Parinov I.A. Microstructure and Properties of High-Temperature Superconductors

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88 2 Composition Features and HTSC Preparation Techniques

(c)

BPSCCO

MgO

BPSCCO

(a)

(b)

MgO

Fig. 2.30. Microstructure of BPSCCO phase in polished and etched cross-sections

parallel to the hot-pressing direction of (a) a monolithic BPSCO, (b)and(c)a

(MgO)

/ BPSCCO composite with 10 and 20% concentration of whiskers, respec-

tively (all specimens have undergone a three-cycle hot-pressing and annealing) [1186]

2.5 Y(RE )BCO Bulks 89

2.5 Y(RE)BCO Bulks

Intercrystalline boundaries can diminish considerably the transport properties

of YBCO bulks and increase sensitivity to applied magnetic ﬁelds. Therefore,

it is very expedient to create large grain specimens with maximum number of

strongly connected intercrystalline boundaries. As a rule, Y(RE)Ba

7−x

crystallites (123 phase), prepared by the melt-processing technique, include

Y(RE)

BaCuO

particle dispersion (211 phase). The elevated critical current

density in these samples can correlate with the concentration of the 211 inclu-

sions [961] and assumes that ﬁne size of these particles causes directly elevated

pinning. At the same time, the suﬃciently large particles of 211 normal (non-

superconducting) phase obviously diminish the superconducting properties.

Therefore, in order to optimize the ceramic properties, it is necessary to con-

trol the concentration and the size of 211 particles in precursor powder used

in YBCO processing [643]. After fabrication of the precursor powder, it is

pressed into pellets, which is then calcined and melted. The pellets, subjected

to insigniﬁcant pressing, demonstrate a lost density compared with pellets

subjected to cold isostatic pressing. Further sintering causes their densiﬁca-

tion next. However, the hard aggregates of grains, formed during the cold

isostatic pressing, can prevent alignment of some crystallites, and, therefore

limit an eﬀective densiﬁcation during the following sintering. Then, it could be

assumed that the sintering is the optimum process to create high-dense HTSC

samples. It should be noted that 211 particles grow into precursor bulk during

sintering. It is not desired with the view of preservation of ﬁne granularity of

the 211 phase. Therefore, it is necessary to exclude this eﬀect. Thus, the cold

isostatic pressing in some cases may be more practical to prepare optimum

precursor samples, which are then subjected to melting [642].

All techniques, based on melt processing and used to prepare YBCO large-

grain ceramics, are characterized by the peritectic reaction at T

= 1015

◦

owing to the formation of 123 phase from 211 phase and liquid component

[642]:

BaCuO

(211 phase)

+3BaCuO

+2CuO

(liquid phase)

→ 2YBa

6.5

(123 phase)

, (2.4)

or in another form [640]:

BaCuO

(211 phase)

+Ba

6.72

(liquid)

+0.42 O

(gas)

→ 2YBa

6.28

(123 phase)

. (2.5)

Pseudo-binary diagram, showing two peritectic transformations, is pre-

sented in Fig. 2.31. The 211 phase and liquid could be formed by fast heating

a pre-sintered precursor sample to a temperature that is considerably above

. Then 123 phase forms very slowly during the cooling of YBCO partial-

melting through point T

. In this case, 123 phase are added up to 30 wt%

of the 211 particles before the melting process with the aim to create local

pining centers and to prevent liquid loss during melting [641]. Solidiﬁcation

processes, depicted by the formulas (2.4) and (2.5), put deﬁnite requirements

90 2 Composition Features and HTSC Preparation Techniques

Melt

011 + Melt

200 + Melt

211 + Melt

211

200

210

123 + Melt

123 + 211

3BaCuO

+2CuO

YBa

7–x

BaCuO

5YO

1.5

+ BaO

1270°C

1015°C

123 + 011

+ 001

Fig. 2.31. Pseudo-binary diagram along the Y

BaCuO

–YBa

7−x

line show-

ing two peritectic transformations [29]

on YBCO precursor samples [642], namely (i) 211 particles must be suﬃciently

ﬁne in the initial sample precursor in order to form ﬁne-grain dispersion in the

sample fabricated. (ii) The material should be able to retain the liquid phase,

appeared in result of the peritectic reaction into samples at the temperatures,

which are considerably above T

. It is necessary in order to form 123 phase

at cooling. Successive process is dependent on the precursor homogeneity and

density and also on the size distribution of the 211 particles. (iii) The precur-

sor sample should not contain foreign compositional and surface admixtures,

which form the heterogeneous nucleation sites of grains, and hence limit the

grain sizes, that may be reached during the grain growth.

Obviously, it is very possible an enlargement of 211 ﬁne-size inclusions dur-

ing melting. As rule, the 211 enlargement occurs above peritectic temperature,

solely. Therefore, the 211 inclusions with diameter up to 50 μmcanforminthe

ﬁnal ceramic sample. There are other problems, namely a very slow rate of so-

lidiﬁcation, a necessity to control thermal gradients during high-temperature

treatment, a limited size of domains obtained, that is accompanied by their

misorientation, a microcracking formation and heterogeneous composition.

The progress could be reached by changing processing parameters and using

doping additives.

Oxygen release increases sharply during melting. This forms voids, caus-

ing the corresponding rise of the ﬁnal specimen volume. At the same time, a

surface tension, connected with the melted state and decreased rate of the oxy-

gen release, leads to specimen densiﬁcation at increasing temperature. As it

2.5 Y(RE )BCO Bulks 91

is shown in tests, the sample densiﬁcation, controlled by the surface tension,

is the dominating factor [642].

Figure 2.32 demonstrates the YBCO cross-sections obtained at various

heating rates. Even, using the same (in sizes) pellets does not preserve from

considerable diﬀerences in geometry, distribution and structure of porosity in

the melt-processed superconductors. Considerable changes of the sample ge-

ometry and homogeneity have been observed at the heating rate of 30

◦

C/h

because of void formation with sizes above 1 mm (Fig. 2.32a). However, the

void sizes diminished together with the heating rate (Fig. 2.32b and c) that

caused their maximum homogeneity in the case of heating rate of 10

◦

C/h

[640]. At the same time, a void diminishing, caused by the following decrease

of the heating rate, can lead to unfavorable consequences, in particular to

elevated loss of liquid phase. This is not desirable to a high degree in view

of stoichiometry and can lead to loss of control for composition and grain

size during superconductor processing. Therefore, improved physical proper-

ties should be reached, using other methods, in particular by changing the

specimen density and processing parameters of the peritectic solidiﬁcation

[642].

Voids

1 mm

(c)

(b)

(a)

1 mm

Fig. 2.32. Optical micrographs of the polished YBCO cross-sections obtained at

various heating rates: (a) 30, (b)20and(c)10

◦

C/h [642]

92 2 Composition Features and HTSC Preparation Techniques

HTSC applications require high-aligned grains and specimens with perfect

texture that is reached more easily in smaller specimens in the presence of

thermal gradients. The use of large thermal gradient, leading to better align-

ment, could cause microcracking of the sample. The latter spreads out as a

consequence of the anisotropic thermal expansion of 123 phase and secondary

phases (BaCuO

, CuO), which precipitate at grain boundaries. The microc-

racks perpendicular to the ab-plane are the most detrimental for HTSC and

can develop during cooling or re-oxygenation of the sample. Then, puriﬁca-

tion, associated with the crossing of the hot zone, can lead to chemical hetero-

geneities in long samples (which are always obtained using a melt-zone-type

technique) owing to the diﬀerent diﬀusion coeﬃcients of the species. One ob-

serves in such a case a progressive decrease of the barium and copper amount,

correlated with an increase of the amount of 211 phase during the process. Ad-

ditives, such as Y

, which increase the viscosity of the melt, can drastically

limit these phenomena. The more rapid heating rate of 123 phase causes the

ﬁner 211 precipitates. This means that the more rapid is the heating above the

peritectic temperature, the higher is the temperature where the decomposition

of 123 phase begins. The lower the stability of the 123 phase, the more rapid

is the precipitation of the 211 phase, that is, the higher the number of nucle-

ation sites. In rapid heating, deﬁning the ﬁne-sized 211 phase a long plateau

allows the particle coarsening by Ostwald ripening, due to the bigger grains

consuming the smaller ones. In fact, as for the growth of the 123 phase, two

mechanisms can be supposed to limit the growth of the 211 particles dispersed

in a liquid, namely (i) diﬀusion of a solute in the liquid; and (b) reaction at the

interface between 211 phase and liquid. Finally, the maximum temperature,

reached above the peritectic temperature, controls the amount of liquid and its

viscosity. So one expects a large inﬂuence on the growth of the 211 phase. Tak-

ing into account these considerations, diﬀerent methods of melt crystallization

have been proposed during the development of melting techniques, namely:

– Melt-textured growth (MTG) [491]. The textured growth of 123 supercon-

ducting phase from melting, in which the 123 ceramic is used as initial

precursor.

– Liquid-phase processing (LPP) [913]. The liquid-phase technique based on

decreasing of prolonged treatment of samples at maximum temperatures

with the aim to prevent an undesired growth of the 211 particles, that

occurs intensively above the peritectic temperature.

– Zone melting (ZM) [684]. This technique applies zone melting to obtain

long samples using 123 ceramic as initial precursor.

– Quench-melt-growth (QMG) [745]. The growth of 123 superconducting

phase from melting, based on the super-fast cooling, in which the 211

phase forms in results of rapid interaction Y

phase with melt into high-

temperature region, where the 211 normal phase to be thermodynamically

stable.

– Melt-powder-melt-growth (MPMG) [297]. This process, using the same

thermal schedule as the QMG process, introduces a drastic crushing of

2.5 Y(RE )BCO Bulks 93

the quenched mixture after quenching in order to obtain a ﬁne distribu-

tion of the Y

phase, leading further to a ﬁner 211 phase, which is

uniformly distributed.

– Powder-melt process (PMP) [1199]. The same as QMG process, starting

from a Y

BaCuO

+3BaCuO

+ CuO mixture instead of pre-synthesized

123 phase. The result is equivalent to the QMG process, but without any

overheating of the sample.

– Solid-liquid-melt-growth (SLMG) [976]. Using the same thermal proﬁle as

the PMP process, the SLMG process diﬀers by the starting material: a

mixture of Y

+BaCuO

+ 2CuO is used here. These two techniques

diﬀer essentially in the morphology of the 211 phase: needle shape in the

case of SLMG instead of quasi-spheres for PMP process.

– Microwave-melt-texture-growth (MMTG) [146]. This technique was devel-

oped to take into account the high thermal gradient for both increasing

the solidiﬁcation rate, and directing this process.

– Magnetic-melt-texturing (MMT) [190]. This technique uses a static mag-

netic ﬁeld in order to modify the microstructure of a material and to obtain

samples with a good texture. This process exploits magnetic anisotropy

of elementary cell that may be elevated by replacing Y ion with ion of

rare-earth element (RE).

–DopingYBCObytheRE ions in order to form local stresses caused by

the deformation mismatch of the YBCO and REBCO crystalline lattices,

that is an important origin of the magnetic ﬂux pinning and corresponding

increase of J

[627,628].

The thermal cycle, used in a classical MTG process, is shown in Fig. 2.33.

Here, the partial melting of a pre-sintered 123 sample is induced between 1100

and 1200

◦

C, preserving practically the geometry of the precursor sample. Fi-

nally, the sample is very slowly cooled (1–2

◦

C/h) under an oxygen atmosphere

in a thermal gradient [1126].

Introduction of silver [992], resin impregnation [647,1069,1070] and also us-

ing other additives improve the structure-sensitive properties of Y(RE)BCO.

In general case, these additives are used to decrease the size and morphology

modiﬁcation of the 211 phase. They act in three ways, namely: (i) changing

the 123/211 surface energy, (ii) changing kinetics of diﬀusion process, and

(iii) forming nucleation sites for the 211 phase. In particular, an additional

large dispersion of the 211 particles (up to 40 wt%) to the stoichiometric 123

phase is able to diminish the 211 particle size, formed during decomposition of

the 123 precursor sample [745]. To improve superconducting microstructure

and properties, the doping additives used are: Ag (Ag

O, Au/Ag) [998,1171],

Pt (PtO

) [509,777], Sn (SnO

, BaSnO

) [186,685,717], Zr (ZrO

, BaZrO

)

[132,322], Ce (CeO

,BaCeO

) [185,187,664], Ca [391] and SnO

/CeO

[664].

The formation of preferable orientation of grains is necessary for maximum

use of anisotropic properties of HTSC in the speciﬁc products. The specimen

texture could be formed and controlled by using the top-seeded-melt-growth

94 2 Composition Features and HTSC Preparation Techniques

max

Imposed Thermal Gradient

Slow Cooling

Under OxygenUnder Air

Room Temperature

Hot Zone

900°C

Fig. 2.33. Typical thermal cycle used in the MTG process [1126]

(TSMG) technique [729]. In this process, the crystallites-seeds from RE equiv-

alents of the 123 phase are introduced.

They have the higher temperature of

the peritectic decomposition. In this case, the 123 phase nucleates and grows in

the speciﬁc direction (Fig. 2.34). Moreover, the weak links decrease consider-

ably accompanied by proper increase of the grains. The peritectic solidiﬁcation

of YBCO by using seeds leads to growth morphology of faceted grains with

symmetry, depending on nucleus crystallite, for example, SmBCO (Sm-123),

NdBCO (Nd-123) or other RE barium cuprate. However, standard procedure

by using seeding (e.g., SmBCO) proposes high overheating of the sample, that

is necessary to (i) avoid numerous formation of nuclei, (ii) increase liquid in

the sample and (iii) eliminate an eﬀect of precursor microstructure [1126].

In order to get over the numerous nucleation and arbitrary growth of 123

crystals, beginning from substrate material at prolonged process of TSMG,

as observed in test, the increasing of the formation temperatures for diﬀerent

REBCO phases together with their ionic radiuses could be used [728]. Then,

a selection of corresponding RE composition permits to diminish an under-

cooling area ΔT = T

− T

(where T

is the peritectic temperature and T

is the temperature of grain growth in the sample under constant conditions)

at slow cooling to homogeneous temperature (Fig. 2.34c). In this case, is the

grain growth rate decreased and unstable solidiﬁcation eliminated. The kind

of cooling process, that is, slow continuous cooling or jump-like decreasing of

temperature during one or some stages renders considerable eﬀect on quality

of the ﬁnal crystallographic structure [479].

Carrying out the TSMG process in an air atmosphere with seeds from

Nd,Sm,EuandGdformsREBCO samples with decreased critical tempera-

ture T

accompanied by broad superconducting transition. This is caused by

There are hot seeding and cold seeding [480]. In the ﬁrst case, the seed is placed

on YBCO sample at room temperature, in the second case, the seed is placed at

temperature T

max

above the peritectic temperature (Fig. 2.34c).

2.5 Y(RE )BCO Bulks 95

100/010

100/010100/010

100/010

001

Hot Seeding

(c)

(b)

(a)

Cold Seeding

max

ΔT

Quenched

Sample

Melt Processed

Sample

Room Temperature

Seed

Boundary

Fig. 2.34. Schematic diagram of the top-seeded melt growth process applied to

form and control specimen texture (a)viewfromabove,(b)viewatthesideand

the formation of hard precipitation of RE

1−x

2−x

7−x

because of near

cation radius between pointed RE and Ba. Carrying out the TSMG process

in the controlled atmosphere with low partial pressure of oxygen (1 or 0.1%

96 2 Composition Features and HTSC Preparation Techniques

of oxygen in Ar) (oxygen-controlled-melt-growth-process (OCMG)) [749], it is

possible to except formation this precipitation. In order to reduce the dura-

tion of TSMG process, the multi-seeded-melt-growth (MSMG) technique [932]

is used with corresponding thermal treatment. This process consists of some

seeds on the YBCO sample. Now, REBCO samples demonstrate the critical

current density that is above 100 kA/cm

(77 K and 0 T) with corresponding

control of composition and microstructure [908]. Considerable magnetic ﬁelds

can be trapped in large-grain melt-processed HTSC (e.g., 10 T at 45 K [880]),

that is much more than in conventional magnets, and it is very important for

various applications.

Experimental Investigations of HTSC

3.1 Experimental Methods of HTSC Investigations

3.1.1 Special Techniques

The complexity of HTSC structures and properties has resulted in a small

number of directed observations and test dependencies of the type “structure–

property.” Among methods of three-dimensional observation the magnetic

vortex structures in HTSC, the small angle neutron scattering and spin pre-

cession of polarized muons have an important application as the means of

microscopic research of the local magnetic ﬁelds, but the Bitter decoration

method is treated as the most often used method of spatial resolution [84].

In particular, the Bitter decoration method utilizes small ferromagnetic par-

ticles for decorating the magnetic domain structure. When these particles are

sprinkled on a material, displaying at its surface an inhomogeneous distribu-

tion of magnetic ﬂux density, the particles are attracted to the regions with

the largest value of the local magnetic ﬁeld. This method has demonstrated

high eﬀectiveness in visualization of the vortex structures localized at defects

[208, 1116].

The great achievement in the magneto-optical characterization of the vor-

tex structures consists of using the Bi-doped iron garnet thin ﬁlms [470], which

operate from below 4.2 K to above 500 K, that is, they very well ﬁt the tem-

perature window used in HTSC. These ﬁlms come in two distinct varieties,

one with magnetization vector perpendicular to the surface of the ﬁlm (ﬁrst

investigations) and another with magnetization vector in-plane. As a result,

a magnetic resolution of 10 μT has been reached. The spatial resolution can

reach values less than the garnet ﬁlm thickness, if the magnetic ﬁeld gradients

are suﬃciently strong. The resolution of 0.4 μm has been reached for imaging

obtained using ﬁlms with thickness of 2 μm. Some achievements in using the

magnet-optical imaging (MOI) technique for visualization of the vortex struc-

tures in diﬀerent HTSC systems have been presented in the overviews [559,

858, 1119]. In combination with a digital camera and an image-processing