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

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168 3 Experimental Investigations of HTSC

CuPt

and a Y-211 phase [777]. A similar heterogeneous nucleation

mechanism acts in the case of CeO

additives. Y

particles, forming as a

result of the reaction between CeO

and a Y-211 phase:

BaCuO

+CeO

→ BaCeO

+CuO, (3.7)

can be similar nucleation sites for the peritectic Y-211 phase [1115].

Y-211 particles coarsen easily in liquid by the Ostwald ripening process

[64]. However, it is diﬃcult to estimate the degree of contribution of Y-211

coarsening to the Y-211 reﬁnement because of the nucleation and growth of

Y-211 particles, taking place simultaneously. The attempt to separate one pro-

cess from the other was concluded in a melt inﬁltration into Y-211 compacts,

where no Y-211 nucleation was involved [533]. The technique was applied to

Y-211 powder compacts with and without CeO

and PtO

additions [533,

539]. The results show that the addition of either CeO

or PtO

signiﬁcantly

suppresses the Y-211 growth in liquid (Fig. 3.53).

Shape control of Y-211 particles is required to yield ﬁne Y-211 particles

because the dissolution behavior of Y-211 particles in liquid during the peri-

tectic reaction depends on the Y-211 shape [538]. It is a function not only

of the initial powder composition, but also of the type of precursor powder

[718]. In general, the shape of peritectic Y-211 particles of a stoichiometric

Y-123 system is a prism-like one, while the shape of a Y-211 excess system

is an almost equiaxed granular with facets [530]. During a peritectic reaction

the faceted Y-211 particles are dissolved in liquid and their shape is changed

to an irregular form with round surfaces. Sometimes, a longer Y-211 particle

is divided into several parts, demonstrating that Y-211 particles with a high

aspect ratio are better in reducing size than block-like or equiaxed Y-211 par-

ticles. Introducing CeO

and PtO

additions changes the Y-211 shape in liquid

0 2 4 6 8 10 12

Holding Time (h)

Without Additives

1 wt% PtO

1 wt% CeO

1100°C

211 Particle Size (μm)

Fig. 3.53. Variation of Y-211 particle size against holding time at 1100

◦

Cinsamples

with and without PtO

/CeO

addition [530]

3.4 Melt-Processed Y(RE )BCO 169

to more anisotropic, that is, the needle-like one (Fig. 3.54). These transforma-

tions can be a result of change by the additives of the Y-211/liquid interfacial

energy.

The melt-processed YBCO samples with BaSnO

addition could be grown

with higher rate than in the case of addition absence. This implies that SnO

doping improves eﬀectively yttrium diﬀusive factor [186]. However, in this

case, ripening of 211 particles is often observed. CeO

/SnO

additives (at

greater amount of CeO

) permit to increase growth rate of 123 oxide and

to reach submicron size of secondary phase particles, demonstrating square

shape and distributed uniformity in the matrix [664]. Partial substitution of

yttrium by calcium in large-grain YBCO introduces additional hole-carriers

Undoped Y-123

(a)

(b)

Y-211

5wt.% CeO

-added Y-123

10 μm

μm

Fig. 3.54. Y-211 particles in liquid of (a) undoped Y-123 sample and (b)5wt%

CeO

-added Y-123 sample [530]

170 3 Experimental Investigations of HTSC

in CuO

planes, ensuring a regime of excessive doping [382, 383]. These ad-

ditives permit to improve the structure and to reach sub-micron size of 211

particles without using Pt [391]. Then, it has been shown on the test data for

monocrystals [555] that the substitution Y/Ca occurred without strain. How-

ever, the substitution Ba/Ca or Cu/Ca, which competes with the substitution

Y/Ca, becomes predominant at the compression or tension strain above ∼ 6%.

This is found due to Y and Ca ionic radii being approximately the same. At

the same time, Ba and Cu ionic radii on 20% are greater and smaller than

Ca ionic radius, respectively. Because the regions are highly deformed near

intergranular boundaries in polycrystallite samples, Ca substitutes mainly Ba

and Cu, but not Y, decreasing the local strain. In this case, the formation

energy of oxygen vacancies in CuO

planes and Cu–O links decreases at the

deformation. Therefore, there is segregation of oxygen vacancies in the un-

doped, by Ca of the YBa

7−δ

samples, near intergranular boundaries,

and superconducting properties of the YBCO diminish because of the decreas-

ing of hole carriers. The Ca additives decrease strain and make the formation

of oxygen vacancies disadvantageous energetically. The chemical composition

of intergranular boundaries and adjoining regions remains to be relative to

stoichiometrical one, and J

increases compared to the value in the samples

without Ca.

Higher temperature incongruent melting of Nd-123 and Sm-123 compared

to Y-123 [729, 1179] and greater solubility of Nd and Sm in liquid compared

to Y [1169] increase both a ripening Sm-211/Nd-422 and a growth of Sm-123/

Nd-123 regions. Addition of PtO

does not signiﬁcantly inﬂuence decreasing

of particle size compared to the eﬀect on Y-211 particles [730]. A precursor

with small Nd-422 particles deﬁnes dispersion of small Nd-422 particles in the

ﬁnal sample [746]. The improved Sm-211/Nd-422 structure can be reached by

combining CeO

additives with grinding of precursor powder [542]. In this

case, an anisotropy of Sm-211/Nd-422 particles into melt due to change of

energy of the RE-211(422)/liquid interface is increased. Doping of precursor

by Au/Ag additives diminishes Nd-422 particles down to sub-micron level that

is approximately three times smaller than that in the absence of additives and

decreases ripening of Nd-422 particles in peritectic region. In this case the

doped grains consist of spherical Nd-422 inclusions compared to needle ones

that are observed usually in monolithic sample [479].

An increase of critical current together with the amount of Ag has dis-

crepant character [316,779]. Nevertheless, the presence of silver (up to 20 wt%)

in YBCO samples, which demonstrate initially weak links, can improve them

signiﬁcantly [914]. This is manifested in the form of corresponding increas-

ing of sample density, orientation and growth of superconducting grains.

As result, an increasing of superconducting properties together with silver

amount (Fig. 3.55) is possible. The increase of J

for small amounts of Ag

(≤ 20 wt%) may be due to the acceleration of densiﬁcation and grain growth.

At the same time, the decrease of J

for large amounts of Ag (> 20 wt%)

may be due to the excessive increasing of the normal (non-superconducting)

3.4 Melt-Processed Y(RE )BCO 171

phase; 20 wt% of Ag

O can increase critical current density in the ab-plane

)by∼ 200% and critical current density along c-axis (J

)by∼ 150%

at 5 K. These results are correlated with a reduction of the microcrack den-

sity parallel to the ab-planes in YBCO/Ag

O composite. The proper increase

of superconducting properties is depicted in Figs. 3.56 and 3.57. Silver addi-

tives can additionally improve the ﬂux pinning that also leads to increasing

of superconducting properties [694].

3.4.5 Mechanical and Strength Properties

Microcracks form in the ab-plane of Nd-123 orthorhombic phase during cooling

even at small numbers of Nd-422 particles due to the diﬀerence of thermal

expansion between matrix and inclusions. These planar defects are absent

both in the Y-123 tetragonal phase [203, 535] and in the Nd-123 tetrago-

nal phase with small Nd-422 particles (Fig. 3.58) [201], that is, they are not

deﬁned by crystallization process [376]. In the samples with much larger Nd-

422 particles, the ab-microcracks also develop in the tetragonal parts of the

sample (Fig. 3.59). The crack spacing in the orthorhombic Nd-123 is much

smaller than in the tetragonal part. This suggests that for both the orthorhom-

bic and tetragonal 123 phases there is a critical size of 422 (211) particles,

below which ab-microcracking associated with 422 (211) particles does not

occur. The higher critical particle size for the tetragonal 123 phase is asso-

ciated with the absence of additional shortening of the c-axis due to oxygen

uptake or/and inﬂuenced by the higher fracture toughness of the tetrago-

nal 123. Besides the ab-microcracks, the c-microcracks are also observed in

melt-grown YBCO. These c-cracks stop at the boundary between orthorhombic

0.10

0.08

0.06

0.04

0.02

0.00

0 50 100 150 200 250 300

1%Ag

0%Ag

10%Ag

60%Ag

50%Ag

20%Ag

J (A/cm

)

V (mV)

Fig. 3.55. V –J diagrams for diﬀerent Ag-doping of YBCO [914]

172 3 Experimental Investigations of HTSC

(a)

0 10203040

T (K)

(K)

50 60 70 80 90

0 102030405060708090

0 wt.% Ag

20 wt.% Ag

(b)

(A/cm

)

(A/cm

)

0 wt.% Ag

20 wt.% Ag

Fig. 3.56. Temperature dependencies of the self-ﬁeld critical current density: (a)

at H ||c and (b)atH ||ab for 0 wt% and 20 wt% Ag

O [694]

and tetragonal phase (see Fig. 3.60), demonstrating formation of the c-

cracking during tetragonal–orthorhombic phase transition [198].

Figure 3.61 shows interesting change of the ﬁve-domain sample shape.

An originally circular-shaped sample can deform during solidiﬁcation. The c-

growth parts of the sample are bulged and the a-growth parts have shortened

radii when the sample is observed from above.

Due to this, there is speciﬁc

macrocracking of the sample.

Higher 211 concentration at 90

◦

boundaries conﬁrms existence of the a-GSs along

◦

boundaries [197].

3.4 Melt-Processed Y(RE )BCO 173

1.4

1.2

1.0

0.8

0.6

0.4

0.2

0.0

01234

0123 4

H (T)

0 wt.% Ag

20 wt.% Ag

= 77 K

(b)

(a)

(A/cm

)

(A/cm

) × 10

0 wt.% Ag

20 wt.% Ag

Fig. 3.57. Magnetic ﬁeld dependencies of the critical current density (at T = 77 K):

(a)atH ||c and (b)atH ||ab for 0 wt% and 20 wt% Ag

O [694]

A deﬁciency of the 211 particles can induce tensile residual stresses in the

c direction in the central c-GS and also tensile tangential stress in the ab-plane

at the sample surface (in the a-GS). These stresses can cause ab-microcracking

and detrimental radial macrocracking [197]. Other reasons for macrocracking

are the great thermal gradients at sample cooling [910] and magneto-elasticity

eﬀects connected with magnetic ﬂux pinning, causing residual stresses dur-

ing the sample magnetization [879]. As a result, the cracks grow along the

(100)/(010) plane and decrease a trapped magnetic ﬁeld [709]. At increas-

ing oxidization duration of the melt-textured YBCO samples, an aging of

mechanical and superconducting properties [918] may be observed. It is ac-

companied by increasing dislocation density, formation of great chippings at

174 3 Experimental Investigations of HTSC

10 μm

Orthorhombic

Phase

Tetragonal

Phase

Fig. 3.58. Dense ab-microcracks in the orthorhombic Nd-123 phase and no

ab-microcracking in the tetragonal Nd-123 phase (small 422 particles) [201]

rather low-angle grain boundaries and around 211 inclusions and also by in-

tensive degradation of Y-123 matrix near microcracks and the 211 particles.

The studies of textured YBCO ceramics at helium temperatures and high

magnetic ﬁelds show that the latter can cause the dynamics of twinning dis-

locations, leading to cleavage stresses, that can fracture the sample along the

planes parallel to the basic ab-plane [240]. 211 particles deﬂect crack, increas-

ing its path. In this case, the crack branching at the 211 inclusions [614] and

de-lamination of 211/123 interface [197] may be observed. As the 211 parti-

cle has a higher Young’s modulus compared to the 123 matrix, they can also

10 μm

Orthorhombic

Phase

Tetragonal

Phase

Fig. 3.59. Higher ab-microcrack density in orthorhombic than in tetragonal Nd-123

phase (large 422 particles) [201]

3.4 Melt-Processed Y(RE )BCO 175

c-crack

Orthorhombic

Phase

Tetragonal

Phase

10 μm

Fig. 3.60. c-crack stop at the boundary between orthorhombic and tetragonal phase

[198]

act as grains-bridges, pinning down the crack opening and decreasing tensile

stresses in the crack tip [338].

Silver dispersion (Ag

O, Au/Ag) decreases considerably cracking and

porosity both in the ab-plane and along c-axis, increasing J

signiﬁcantly [376,

694, 998]. Figure. 3.62 shows the magnetic ﬁeld distributions in YBCO bulk

(0 wt% Ag) at 9.0 and 8.5 T in the decreasing ﬁeld process from 10 T. Magnetic

1 mm

Fig. 3.61. Polarized-light photomicrograph, revealing the higher density of 211 par-

ticles (gray diagonal regions) in a-GSs along 90

◦

-grain boundaries and also bulging

of c-GSs and macrocracks parallel to ab-plane [197]

176 3 Experimental Investigations of HTSC

(a)

(c)

(b)

(d)

Fig. 3.62. Magnetic ﬁeld distribution in YBCO samples during decreasing process

from 10 T down to: (a)9.0T,0wt%Ag;(b)8.5T,0wt%Ag;(c) 0 T, 10 wt% Ag;

(d) 0 T, 20 wt% Ag [709]

Twin Planes

Crack

Fig. 3.63. Light micrograph of YBCO sample (without Ag) near the crack edge

with higher magniﬁcation observed under polarized light. Twin planes can clearly be

seen on the sample surface. Note that the crack crosses twin planes at 45

◦

, showing

that the cleavage surface is (001) or (010) plane [709]

3.4 Melt-Processed Y(RE )BCO 177

ﬁeld distributions show clearly the sample fracture during the decreasing ﬁeld

process from 9.0 to 8.5 T. At the same time, the additives of 10 wt% Ag and

20 wt% Ag prevent the sample fracture at the same decreasing ﬁeld process

even after total removal of external ﬁeld [709]. The observed fracture of the

sample into two pieces by macrocrack occurs along (001) or (010) planes,

showing that those are the cleavage planes in YBCO. This is followed from

the concept that twin planes are {110}, but the fracture plane crosses at 45

◦

from twin planes (Fig. 3.63). Then, it may be concluded that the cleavage oc-

curs at (001) or (010) planes. The Ag addition (Ag

O) decreases considerably

the ab-microcracking, introduced in TOPT by microstresses, caused by 211

(a)

(b)

Fig. 3.64. Microcracks in YBCO/Ag

O samples after etching: (a)0wt%Ag

Oand

(b)12wt%Ag

O [198]