphere shows oxygen
vacancies play a crucial role in producing ferromagnetism [18]. C.M.Liu et al.(2007) had
investigated the influence of Nickel dopant on the microstructure and optical properties
of SnO2 nanopowders. XRD measurements indicate that the grain size becomes larger
and the structure of lattice grows better with low nickel content and the obtained samples
are polycrystalline with rutile structure of pure SnO2, no impurity phase cannot present.
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From the optical studies the band gap diminishes with the increase of nickel content. The
nickel doping does not change the PL peak position. The sintering temperature increases
the ratio of UV/blue decreases, the band gap increases and finally saturated [29]. C.M.
Liu et al. (2009) explored the role played by nickel doping in magnetic property and PL
of SnO2 nanopowders. They found that there is no need for transition metal doping to
change the defects into magnetic polarization as the undoped SnO2 nanopowders are
magnetic and nickel doping disturbs the magnetic interaction of pure SnO2 nanopowders.
The PL peaks position changes little with dopant content, UV emission is separated into
two luminescence peaks at higher doping [31]. X.Liu et al. (2010) fabricated Ni- doped
SnO2 hollow spheres and characterized by x-ray diffraction, x-ray photoelectron
spectroscopy,
inductively
coupled
plasma–optical
emission
spectroscopy
and
transmission electron microscopy[44].
2.6.5 Fe-doped SnO2 and Cu-doped SnO2
Aditya Sharma et al. (2011) had developed Fe and Ni doped SnO2 nanoparticles
using simple wet chemical method. The XRD pattern indicates that Fe and Ni are
incorporated into SnO2 lattice without forming any oxide phases. UV-Visible absorption
spectra results indicate the decrease of optical band gap with increasing concentration.
From the room temperature magnetization measurements, they observed ferromagnetic
hysteresis loops and magnetic moments per Fe\Ni ions were found to decrease with
increase of Fe and Ni concentrations.
The decrease in the magnetic moment with
increase in concentration is consistent with the known super-exchange interaction among
the transition metal ions [26]. C. B. Fitzgerald et al. (2004) synthesized SnO2 doped
(Mn, Fe, or Co) dilute magnetic semiconductors and they found room temperature
ferromagnetism in it. They found the formation of a rutile-structure phase from x-ray
diffraction. For (Sn0.95Fe0.05) O2 the saturation magnetization is 0.2 and 1.8 Am2kg-1.
The Curie temperature is found to be 340 and 360 k respectively. From the x-ray
diffraction pattern, the magnetization cannot be attributed to any impurity phase. The
57
Fe Mossbauer spectra of the Fe-doped SnO2 samples, recorded at room temperature and
16k, show that about 85% of the iron is in a magnetically ordered high spin Fe3+ State,
the remainder being paramagnetic [33]. L. M. Fang et al. (2009) prepared Fe3+-doped
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SnO2 nanoparticles by sol-gel-hydrothermal route. The SnO2 crystalline with tetragonal
rutile structure could form directly during the hydrothermal process without calcination
Fe formed stable solid solutions in SnO2 nanoparticles. Compared to sol gel calcination
route, sol-gel hydrothermal route led to better disperse spherical Fe3+-doped SnO2 nano
particles with narrow size distribution and larger specific surface area. The composite
nanoparticles prepared by the sol-gel-calcination route. Large surface areas of Fe3+-doped
SnO2 nanoparticles prepared by sol-gel-hydrothermal route make it particularly appealing
in applications for gas sensors and optoelectronics devices [36]. C. Van komen et al.
(2008) reports the structural and magnetic behavior of transition metal doped SnO2
nanopowders by sol-gel based chemical method. Relatively stronger ferromagnetism was
observed in Co, Fe and Cr ions doped samples. In these systems, strong correlations
between the changes in the structural and ferromagnetic properties were observed. A
limiting dopant concentration XL was observed below which the SnO2 lattice contracts
and ferromagnetic magnetization strengthens, both in proportion to x. This is attribution
due to substitutional incorporation of the transition metal ions as the dominant process
the SnO2 lattice expands and the ferromagnetic behavior is destroyed. They attribute this
behavior to additional interstitial incorporation of the dopant ions and depletion of the
free carriers [43]. J. M. D. Coe et al. (2003) fabricated thin films by pulsed –laser
deposition. The Fe-doped SnO2 thin films are transparent ferromagnets with Curie
temperature and spontaneous magnetization of 610k and 2.2 Am2kg-1 respectively. The
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Fe Mossbauer spectra show the iron is all high-spin Fe3+. But the films are magnetically
inhomogeneous on an atomic scale, with only 23% of the iron ordering magnetically.
The net ferromagnetic moment per ordered iron ion, 1.8µB, is greater than for any simple
iron oxide with super exchange interactions. Ferromagnetic coupling of ferric ions via an
electron trapped in a bridging oxygen vacancy is proposed to explain the high Curie
temperature [45]. S. Sambasivam et al. (2011) prepared a series of Fe-doped SnO2
nanoparticles by sol-gel method. Structure of the samples is found to be rutile and the
crystalline size is found to decrease with increase of Fe content. The g- values of the
ESR spectra reveal that the nature of Fe in the prepared samples is isolated rhombic Fe3+
ions in rutile phase, and the line width of ESR increase with increase in Fe3+-content due
to the ion induced disorder effect. The temperature dependent magnetization confirms
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that the Fe3+-ions have weak antiferromagnetic interactions. However, for the samples
with high Fe content, an extra spin-pumping is observed below 250k. The optical
absorption spectra of Fe doped SnO2 nanoparticles showed a blue shift in the band gap
compared to the undoped SnO2 samples [47]. Junko sakuma et al. (2007) prepared
transparent conducting SnO2 powders by a polymerized complex method. It is confirmed
from XRD data that the powder samples prepared from acidic solution have the rutile
structure of
The
57
SnO2 and that the annealing temperature affects the structural parameters.
Fe Mossbauer spectra consist of only doublets, including no magnetic sextets.
However, the samples prepared at 500◦c and 600◦c showed ferromagnetism. The
magnetization of the sample at 550◦c has larger magnetization than the sample at 600◦c
[48]. Khaled melghit et al. (2005) prepared nanosize solid solution of Fe-doped SnO2 at
low temperature. The XRD data of the samples shows that the sample is rutile phase
without any trace of an extra phase. SEM and TEM reveal a homogeneous composition
of the sample. The size was found to increase slightly with temperature that Fe3+
substitute for Sn4+ [49].
2.6.6 Co-doped SnO2
K. Srinivas et al. (2008) had investigated the structural, optical and magnetic
properties of nanocrystalline Co doped SnO2 based diluted magnetic semiconductors.
From the XRD pattern of all the samples reveals that all the samples have tetragonal
rutile structure. The average crystalline size varies from 15 to 40 nm. From the FTIR
spectra it is clear that there are clear changes in the position, size and shape of IR peaks
indicating the Co ions incorporate in SnO2 host. From the magnetization studies they
conclude that cobalt doping enhanced the specific magnetization value. The Room
temperature ferromagnetism have been achieved by modifying oxygen vacancies and
electronic structure through stabilization of unstably bonded Co2+ ions by controlling the
particle size, shape and surface stoichiometry. In the case of ferromagnetic nanocrystals,
the strength of exchange coupling between magnetization of adjacent ferromagnetic
grains also depends on the magnetic nature of the matrix [27]. Weibing chen et al. (2011)
reported that SnO2 doped with less than 5% Co synthesized by sol-gel-Hydrothermal
method exhibits room temperature ferromagnetism. From the XRD pattern any other
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secondary phases or clusters of cobalt cannot present and the sample holds a curie
temperature as high as 400K and shows a combining character of ferromagnetism and
antiferromagnetism [28].
C. B.
Fitxgerald et al. (2006) reported that the high
temperature ferromagnetism in thin films of transition metal doped SnO2 is an intrinsic
effect. The moments at low concentrations approach the spin-only values, and they
cannot be explained in terms of secondary impurity phases. Nor can the anisotropy of the
magnetization of the films be explained in these terms. Absence of anomalous Hall
Effect indicates that there is negligible 3d density of states at EF. Much of the magnetic
moment is associated with electronic defects of lattice defects. The dopants somehow
serve to activate the defect moment, since the undoped SnO2 is net ferromagnetic. The
high Curie temperature observed for low dopant implys that the defect related magnetism
is inhomogeneous [51]. X. F. Liu et al. (2009) fabricate Co-doped SnO2 insulating films
using magnetron sputtering. X-ray photoelectron spectrum reveal a solid solution of Co
dopants in SnO2 lattice, where Co is in 2+ oxidation state and substitutes for Sn4+. The
films exhibit clearly room-temperature ferromagnetism and with more structural defects
show higher saturated magnetic moment, while increase of annealing time, crystallinity
of the films is improved and the defects decrease, leading to the decrease of
ferromagnetism. This shows that the ferromagnetism in Co-doped SnO2 is intrinsic and
can be considerably influenced by the concentration of structural defects [52]. Hongxia
Wang et al (2009) reported the effects of Co-dopants and oxygen vacancies on the
electronic structure and magnetic properties of the Co-doped SnO2 are studied by the first
principle calculations in full-potential linearized augmented plane wave formalism within
generalized gradient approximation. The Co atoms favorably substitute on neighboring
sites of the metal sub lattice.
For Co-doped SnO2 without oxygen vacancies, spin
polarization occurs mainly at Co site and Co is in its low spin state, which is independent
of the concentration and distribution of Co atoms. Furthermore only nearest-neighbor Co
atoms are ferromagnetically coupled through direct exchange and super exchange
interaction. Oxygen vacancy prefers to locate near the Co atom and strongly increases
the local magnetic moment of Co atom and strongly increases the local magnetic moment
of Co atoms. The presence of the vacancy cause calculated magnetic moment to depend
sensitively on the concentration and distribution of Co atoms and induce the long-range
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ferromagnetic couplings between Co atoms with large distance through the spin-split
impurity band exchange model which contribute to the room temperature ferromagnetism
observed experimentally in the Co-doped SnO2 [53]. L. Yan et al. (2009) has reported
the Co-doped SnO2 thin films deposited on Si substrates by pulsed- laser deposition. The
XRD results showed that the 5at% Co-doped SnO2 with rutile- structure could be well
grown on Si substrate. The Co-doped SnO2 thin film has good room temperature
magnetic property with the saturated magnetic moment of 1.3µB/Co at 293 k. The x-ray
photoelectron spectroscopy measurements for the Co-doped SnO2 thin film on Si
predicted that Co has oxidation states of +3. The good magnetic property of the Codoped SnO2 may be deduced the exchange between Co3+ and Co3+ through the oxygen
vacancy [54]. Hongxia Wang et al. (2009) studies the electronic structure and
ferromagnetic stability of Co-doped SnO2 using the first-principle density functional
method within the generalized gradient approximation (GGA) and GGA+U schemes. The
addition of effective Uco transforms the ground state of Co-doped SnO2 to insulating from
half-metallic and the coupling between the nearest neighbor Co spins to weak
antimagnetic from strong ferromagnetic. GGA+Uco calculations show that the pure
substitutional Co defects in SnO2 cannot induce the ferromagnetism. Oxygen vacancies
tend to locate near Co atoms. Their presence increases the magnetic moment of Co and
induces the ferromagnetic coupling between two Co spins with large Co-Co distance.
The calculated density of state and spin density distribution calculated by GGA+Uco
show that the long-range ferromagnetic coupling between two Co spins is mediated by
spin-split impurity band induced by oxygen vacancies. More charge transfer from
impurity to Co-3d states and large spin split of Co-3d and impurity states induced by the
addition of Uco enhance the ferromagnetic stability of the system with oxygen vacancies.
By applying a coulomb U0 on Oxygen 2s orbital, the band gap is corrected for all
calculations and the conclusions derived from GGA+Uco calculation are not changed by
the correction of band gap [55]. Yongbin xu et al. (2009) prepared Co-doped SnO2
diluted magnetic semiconductor nanocrystals successfully and obtained roomtemperature ferromagnetism. The magnetic field-assisted approach has been used in the
synthesis of Co-doped SnO2 diluted magnetic semiconductor nanocrystals. They found
that the external high magnetic field can influence the growth behavior of Co2+ doped
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SnO2 nanocrystals, and ferromagnetic properties would improve obviously various
techniques such as x-ray diffraction, transmission electron microscope, UV-Visible
spectrometry, Raman spectrometry and Vibrating sample magnetometer have been use to
characterize the obtained products [56]. N. Brihi et al. (2007) have used the coprecipitation technique to synthesize polycrystalline Co-doped tin oxide diluted magnetic
semiconductors of different concentrations. The X-ray diffraction analysis confirmed the
rutile structure of SnO2. Transmission electron microscopy confirms the uniform
distribution of Co inside the samples. Optical absorption measurements showed an
energy band gap which decreases when increasing the concentration of Co. Raman
spectroscopy indicates the significant structural modifications and disorder of the SnO2
lattice. Magnetic measurements revealed a mixture of para and antiferro behavior of Codoped SnO2 [60].
2.6.7 Vanadium (V) -doped SnO2 and Tungsten (W)-doped SnO2
Sang-Do Han et al. (2000) prepared V-doped SnO2 nano crystalline powder by
co-precipitation method. The average crystal size decreases with increasing amount of
doped Vanadium. The morphology of nanocrystalline vanadium-doped SnO2 particle is
spherical and the distribution of crystal size is relatively uniform [20]. Li Zhang et al.
(2005) reported that the V-doped SnO2 films on Si (III) shows RTFM with no trace of
any impurity phases. The magnetic moment per V drops rapidly with the increase in V
content. The annealing study of the films in rich-oxygen and poor-oxygen atmospheres
reveals that the rich-oxygen annealing may introduce more oxygen ions, hence reduce
oxygen vacancies, which are responsible for the decrease in Ms, while the vacuum
annealing may remove oxygen vacancies, and so favor for room- temperature
ferromagnetism. Therefore, these results that the oxygen vacancy plays a crucial role in
inducing the ferromagnetism in V-doped SnO2 and the origin of ferromagnetism can be
explained by the BMP Model [21].
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Figure 2.8 Vanadium atom (green) in the SnO2 rutile structure.
Nguyen Hoa Hong at al. (2005) developed V-doped SnO2 thin films were grown on LAO
(LaAlO3) , STO (SrTiO3) and R-Cut Al2O3 substrates. The V-doped SnO2 Films grown
on Al2O3 do not show ferromagnetism in the whole range of temperature, while the films
grown on both LAO and STO show room temperature ferromagnetism. It is noted that
the V-doped SnO2 Films grown on LAO substrate is larger than that of the V-doped
SnO2 films grown on STO substrate. From the above results the substrate is one of the
factors for inducing RTFM [22]. N. Tahir et al. (2009) synthesized V-doped ZnO
nanoparticles by sol-gel method. XRD patterns of samples show that all samples exhibit
wurtzite structure with no secondary phases. The presence of V-doping was confirmed
by the EDX analysis. SEM images reveal that particles are almost of spherical in nature.
The TPR analysis indicates the electronically modified material with V doping. The air
annealed sample show no ferromagnetism, while samples annealed under reducing
atmosphere show clearly room temperature ferromagnetism and there is an increase in
saturation magnetization with increase in vanadium concentration. They believe that
RTFM is due to the defects or oxygen vacancies. The blue shift in band gap energy has
been observed. This might be due to the defects such as oxygen vacancies leading to the
Burstein Moss shift [24]. X. L. Wang et al. (2008) had investigated the ferromagnetism
in SnO2 doped with transition metals. Their calculation show that ferromagnetism is
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obtained by doping SnO2 with Fe and Co ions and paramagnetism is ground state in Vand Mn- doped SnO2. They found that oxygen vacancy has a strong influence on the
magnetic properties of the transition metal (TM) ion doped compounds, TM-oxygen-TM
groups will be common in Fe- and Co- doped SnO2, but are not formed in V- and Mn
doped SnO2. Additional n-type carriers are required to stabilize the ferromagnetic state
and the Curie temperature increases with increasing densities of n-type carriers in Vdoped SnO2. They also propose that raising the density of n-type carriers is a possible
practical way to obtain high Curie temperature [30]. Yanwei Huang et al. (2009) had
investigated on structural, electrical and optical properties of tungsten-doped tin oxide
thin films by pulsed plasma deposition method with post annealing. They reported that
the electrical conductivity is enhanced greatly, while maintaining high optical
transmittance in the range of 400 nm-2500 nm. The polycrystalline films posses’
tetragonal rutile structure of SnO2 and higher-valence element tungsten doping can
efficiently improve the conductivity of the thin film. The lowest resistivity of 6.67x104
ohm cm was reproducibly obtained and the average optical transmittance is 86% in the
visible region and about 85% in the near-Infra red region [35].
2.6.8 Transition metal codoped SnO2
S. J. Liu et al. reported that the Zn and Mn codoped SnO2 films shows room –
temperature ferromagnetism. They deduced that the valence of Mn ions in the (Mn, Zn)
codoped SnO2 films is 2+ where as Mn doped SnO2, the valence of Mn ions is 3+. The
difficult in the valence state due to the codoping of Zn. The Mn doped SnO2 films
exhibit diamagnetic and paramagnetic respectively. By introducing a few percent of Zn
(5at %) the codoped films exhibit stable and hysteretic ferromagnetism at room
temperature. It is believed that the magnetic moments in the codoped SnO2 films come
from the Mn2+ions whose numbers of unpaired 3d electrons (3d5) and ion radius (80pm)
are both larger than those of Mn3+ ions (3d4 and 58pm).Thus more unpaired electron and
larger ion radius would increase the probability of ferromagnetic coupling between
magnetic Mn2+ ions. On the other hand, defects resulted from Zn doping and delocalized
carriers originated from oxygen vacancies also enhance the ferromagnetic interaction. It
can be seen that the saturated magnetization decrease with decreasing the carrier
38
Chapter II
Literature Review
concentration. This result reveals that the itinerant carrier plays a dominant role in the
ferromagnetic interaction between the magnetic Mn2+ ions.
The dependence of the
saturated magnetization on the carrier concentration is consistent with the carrier –
mediated model [13]. Xing Li et al. (2007) observed room-temperature ferromagnetism
in Fe and Ni-codoped In2O3 samples. (InO.9 Fe
o.1-x
Ni x)2 O3, O< x < 0.1 samples were
prepared by citric acid sol-gel method. All of the samples with intermediate x values are
ferromagnetic at room- temperature. The Fe-doped sample is found to be weakly
ferromagnetic, while the Ni-doped sample is paramagnetic, but the co-doped sample has
enhanced room-temperature ferromagnetism. The magnetic and magneto-transport
including Hall effects studies on the samples indicate the observed ferromagnetism is
intrinsic rather than from the secondary impurity phase [23]. Rezq naji Aljawfi et al.
(2011) reported the properties of Co/Ni codoped ZnO based nanocrystalline DMS. There
is no change of lattice parameters indicates that the ionic radius of Co2+,Ni2+ and Zn2+ are
more or less same and the XRD, XPS spectra shows the incorporation of the Co2+ and
Ni2+ ions into the ZnO lattice and replace Zn2+ sites without changing the wurtzite
structure. From the UV spectra the absorption edge of sample shifts towards lower
energy. This is mainly due to the Co2+ ions got incorporated into the lattice structure of
ZnO. From the FTIR spectra, the gradual shift in the absorption frequency with different
Co/Ni doping is caused by the difference in the bond lengths that occurs when Co2+ and
Ni2+ ions replace Zn2+ ions. And confirm the incorporation of Co and Ni into ZnO lattice
structure. The RTFM has been presented in terms of vacancy in the frame of BPM
model, where the Co ions at surface of the nanoparticles can be ferromagnetically
coupled and mediated by oxygen vacancies. Because of low carrier concentration in the
system, they cannot use the RKKY theory to explain the RTFM [25].
39
Chapter II
Literature Review
Reference
[1]. Jacob Millman, “Integrated electronics,” Tata McGraw-Hill edition, New Delhi
(1999).
[2]. E.M. Conwell, Phys. Rev, 103, 51 (1956).
[3]. N.F. Mott, J.Phys. 34, 1356 (1956).
[4]. A.Miller and E.Abrahams, Phys.Rev. 120, 745 (1960).
[5]. N.F.Mott, Phil.Mag. 19, 835 (1969).
[6]. C.L.Chien and C.R.Westgate, “The Hall effects and its application,” New York:
Plenum (1980).
[7]. Saxena , Gupta, Saxena, “Fundamentals of solid (2004).
[8]. K.H.J.Buschow and F.R.De Boer, “Physics of magnetism and magnetic materials,”
Kluwer Acadecmic/Plenum publishers, New York (2003).
[9]. K.K.Tewari, “Electricity and magnetism,” S.Chand& Company Ltd, New Delhi
(1998).
[10]. Charles Kittel, “Introduction to solid state physics,” John wiley &sons, New York
(2000).
[11]. Nan zheng, “Introduction to dilute magnetic semiconductors,” University of
Tennessee Knoxville (2008).
[12]. S.A.Wolf, J.Supercond.13, 195 (2000).
[13]. S.J.Liu, C.Y.Liu, J.Y.Juang and H.W.Fang, J.of .App.Phy. 105,013928 (2009).
[14]. R.Mattan, J.M.George, H.Jaffres, F.Nguyen Van Dua, A.Fert, B.Lepine ,
A.Guivarc’h, and G.Jezequel, Phys,Rev, Lett,90,166601 (2003).
[15]. Subhasb C.Kashyap, K.Gopinadhan, D.K.Pandya, Sujeet chaudhary, JMMM.
321,957-962 (2009).
40
Chapter II
Literature Review
[16]. Ajay kaushal, Prerna Bansal, Ritu vishnoi, Nitin chaudhary, Davinder Kaur, Physica
B.404, 3732-3738 (2009).
[17]. Nguyen Hoa Hong, Antoine Ruyter, W.Prellier, Joe sakai and Ngo Thu Huong,
J.Phys.: condens, Matter.17 ,5633-6538 (2005).
[18]. Junying Zhang, Qu Yun, Qianghong Wang, Mod.Appl.Sci.4,11 (2010).
[19]. J.Dubowik, I.Goscianska, Y.V.Kudryavtsev, A.Szlaferek,JMMM. 310, 2773-2775
(2007).
[20]. Sang –Do Han, Hua Yang , Li Wang , Jong- won Kim, Sensors and Actuatirs B.
66,112-115 (2000).
[21]. Li Zhang, Shihui Ge, Yalu zuo ,Xueyun Zhou, Yuhua Xiao, Shiming Yan, Xiufeng
Han, Zhenchao wen, J.Appl.Phys. 104, 123909 (2008).
[22]. Nguyen Hoa Hong, Joe Sakai, Physica B. 358,265-268 (2005).
[23]. Xing Li, Changtai Xia, Guangqing Pei, Xiaoli He, JPCS. 68,1836-1840 (2007).
[24]. N.Tahir, S.T.Hussain, M.Usman, S.K.Hasanain ,A.Mumtaz, Appl.Sur. Sci. 255,
8506-8510 (2009).
[25]. Rezq Naji Aljawfi, S.Mollah, JMMM. 323,3126-1332 (2011).
[26]. Aditya Sharma , Mayora Varshney, Shalendra kumar, K.D.Verma and Ravi Kumar,
Nanometer.nanotechnol. 1(1),29-33 (2011).
[27]. K.Srinivas, M.Vithal, B.Sreedhar, M.Manivel Raja and P.Venugopal Reddy,
P.PHYS. Chem.C. 113.3543-3552 (2009).
[28]. Weibing chen and Jingbo Li, J. Appl. Phys. 109,083930 (2011).
[29]. Liu Chun –Ming ,Fang Li- Mei, Zu Xiao-Tao and Zhou Wei-Lie, Chinese Physics.
16,1,95-99 (2007).
41
Chapter II
Literature Review
[30]. X.L. Wang , Z.X.Dai and Z.Zeng, J.Phys.: Condens. Matter. 20, 045214 (2008).
[31]. C.M.Liu, L.M. Fang, X.T.Zu and W.L.Zhou, Phys.Scr. 80, 065703 (2009).
[32]. Susmita Kundu, Dipten Bhattacharya, Jitten Ghosh, Pintu Das, Prasanta K.Biswas,
Chemical .Phys.Lett. 469, 313-317 (2009).
[33]. C.B.Fitzgerald, M.Venkaesan, A.P.douvalis, S.Huber and J.M.D.Coey, J.Appl.Phys.
95 (2) (2004).
[34]. Z.M.Tian, S.L.Yuan, J.H.He, P.Li, S.Q.Zhang , C.H.Wang, Y.Q.Wang, S.Y.Yin,
L.Liu, J.All.Com. 466, 26-30 (2008).
[35]. YanWei Huang, Guifeng Li,Jiahan Feng, Qun Zhang, Thin solid films. 26409
(2009).
[36]. L.M.Fang, X.T.Zu, Z.J.Li, S.Zhu, C.M.Liu, W.L.Zhou. L.M.Wang, J.All.Com. 261267 (2008).
[37].Yanxue Chen, Jun Jiao, IJMPB. 23,6&7,1904-1909 (2009).
[38].B.Sathyaseelan, K.Senthilnathan, T.Alagesan,R.Jayavel, K.Sivakumar,
J.Mat.Che.Phys. 124,1046-1050 (2010).
[39].H.Kimura T.Fukumara, H.Koinuma, M.Kawasaki, Physica E. 10,265-267 (2001).
[40]. Feng Gu, Shu Fen Wang, Meng Kai Kii, Yong Xin Qi, Guang Jun Zhou, Ding Xu,
Duo Rong Yuan , Inorg.Che. Comm. 6, 882-885 (2003).
[41]. A.J.Nadolny, J.Sadowski, B.Taliashvili, M.Arciszewska, W.Dobrowoslski,
V.Domukhovski, E.Lusakowska, A.Mycielski, V.Osinniy, T.Story, K.Swiatek, R.R
Glazka, R.Diduszko, JMMM. 248,134-141 (2002).
[42]. Y.W.Heo, J.Kelly, D.P.Norton, A.F.Hebard, S.J.Pearton, J.M.Zavada, L.A.Boatner,
Ele.Chem and Solid- State Lett. 7,12,309-312 (2004).
42
Chapter II
Literature Review
[43]. C.Van Komen , A.Thurbe, K.M .Reddy, J.Hays, A.Punnoose, J.Appl.Phys.
103,07D141 (2008).
[44]. Xiaghong Liu, Jun Zhang, Xianzhi Guo, Shihua Wu, Shurong Wang, Sensors and
Actuators B.(2011).
[45]. J.M.D.Coey, A.P.Douvalis, C.B.Fitzgerald, M.Venkatesan, Appl.Phys.Lett.
84,8,1332-1334 (2003).
[46]. Pradip KR Kalita, B.K.Sarma, H.L.Das, Bull.Mater.Sci. 23,313-317 (2000).
[47]. S.Sambasivam, Byung Chun choi, J.G.Lin, JSSC. 184,199-203 (2011).
[48]. Junko Sakumar, Kiyoshi Nomura , Cesar Barrero, Masuo Takeda, Thin solid films.
515,8653-8655 (2007).
[49]. Khaled Melghit, Khalid Bouziance, Mat.Sci and Engi B. 128, 58-62 (2006).
[50]. K.Gopinadhan, Subhash C .Kashyap, Dinesh K.Pandya and Sujeet Chaudhary,
J.Appl.Phys. 102,113513 (2007).
[51]. C.B.Fitzgerald, M.Venkatesan, L.S. Dorneless, R.Gunnung, P.Stamenov, and
J.M.D.Coey, Phys.Rev. B. 74,115307 (2006).
[52]. X.F.Liu ,W.M.Gong, Javed lqbal ,B.He, R.H.Yu, Thin solid films. 517, 6091-6095
(2009).
[53]. Hong Xia Wang, Yu Yan , Y.S.H.Mohammed ,Xiaobo Du, Kai Li, Hanmin Jin,
JMMM. 321,337-342 (2009).
[54].L.Yan,J.S.Pan ,C.K.Ong,MSEB.128,34-36 (2006).
[55]. Hong Xia Wang, Yu Yan ,Y.Sh. Mahammmed, Xiaobo Du, Kai Li, Hanmin Jin,
JMMS. 321,3114-3119 (2009).
[56]. Yongbin Xu, Yongjun Tang, Chuanjun Li, Guanghui Cao, Welin Ren, Hui Xu,
Zhongming Ren, J.All. Com. 481,837-840 (2009).
43
Chapter II
Literature Review
[57]. T.Krishnakumar, R.Jayaprakash, M.Parthibavarman, A.R.Phani, V.N.Singh,
B.R.mehta, Mat.Lett. 63,896-898 (2009).
[58]. P.S.Patil, S.B.Sadale, S.H.Mujawar, P.S.Shinde, P.S.Cjigare, App.Su.Sci, 253,685085567 (2007).
[59]. M.Parthibavarman, V.HariHaran, C.Sekar, V.N.Singh, J.Opto.Ele. &. Adv.Mat. 12,
9, 1894-1898 (2010).
[60]. A.Bouaine and N.Brihi, G.Schmerber, C.Ulhaq-Bouillet, S.Colis, A.Dinia,
J.Phys.Chem. 111, 2924-2928 (2007).
44