Chapter 1 Introduction [601886]

Chapter 1 Introduction

1
Chapter 1 Introduction
1.1 Superconductivity Discovery
Generally speaking, the discovery of superconductivity has been recognized as one of
the greatest scientific achievements of the twentieth century. Superconductivity
phenomenon can be defined as the state at which the current can flow in a material without
noticeable energy dissipation . It is accompanied by a sudden drop of the electrical
resistance to zero by cooling below a critic al temperature called the superconducting
transition temperature Tc, see Figure 1 .1(a). Superconductivity was first discovered in
mercury by Onnes [1] in 1911. The temperature at which mercury becomes
superconducting was found to be very close to the boili ng point of liquid helium at 4.2 K.
Subsequently , superconductivity was discovered in many metals, alloys and intermetallic
compounds. Besides the lack of resistance, another characteristic of superconductivity is
the "Meissner effect", discovered by Meissner and Ochsenfeld in 1933 [2]. It is found that
the superconductor expels the magnet ic field from inside it, while the ideal conductor
maintains its interior field . Thus, superconductor s are not just perfect conductor s but also
perfect diamagnet ism. This is called the Meissner effect as shown in Figure 1.1(b).

(a) (b)
Figure ‎1.1 Superconductor properties (a) Zero resistivity, (b) Meissner effec t [3].
As a matter of fact, t he superconductor materials that totally exclude an applied
magnetic flux are known as type -I superconductors and the field at which this happens is
called the thermodynamic critical magnetic field B c (see Figure 1.2a ). Moreover, other
superconductors, called type-II superconductors, are also perfect conductors but their
magnetic properties are more complicated. This type of superconductors has two different
critical magnetic fields denoted by B c1 and B c2, where B c1 and B c2 are the lower and the
upper critical magnetic fields , respectively. Flux is totally excluded when the applied
magnetic field is below B c1, partially excluded in the range from B c1 to B c2, and the
material becomes normal above B c2. The state between B c1 and B c2 is called the vortex or
mixed s tate, sometimes called "fluxons" because the flux carried by these vortices is
quantized (see Figure 1.2 b). Inside the vortex , there is no superconducting state, whereas
outside the vortex there is. Upon increasing the extern al magnetic field even more, the
vortices overlap , and the material becomes normal.

Chapter 1 Introduction

2

Figure ‎1.2 Phases diagrams of (a) type -I and (b) type -II superconductors.
1.2 Historical Background
During the first 75 years of superconductivity research , the highest T c of
conventional (elements and alloys) superconductors was still below 30 K . However, the
highest T c known was limited to 23.2 K in the Nb 3Ge alloy [ 4]. This circumstance impeded
the practical applications of s uperconductors in tremendous degree due to the high cost of
liquid helium and difficulties in its preparation. The situation changed dramatically in 1986
with the discovery of the so -called high temperature supercondu ctors (HTSC s) in
nontraditional compoun ds, namely cuprates. The distinctive feature of HTSC s is that all
these compounds have atomic CuO 2 -plane, playing the key role for the origin of
superconductivity. Actually, the first HTSC was discovered by Bednorz and Müller [ 5] at
Tc = 30 K in La -Ba-Cu-O compound. The value of T c in La2-xBaxCuO 4 was found to
increase up to 57 K with the application of high pressure [6]. This observation in
La2-xBaxCuO 4 material raised the hope of attaining even higher transition temperatures in
cuprate oxides. Hence, Wu et al. [ 7] reported a remarkably high superconducting transition
temperature of 92 K with replacing La3+ ions by Y3+ ions in nominal composition
Y1.2Ba0.8CuO 4-δ. Later, different groups [ 8-10] ide ntified that the superconducting phase
responsible for 90 K has the composition YBa 2Cu3O7-δ (YBCO). Consequently, t he
discovery of superconductivity above the boiling point of liquid nitrogen led to an
extensive search for new superconducting materials. S uperconductivity at transition
temperatures of 105 K in the multiphase sample of the Bi–Sr–Ca–Cu–O (BSCCO)
compound was reported by Maeda et al. [1 1] in 1988. The highest T c of 110 K was
obtained in the Bi –Sr–Ca–Cu–O compound for Bi2Sr2Ca2Cu3O10+δ composit ion [ 12].
Sheng et al. [13] substituted the nonmagnetic trivalent Tl for R E in RE-123, where R E is
a rare -earth element. By reducing the reaction time to a few minutes so as to overcome the
high-volatility problem associated with Tl 2O3, they detected superconductivity above 90 K
in TlBa 2Cu3Ox samples in 1987. By partially substituting Ba2+ ions by Ca2+ ions, Sheng et
al. [14] discovered a T c ~ 120 K in the multiphase sample of Tl -Ba-Ca-Cu-O (TBCCO)
compound in 1988. The superconducting transition tem perature of 125 K for
Tl2Ba2Ca2Cu3O10+δ composition was reported as the h ighest record in T c among TBCCO
compound s [15]. In September 1992, Putillin et al. [16] found that the HgBa 2CuO δ (Hg-
1201) compound with only one CuO 2-plane showed a T c up to 94 K. Furthermore, in April
1993, Schilling et al. [17] reported the detection of superconductivity at temperatures up to

Chapter 1 Introduction

3
133 K in HgBa 2Ca2Cu3O8+δ, and its T c increased to reach 153 K under applied pressure of
31 GPa [18]. Transition temperature has been found to increase as the number of CuO 2-
planes increase to three in Bi –Sr–Ca–Cu–O, Tl –Ba–Ca–Cu–O, and Hg –Ba–Ca–Cu–O
compounds. The discovery of superconductivity in the cuprates was surprising and
exciting, not simply because of the large increase in T c, but also because no previous oxide
superconductors had ever been found. Furthermore, in their stoichiometric form (with no
additional oxygen or other dopant atoms added), these materials are antiferromagnetic
Mott insulators. It is commonly accepted that magnetism cannot coexist with
superconductivity. For example, Abrikosov and Gor’kov showed that magnetic impurities
disrupt superconductivity and depress T c [19].
However, a huge surprise came in 2001 with the discovery of superconductivity up to
40 K in the intermetallic MgB 2 phase by Nagamatsu et al. [ 20]. One of the most important
features of MgB 2 is that it does not exhibit weak link electromagnetic behavior at grain
boundaries or fast flux creep . Therefore, this limits the performances of YBCO cuprates,
oxidized Nb, Nb 3Sn and Nb films [21]. In January 2008 , a new family of iron-based
superconductor s " Pnictide s" has been discovered with a maximum Tc up to 55 K for
SmFeAsO 1-xFx [22]. In contrast to the cuprate HTSC s, the parent compounds are here
metals with a spin -density wave long -range magnetic order. Recently, Tc values up to 10 9
K were claimed to be a monolayer of FeSe on a SrTiO 3 substrate [23]. In 2015, hydrogen
sulfide ( H2S) has been observed to exhibit superconductivity at 203 K but at extremely
high pressures 100 GPa [24]. The driving force in the research for new materials is
based on the desire to find superconductivity at room temperature , which would have
enormous effect on the use of this phenomenon in a variety o f applications.

Figure ‎1.3 History of superconductivity discovery .
020406080100120140160180200
1900 1920 1940 1960 1980 2000 2020HgPb NbNbN0.96 Nb3Sn
Nb3(Al0.75Ge0.25)Nb3GLa BaCuO4YBa2Cu3O7Bi2Sr2Ca2Cu3O10Tl2Ba2Ca2Cu3O10HgBa2Ca2Cu3O8
MgB2SmFeAsO1-xFx77K LN2
YearTc (K)H2S
FeSe
4.2K LHe

Chapter 1 Introduction

4
1.3 Properties of Sm -123 Superconducting Phase
Since the discovery of high -temperature superconductors in the La -Ba-Cu-O system
[5], a number of studies have been focused on developing higher Tc superconductors.
YBa2Cu3O7-δ is the first high Tc superconductor that shows superconductivity above 77K
[8]. The Y-site can be completely replaced by 4f magnetic ions, lanthanide series, except
Pm [25], Pr [26, 27], Ce and Tb [28, 29].The elements of these series are either light rare
earth (LRE) elements; Nd, Sm, Eu and Gd, or heavy rare earth (HRE) elements; Yb, Tm,
Er, Ho and Dy. The decrease in the ionic radius of these magnetic ions causes almost linear
decrease of the lattice parameters and the separation distance between the CuO2 planes
[30]. However, T c does not depend heavily on the rare -earth [ 31]. This is attributed to the
reason that the superconducting electrons in the CuO2-planes do not experience spin flip
scattering -off of the far sitting magnetic rare earth moments [32]. Due to the close ionic
radius between LRE and Ba ions, the phenomenon of LRE -Ba substitution is present [ 33].
This phenomenon result s in the formation of LRE1+xBa2−xCu3Oy solid solutions. The
possibility of an interchange between Ba2+
and RE3+ ions depends on the irrelative ionic
radii. The more similar are Ba2+
and RE3+
radii, the easier would be the interchange between
them. The LRE -Ba substitutions cause reduction in Tc and broadening in the
superconducting state transition ΔT= (T c-To) [34]. On the other hand , the substituted
phases may work as pinning centers under magnetic field [ 35]. Further, it is known that
synthesis at reduced oxygen partial pressures or addition of Ba -excess (in the form of such
precursors as BaO, BaO 2, BaCuO x etc.) can effectively suppress the substitution and
therefore the formation of solid solutions of the LRE 1+xBa2-xCu3Oy type. [36]. Manka et al.
[37] studied light non -stoichiometry effects in Sm -Ba-Cu-O sintered samples of two series
of single -phase Sm 1+xBa2-xCu3O7+δ and Sm 1-xBa2+xCu3O7+δ samples with x from 0 to 0.1.
For Ba excess x in the Sm 1-xBa2+xCu3O7+δ samples suppresses the substitution of Sm3+ for
Ba2+ and thereby reduces degradation of superconducting properties, such as Tc, ΔT, and
the maximal value of volume magnetization when compared to Sm excess x in Sm 1+xBa2-
xCu3O7+δ samples at higher values of x. However, the Sm excess sample with composition
deviation x = 0.04 shows the highest hysteresis of ΔM from all the studied samples at
higher values of the applied magnetic field. Oh et al. [ 38] reported that higher critical
current densities, Jc, of Sm -Ba-Cu-O coated conductors were obtained in the Sm-rich, Ba –
poor and Cu -rich composition regions, as compar ed to the stoichiometric region .
In recent years, a lot of attention has been paid to SmBa 2Cu3O7-δ partly due to its
higher T c, and mainly due to its higher critical current density as co mpared with that of Y –
123, both in low and high magnetic fields [ 39]. Actually, Sm-123 superconductor can trap
a magnetic field of 2.1 T at 77 K and can increase its maximum trapped field to 9 T at 25 K
in the sta tic field cooling magnetization [40]. Theref ore, Sm-123 is potentially attractive
for high field power applications and the synthesis of superconducting bulks, tapes, and
wires [41].

Chapter 1 Introduction

5
1.4 Crystal Structure of SmBaCu 3O7-δ
In general, the structure of HTSCs is commonly related to perovskite structure. The
unit cell of perovskite consists of two metal atoms (A and B) and three oxygen atoms, with
the general formula given as ABO 3. The ideal perovskite structure is shown in Figure 1.5.
Atom A, sitting at the body -centered site, is coordinated by 12 oxygen atoms. On the other
hand , atom B occupies the corner site and the oxygen atom occupies the edge -centered
position [42].

Figure ‎1.4 The perovskite structure ABO 3 [42].
Additionally , the crystal structure of SmBa 2Cu3O7-δ as shown in Figure 1.6 can be
described as the superposition of three oxygen -deficient perovskite cells (Sm is the central
ion of central cell, Ba is th e central ion of external cells) in the layered sequence BaO –
CuO 2-Sm-CuO 2-BaO -CuO, SmBa 2Cu3O7-δ contains two CuO 2 planes per unit cell
separated by a Sm atom. The CuO chain s lie between the BaO layers . These CuO chains
not only lead to an in -plane orthorhomb ic structural distortion but also give rise to a
pronounced electronic and magnetic in plane anisotropy. Moreover, t he CuO 2 planes
consist of CuO 5 pyramids whose bases are connected via their corner oxygens to form the
CuO 2 planes. Via the so -called apical oxygens, they are connected with the CuO chains
that extend along the b -axis direction.

Figure ‎1.5 Crystal structure of Sm -123 [43].

Chapter 1 Introduction

6
1.5 Oxygen Stoichiometry of Sm -123
One of the most interesting properties of the (LRE)Ba 2Cu3O7-δ compound is its
ability to support a large oxygen non -stoichiometry. In fact, its oxygen content can be
varied from six (δ = 1) to seven (δ = 0) in a continuous way. The oxygen deficiency occurs
in the CuO chains, which co ntain copper Cu(1) and oxygen atoms O(1).The chain plane
acts as a charge reservoir for the CuO 2 double planes. Actually, a nnealing at temperature
above 300 oC is often a convenient way to control charge density. The structure of
SmBa 2Cu3O7-δ depends on the oxygen content δ [44] as shown in Figure 1.7 . With δ=1
the O(4) and O(5) in the Cu(1) layer are vacant and the structure is tetragonal (space group
P4/mmm) . O(4) and O(5) sites are equivalent and also the O(2) and O(3) sites. The
tetragon al phase of SmBa 2Cu3O7-δ (δ > 0.5) is insulating and does not superconduct.
However, i ncreasing the oxygen content slightly causes more of O(4) and O(5) sites to
become occupied, with occupational factor equal to δ/2. For δ < 0.5, Cu-O chains along the
b-axis of the crystal are formed, O(5) sites are vacant and O(4) sites are occupied with
occupational factor equal to δ. Elongation of the b -axis changes the structure to
orthorhombic (space group Pmmm) and it undergoes a tetragonal to orthorhombic phase
transition. The orthorhombic phase of SmBa 2Cu3O7-δ is superconducting , and the optimum
superconducting properties occur when δ =0.93 and all of the O(4) sites are occupied with
few vacancies. Charge transport is confined to the Cu(2)O planes while the Cu(1)O(4)
chains act as charge reservoirs, which provide carriers to the CuO planes.

Figure ‎1.6 Sketches of the SmBa 2Cu 3O7-δ unit cells, for the non superconducting tetragonal
phase, δ > 0.5 (a), and for the superconducting orthorhombic phase, δ < 0.5 (b) [44].

Chapter 1 Introduction

7
1.6 Phase Diagram of Sm -123
For the Sm -123 system, the phase diagram at 950°C in the Sm 2O3-BaO-CuO system
made by the quenching method was reported , and the Sm -123 phase is considered to
crystallize in a triangle of (Sm-123)-BaCuO 2-CuO [ 45]. To determine precise liquidus
lines, where primarily the Sm -123 phase crystallizes, seven pseudo -binary phase diagrams
in the region of (Sm-123)-BaCuO 2-CuO were made [46]. Figure 1.8 shows a constructed
pseudo -binary phase diagram by using (Sm-123)–Ba7Cu18O25 pseudo -binary system as an
example [47]. It shows that the Sm -123 phase primarily crystallizes in the composition
range from 5 mol% to 40 mol%. Above 45 mol% concentrati on, the peritectic reaction [48]
Sm-123 → Sm -211 + liquid occurs, and the Sm -123 crystals decompose at about 1060°C
to form the Sm -211 phase. The liquidus line of the Sm -211 phase lies above 1080°C. Three
phases Sm -123, Sm -211, and the liquid coexist in the composition range between 45 mol%
and 90 mol%, and in the temperature range between 1060°C and 1080°C. As temperature
increases in this region, the Sm -123 crystal dissolves but the Sm211 grows. Conversely, as
temperature decreases, the Sm -211 dissolves and the Sm123 grows. These processes were
directly observed.

Figure ‎1.8 Pseudo -binary phase diagram of Sm -123 system in air [46].

1.7 Nanoparticles Spinel Ferrite
The cubic spinel ferrite has the general formula, MeFe2O4 (Me=Fe2+, Co2+, Mn2+,
Ni2+, Zn2+, etc). These ferrites crystallize in the spinel structure. The spinel lattice is
composed of a close packed oxygen anion arrangement in which 32 oxygen ions form the
unit cell. These anions are packed in a face -centered cubic (FCC) arrangement, leaving two
kinds of spaces between anions. The first one is tetrahedrally coordinated sites (A sites) ,
surrounded by four nearest oxygen atoms while the second one is octahedrally coordinated
sites (B sites) , surrounded by six nearest neighbor oxygen atoms, as shown in Figure 1.4.
In the unit cell of 32 oxygen ions , there are 64 tetrahedral sites , only 8 are occupied and out
of 32 octahedral sites , only 16 are occupied, resulting in a structure that is electrically
neutral [ 48]. The localization of ions either in A or B sites depends fundamentally on the
ion and lattice sizes. In general, divalent ions are larger than trivalent ions. This is because

Chapter 1 Introduction

8
trivalent ion nuclei produce a greater electrostatic attraction ; hence , their electron orbits
contract. The octahedral sites are larger than the tetrahedral sites; thus, the divalent ions are
localized in the octahedral sites (+2)×8= +16 whereas the trivalent ions are in the
tetrahedral sites (+3)×16=+48 [49] or a total of +64 which is needed to balance the
(-2)×32= -64 for the oxygen ions. Due to the exchange interaction forces, the magnetic
moments of cations in the A sites are aligned parallel with respect to each other and the
magnetic moments of the cations in B sites align parallel to each other. However, the
arrangement of the magnetic moments between the cations in A and B sites are
antiparallel , as shown in Figure 1.4. Due to the difference in the number of cations in A
and B sites, there is a net magnetic moment and therefore spinel ferrite is a ferrimagnetic
material [50]. In reality, the magnetic moment ordering is more complicated than the ideal
situation above because cationic distribution over A and B sites varies with compounds,
and synthesis conditions such as temperature, pressure, and starting metal precursors.

Figure ‎1.7 Schematic diagram for two subcells of a unit cell of the spinel structure, showing
tetrahe dral and octahedral sites, A and B respectively [50].
The structural formula for a generic spinel compound M eFe2O4 can be written as
[Me1-iFei]A[MeiFe(2-i)]BO4, where the amounts in brackets represent the average occupancy
of A sites and B sites and i is the inversion parameter. Depending on cation distribution,
the spinel structures can be normal, inverse or partially inverse. In the case of the normal
spinel, the divalent ions are all on A sites and trivalent ions occupy B sites. ZnFe 2O4 is a
good example of a normal spinel ferrite with the composition of [Zn2+]A[Fe 2 2+]BO4 where
i=0. Zn2+ has no unpaired d electron and therefore should have no contribu tion to the
magnetic moment of the system. Since all the Zn2+ cation occupies the tetrahedral site and
all the Fe3+ occupies octahedral sites, instead of antiferromagnetic ordering between A
and B sites, antiferromagnetic ordering within B site is exp ected. Thus, net magnetic
ordering is expected to be antiferromagnetic. However, careful experiments prove that
there is no antiferromagnetic ordering between Fe3+ ions [ 51]. In the inverse spinel, i=1,
the divalent ions occupying only B sites while trival ent ions are located on both A and B
sites in equal proportion. If the divalent cations are present on both tetrahedral and
octahedral sites, the spinel is partially inverted and 0 < i < 1. Previously, it was found that
80% of Mn2+ occupies the A sites and 20% may go to the (B) sites, which makes it a
partially inverse spinel with the composition of [Mn 1−i 2+ Fe i3+ ]A[Fe 2−i3+ Mn i2+]BO4, where i
=0.20 [52].

Chapter 1 Introduction

9
1.8 Effect of Chemical Substitution on RE -123 Phases
The sup erconducting properties of LRE -123 system can be controlled via chemical
substitution and/or changing the oxygen stoichiometry. In principle, non-isovalent
substitutions in LRE -123 system offer a powerful tool to vary the charge concentration
CuO 2-plane, and therefore, to study the behavior of structural and superconducting
properties with varying carrier concentration. In effect, t his process depends on the
substituting elements that having different ionic radii and bonding character. Similarly , the
superconducting properties are either enhanced or destroyed, depending on the
characteristics of the dopant in the crystal structure. As a matter of fact, t he substitution
process has a limit named the critical solubility limit over which impurities begin to form
[53]. Substitutions in the rare earth position, Ba site and the two Cu sites are described in
some detail as follows:
1.8.1 Substitution in R Site
The superconducting properties are strongly influenced by partial substitution of
Ca2+ on the (RE)3+ site in RE -123 systems. In Y 1- xCaxBa2Cu3 O7−δ, the solubility limit for
Ca is just below 20%. Above this level , the BaCuO 2−δ impurity phase is formed [54,55].
Tokura et al. [ 56] found that the sample with x = 0.2 0 display a maximum T c at
intermediate oxygen concentrations, followed by a decrease of T c. Polycrystalline samples
of Gd 1- xPrxBa2Cu3O7−δ for 0.0 ≤ x ≤ 1.0 were investigated by Akhavan [ 57]. In fact, the
electrical resistivity measurements showed the suppression of superconductivity with
incre asing x until for x = 0.45 the superconductivity vanished. Wen et al. [58] studied
superconductiv ity and crystalline structure of SmBa 2Cu3O7-δ doped with Pr and Ca. It
has been found that T c decreases from 91.5 K to zero when doping concentration x in
Sm 1−2xCaxPrxBa2Cu3O7−δ increases from zero to 0.4 . This is due to Pr may have an effect
of pair breaking in addition to hole -filling. On the other hand , Ca may have an effect of
hole creation in the CuO sheets as x < 0. 25. When x < 0.25, Ca and Pr substitute for Sm in
the same fraction . The solubility limit of Ca in orthorhombic phase is about x=0.25.

1.8.2 Substitution in Ba Site
For Sm(Ba2−xPrx)Cu3O7-δ polycrystalline samples with 0.0 ≤ x ≤ 0.4 were
investigated by Colonescu et al. [59]. The results showed that the Tc decreased and an
orthorhombic –tetragonal (O /T) transition for x=0.35 which corresponds to the O(5) (0.5,
0.0, 0.0) occupation due to the appearance of Pr3+ at the Ba2+ site, and requires a more
negative charge . The neutron diffraction data of Nd(Ba 2−xPrx)Cu 3O7 has also indicated that
when Pr occupies the Ba site, it results in the (O/T) transition . Furthermore, it behaves like
all other trivalent rare -earth ions [60]. The substit ution of Sr in (RE)Ba2−xSrxCu3O7-δ, did
not change the oxygen content significantly and Tc decreased linearly [ 61–63]. The
substitution of La in the Ba site added electrons into the R -123 system. For
YBa2−xLaxCu3O7-δ, Tc first increased with x by a about 2 K at small values of x and then
decreased at higher doping levels [64,65].

Chapter 1 Introduction

10
1.8.3 Substitution in Cu Site
The chemical substitution effects at copper had shown that the copper is very essential
for understanding the superconductivity [ 66]. Sm-123phase has two non -equivalent Cu
sites, the linear chain Cu(1) in the O(1) -Cu(1) -O(1) units and the planar Cu(2) in the CuO 2
planes containing O(2) and O(3). It is believed That the superconductivity of this system
takes place in the CuO 2 planes through hol e charge carriers, while the oxygen content of
CuO chains governs the carries concentration (hole doping) in the CuO 2 planes [ 67, 68].
The presence of any modification in the CuO 2 layers strongly affects the electronic
structure and the density and mobilit y of the charge carriers and thus the superconductivity
is suppresse d [69, 70]. Xue et al. [ 70–72] studied the effect of partial substitution of Cu2+
ions by Zn2+ and Fe3+ ions in Sm -123 phase. They found that for SmBa 2Cu3−xFexO7−δ
phase , a structure phase transition from orthorhombic to tetragonal (O /T) in x ≥ 0.1 [70,
72]. Furthermore , the substitution of Cu2+ ions by Fe3+ ions resulted in the localization of
carriers and weakening of the Cu -O chains as carrier reservoir . On the other hand, they
reported that the substitution of Cu2+ ions by Zn2+ ions maintained the orthorhombic
structure for SmBa 2Cu3−xZnxO7−δ phase and suppressed the superconductivity both by
reducing the effective coupling strength and destroying the phase coherence [71,72].
Moreover, they concluded that the superconductivity suppression was stronger by Zn
substitution than Fe substitution in Sm -123 phase. This means that the depression of Tc did
not have direct correlation with the magnetism of Fe3+and Zn2+ions [ 72]. Bara kat et al.
[73] reported that for SmBa 2Cu3−xRuxO7−δ with 0.0 0 ≥ x ≥ 0.5 0, an improvement in the
phase formation and T c for Sm-123 phase with x up to 0.05. For x > 0.05 , the suppression
of phase formation and Tc were observed with (O /T) phase transition around x = 0.50. The
depression in T c could be due to Cooper -pair breaking mechanism as a result of magnetic
ion substitution of Cu2+ by Ru4+ ions which creates a disorder in the internal magnetic
state [ 74,75].
1.9 Importance of Nanosized Particles Addition in HTSCs
For practical applications, especially power applications, HTSCs need to possess a
high critical current density , Jc, under high magnetic fields. Increasing Jc at higher fields
can be accomplished by introducing effective vortex pinning centers. Pinning occurs
naturally via material inhomogeneity and microstructure defects, and can be enhanced
through the addition of artificial pinning centers on the order of magnitude of the
coherence length , ξ [76]. However, other studies showed that the optimum size of pinning
centers should be comparable to the penetration depth, λ , rather than ξ [77]. One of the
effective methods to enhance Jc is the addition of nanoparticles, which introduce strong
flux pinning centers into the bulk superconductors [ 78–83]. Furthermore, t he pinning
strengths of the flux lines in HTSCs can be enhanced by direct magnetic interaction of
vortices with magnetic pinning centers [ 84-87]. For magnetic nanoparticles with average
size d where ξ < d < λ, strong interaction between flux line network and magnetic system
can be expected [88]. Accordingly, t he general features for the effect of nanosized particles
addition on HTSCs have been summarized as follows:

Chapter 1 Introduction

11
1.9.1 For Bi -cuprates
Albiss et al. [89] investigated the effect of nanosized NiO addition on the phase
formation and flux pinning of the polycrystalline (Bi,Pb) -2223.The results show that, with
increasing nanosized NiO addition, the volume fraction of high T c phase (2223) decreased
gradually accompanied with an increase in the low Tc phase (2212). However, almost
no change in the T c for all samples was observed. Vortex pinning forces were found to be
suppressed by nanoparticles addition for all concentrations with x up to 0.005 at all
temperatures. Wan et al. [90] studied the effect of nanosized MgO on (Bi,Pb) -2223 silver –
sheathed. It was found that Jc of samples added with nanosized MgO is higher than that
with zero MgO samples by a factor of 1.2 in zero applied field and about 1.4 in a magnetic
field ranging from 0.1 to 1T at 77 K. Jia et al. [79] investigated the effect of nanosized
ZrO 2 addition on flux pinning capability of the bulk (Bi,Pb) -2223. They concluded that the
nano – ZrO 2 particles have not remarkable effects on the T c. However, the measurements of
Jc showed that appropriate amount of nanosized ZrO 2 addition greatly enhanced the values
of J c under magnetic fields. Annabi et al. [ 91] added nanosized Al 2O3 to (Bi,Pb) -2223
precursor powders during the final sintering cycle o f a multi -step preparation process.
They studied the influence of nanosized Al 2O3 on the melting temperature, phase
formation, microstructure and transport properties of (Bi,Pb) -2223. They reported that the
addition of a small amount of Al2O3 (0.2 wt%) increased Jc at 77K by ~30% and improved
Jc behavior in applied magnetic field either parallel or perpendicular to the sample wide
surface. Abou -Aly et al. [92] examined the effect of nanosized SnO 2 addition with
different conc entrations (x = 0.0 to 2.0 wt %) to the phase formation, microstructure,
electrical and thermal properties of (Bi,Pb) -2223 phase. They investigated that the
addition of nanosized SnO 2 up to x=0.4 wt% increase the volume fraction, Tc and Jc.The
improvement in Jc was believed to be due to the pinning effect of nanosized SnO 2.While,
the higher concentrations of nanosized SnO 2 , x > 0.4 wt% reduced the phase formation, T c
and Jc of (Bi,Pb) -2223 phase. This was attributed to the high concentration of nanosized
SnO 2, induced lar ge agglomerations between superconducting grains , and hence reduced
superconducting grain connectivity and deteriorated intergranular critical current density.
1.9.2 For Tl -cuprates
Sengupta et al. [93] studied melt-processed pellets from Tl-1223 phase with and
without nanosized Al 2O3 addition. They found that nanosized Al 2O3 additions had a
negligible effect on Tc. However, for temperatures ≤ 35K, Jc increased dramatically for all
samples with the nanosized Al2O3 additions. At 77 K, the nanosized Al2O3 additions had a
little effect on J c. It was demonstrated analytically that high concentrations of defects on the
order of the nanosized additions can be effective in pinning flux lines.
1.9.3 For (Cu,Tl) -cuprates
Awad [ 94] investigated the effect of nanosiz ed MgO addition on the phase
formation, microstructure, electrical and mechanical properties of (Cu 0.25Tl0.75)-1234
phase. The nanosized MgO addition up to 0.6 wt% improved the phase formation, grain
connectivity, transport critical current density and microhardness. The improvement in Jc
was believed to be due to the pinning effect of nanosized MgO whereas the increase in
microhardness was due to the increase in grain connectivity and the crack resistance

Chapter 1 Introduction

12
propagation. However, the addition of a larger amount of nanosized MgO > 0.6 wt%
decreased the phase formation, grain connectivity, transport critical current density and
microhardness. This decrease was believed to be due to large amount of nanosized MgO
agglomerated between the grains and hence reduced superconducting grain connectivity
and deteriorated intergranular critical current density. Mumtaz et al. [95] examined the
effect of the incorporation of nanosized CuO in (Cu 0.5Tl0.5)-1223 superconductor matrix
for healing the inter-grain voids, pores and cracks. Elokr et al. [96] prepared nanosized
ZnO and (Cu 0.5Tl0.25Pb0.25)-1223 superconducting phase using the co -precipitation method
and single step solid -state reaction, respectively. The addition of nanosized ZnO up to y =
0.8 wt % improved the volume fraction and enhanced both T c and J c of
(ZnO) y(Cu 0.5Tl0.25Pb0.25)-1223 phase. The melting point of this phase increased from
879°C for y=0 to 893°C at y = 0.4 wt %. The further increase in ZnO addition (y > 0.8 wt
%) reduced the volume fraction, T c, Jc and melting point of (Cu 0.5Tl0.25Pb0.25)-1223 phase.
Abdeen et al. [ 97] studied the influence of nanosized Ag addition on phase formation and
electrical properties of (Cu 0.5Tl0.5)-1223 superco nducting phase. They found that t he low
addition of nanosized Ag up to 1.5wt% enhanced the phase formation and improved the T c,
Jc, and the melting temperature. Nevertheless , for x >1.5wt%, a reverse trend was
observed. Mohammed et al. studied the effect of nanosized SnO 2 [98], nanosized In2O3
[99] and nanosized Fe2O3 [100] on the phase formation, microstructure, electrical and
mechanical properties of (Cu 0.5Tl0.5)-1223 superconducting phase. The analysis revealed
that the volume fraction and Tc increased with nano -SnO 2 addition up to 0.6 wt%, and the
microhardness increased with the increase of nano -SnO 2 addition up to x=1wt%. However ,
nano -In2O3 addition nearly did not alter the percentage of the main phase for all values of
addition, x, and T c is slightly changed with x. The microhardness increased with increasing
the addition of nanosized In2O3 up to 1 wt% . Afterwards , it decreased with further
addition. The addition of nanosized Fe2O3 up to 0.2 wt% increase d the volume fraction and
Tc. On the other hand, the microhardness increased with the increase of nanosized Fe 2O3
addition up to 1wt%.
1.9.4 For Rare Earth (RE -123) Cuprate Superconductors
Xu et al. [ 80] studied the effect of nanosized ZrO2 and ZnO inclusions on
superconductive properties of the melt -processed GdBa2Cu3O7-δ bulk superconductor. The
experimental results clarified that a suitable amount of nanoparticles ZrO2 and ZnO
(x=0.004 mol%) inclusion can enhance the Jc and irreversibility field Birrof the Gd-123
bulk supercon ductors. Chen et al. [101] concluded that the addition of small amounts
(0.004 wt% and 0.4 wt%) of nano -sized Sm 2BaCuO 5 (Sm-211) particles in melt textured
growth Sm -123 can improve both T c and J c values especially at the high field region due to
the enhan cement of flux pinning of nanosized Sm -211 doped samples. Dadras et al. [102].
examined t he effects of carbon nano -tubes (CNTs) on the crystal structure and
superconducting properties of YBa 2Cu3O7-δ (Y-123).They found that T c does not change
much with the CNT doping (91 –92 K) and the electrical links between superconducting
grains were improved to result in the J c increase with doping increase up to 0.7 wt%.
Albiss et. al [103] investigated the effect of nanosized Al2O3 additi on on the microstructure
of polycrystalline Y -123. By adding nano -Al2O3 to the YBCO system, an improvement of
the microstructure occurs. Additionally, the number of weak links decreases, the number of
new pinning centers will increase and an enhancement of the critical current density is
observed with increasing x up to 0.01 wt% addition .

Chapter 1 Introduction

13
1.10 Effect of Nanosized Ferrites Addition on HTSCs
The addition of spinel ferrites nanostructures to the HTSCs has attracted great
interest due to their mechanical hardness, chemical stability, excellent electrical and
magnetic properties [10 4]. Nanosized ZnFe 2O4 and CoFe 2O4 addition s into CuTl -1223 had
healed up the inter -grain voids space and increased the inter -grain coupling [105,106].
Both the XRD and FTIR spec tra show that the presence of nanosized ZnFe 2O4 and
CoFe 2O4 have not altered the structure of unit cell of host (Cu 0.5Tl0.5)-1223
superconductor. Inclusion of nanosized ZnFe2O4 has overall decreased the Tc (0) and
magnitude of diamagnetism whereas nanosized CoFe 2O4 increased T c(0) which could be
most probably due to the improvement of weak -links by the addition of these
nanoparticles . Nanosized NiFe 2O4 was added into (Cu 0.5Tl0.5)Ba 2Ca2Cu3O10 [107], the
results manifested that T c(0) decreased with increasing nanosized concentration . The
sample with x =0.25 wt% revealed an increase in activation energy (U) as compared to
host CuTl -1223 superconductor sample, which indicates a strong flux pinning at x=0.25
wt%. Barakat et al. [ 108] reported that the addition of nanosized Co 0.5Zn0.5Fe2O4 into
(Cu0.5Tl0.5)-1223 up to x=0.08 wt% improved the grain connectivity and Tc. Kong and
Abd-Shukor [78] investigated the effect of nanosized ferrite NiFe 2O4 on the fo rmation of
(Bi,Pb) -2223 and it s superconducting properties, and reported that the addition with 0.01
wt% enhanced both T c and J c. Hafiz et al. [109] studied t he effect of nanosized CoFe 2O4
(60 nm) addition on the J c of (Bi 1.6Pb0.4)Sr2Ca2Cu3O10(CoFe 2O4)x, (x = 0 -0.05 wt%). The
optimal J c was observed in the x = 0.01 wt% pellets. Using this optimal wt%, Ag -sheathed
(Bi 1.6Pb0.4)Sr2Ca2Cu3O10(CoFe 2O4)0.01 superconductor tapes were fabricated using the
powder -in-tube method. The effect of nanosized ferrite CoFe 2O4 on the physical properties
of YBa 2Cu3O7−δ thin film was reported by Wimbush et al.[ 110]. In effect, a reduction in J c
was observed as a result of Y(Fe,Co)O 3 formation instead of CoFe 2O4. Awad et al. [111,
112] investigated the effect of nanosized ZnFe 2O4 and CoFe 2O4 addition on the
superconducting parameters and mechanical properties of GdBa 2Cu3O7-δ superconducting
phase. The analysis revealed that for (ZnFe 2O4)xGdBa 2Cu3O7-δ Tc and Jc enhanced with x
up to 0.06 wt% . However, for (CoFe 2O4)xGdBa 2Cu3O7-δ superconducting samples, the Tc
decreased and Jc enhanced up to x= 0.01 wt%. The microhardness values increased
gradually with increasing the nanosized ZnFe 2O4 and CoFe 2O4 additions .

1.11 Work Strategies
The present work aims to study the effect of nanosized ZnFe2O4 and MnFe2O4
additions on the phase formation, microstructure, electrical, mechanical and magnetic
properties of SmBa2Cu3O7-δ superconducting phase . For this purpose the preparation of
nanosized ZnFe2O4, MnFe2O4 and SmBa2Cu3O7-δ superconducting phase by co-
precipitation method and solid -state reaction technique, respectively was done . The
prepared samples were characterized using X -ray powder diffraction (XRD), scanning
electron microscope (SEM) , proton induced X -ray emission (PIXE) and Rutherford
backscattering spectrometry (RBS) . The superconducting properties of the prepared
samples were investigated using electrical resistivity, transport current density, Vickers
microhardness, Ultrasonic pulse echo te chnique, electron paramagnetic resonance (EPR)
and vibrating sample magnetometer (VSM). Finally, the obtained results are discussed
a n d interpreted with the aid of superconductivity theories and models.