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

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118 3 Experimental Investigations of HTSC

( j)

(i)

Longitudinal Cracks

0.05 mm

Fig. 3.10. (i) longitudinal cracks in BPSCCO superconducting core [790], (j)scheme

of macrodefect (black area in gray intermediate layer) in Josephson junction [1014]

epitaxial ﬁlms onto bicrystalline SrTiO

substrates [205] has conﬁrmed the

behavior of intercrystalline boundaries of Josephson type. Earlier, this result

has been predicted by using tests on bulks [841]. On the other hand, the tests

on oriented YBCO bicrystals [38] showed that some grain boundaries, includ-

ing high-angle ones, demonstrate under high magnetic ﬁelds much more high

transport properties compared with behavior that is proper for weak links

subjected to Josephson eﬀects. Analogous phenomena take place in bulks,

3.2 Intergranular Boundaries in HTSC 119

in which there is often small intergranular current component that does not

depend on ﬁeld [248].

The general results are usually based on the bicrystal investigation of three

main types, namely (i) epitaxial thin ﬁlms that are grown on bicrystalline

[205, 353, 474] or otherwise specially prepared [324, 619, 1122] substrates

with predetermined misorientation relationships and boundary planes (on a

macroscopic scale); (ii) very large grain (millimeters) melt-processed bicrys-

tals (MPB) that contain considerable volume fractions of secondary phases

and have random misorientation relationships and macroscopically meander-

ing boundary planes [279, 795]; and (iii) ﬂux-grown bicrystals (FGB) with di-

mensions of the order of a few hundred microns, misorientation relationships

of the θ [001] type and comparatively straight, but uncontrolled boundary

planes with large tilt sites [38, 39, 599].

The study of electric resistance of the [001] symmetrical tilt boundaries

at the normal state of sample revealed common dependence of electromag-

netic properties on the misorientation angle θ.ThevalueofJ

(gb)/J

(g)

(where J

(gb)andJ

(g) are the critical current densities of the boundaries

and neighboring grain interiors, respectively) decreased rapidly with a 1/θ-

like dependence for θ values up to about 20

◦

.Forθ>20

◦

, an approximately

θ-independent, uniformly low value of J

(gb)/J

(g) ≈ 0.01 was observed [205,

206]. The same behavior of grain boundaries with low/high angles of mis-

orientation is observed in all HTSC. In particular, epitaxial bicrystal ex-

periments using thin ﬁlms of Bi-2212 [20, 681, 1129], Tl-1223 [755], Tl-2212

[122] and Tl-2223 [921, 1129], all showed a qualitatively similar dependence

of J

(gb)onθ.The1/θ dependence/plateau behavior of J

(gb)/J

(g)vsθ

led researchers [122, 205, 754] to postulate that the dislocations, covering the

structure of intercrystalline boundaries, determine their electromagnetic prop-

erties. As the misorientation angle of the boundary increases, barriers with

weak link properties progressively pinch oﬀ the supercurrent. At some critical

transition angle, the barrier regions completely control the boundary proper-

ties and a weak link forms. The dependence on θ suggests that the primary

grain boundary dislocations (PGDBs), constituting the boundary structure,

cause the weak link behavior [205, 353]. In this model, the boundary is fre-

quently described as a uniform array of like edge dislocations with spacing,

which are characteristics of symmetrical tilt boundaries as speciﬁed by Frank’s

formula [877].

There are many attributes of dislocation network, including the weak link

behavior, namely (i) overlapping of dislocation cores with cation stoichiome-

try [754]; (ii) overlapping of elastic strain ﬁelds, caused by PGDBs [887]; (iii)

agreement of lattice parameters through stoichiometry on oxygen for elimina-

tion of lattice mismatch at intercrystalline boundary [1202, 1203]; (iv) local

alterations of T

and chemical potentials due to elastic strains [373]. The test

120 3 Experimental Investigations of HTSC

data show the transition from strong to weak coupling at θ ≈ 10

◦

which is

most evident in thin-ﬁlm YBCO bicrystals [205, 353]. This transition could

be explained by narrowing of superconducting connecting dislocation cores

[205] or overlapping of their strain ﬁelds [149]. This model assumes absence

of superconductivity of dislocation cores and limitation of the supercurrent

paths by channels, connecting these cores, and also by neighboring boundary

regions. The regions of perfect crystalline structure into space between dis-

locations or symmetric facets in the absence of the void-like defects can be

potential sources of current-carrying paths deﬁned by strong links [94]. The

orthorhombic crystalline structures of cuprate superconductor could consider-

ably complicate investigation of their intercrystalline boundaries. In addition

to diﬀerent lattice parameters a and b, YBCO family possesses a good devel-

oped twinning structure [344], but BSCCO family demonstrates non-regulated

modulation along b-axis [1117]. Both the lattice parameters and YBCO su-

perconducting properties change with oxygen concentration. In particular, the

lattice parameters a and c increase with decreasing oxygen in the structure,

causing an expansion of unit cell. Therefore, the deﬁciency of oxygen and

corresponding weakening of superconducting properties may be intrinsic to

intercrystalline boundaries in YBCO.

The electromagnetic properties of weakly coupled boundaries vary sub-

stantially, and perhaps systematically, with position along the grain boundary.

The patches of “better” material are separated by weak or non-superconducting

regions, as shown schematically in Fig. 3.11. One of the main causes of these

structure and composition alterations could be the local oxygen depletion and

oxygen disorder in YBCO structure [94, 218, 1201]. Other possible microstruc-

ture sources of heterogeneity include: (i) cation composition modulation, ex-

isting within the boundary plane [372, 1107]; (ii) “wavy” boundaries, facets

with arbitrary conﬁgurations and facet junctions, observed in epitaxial thin-

ﬁlm bicrystals of YBCO and Bi-2212 [09, 420, 1077]; (iii) strain ﬁelds, caused

by intrinsic and extrinsic intercrystalline dislocations and also by regular dis-

tribution of facets [1079, 1080]; (iv) the intersections of twin planes with the

intercrystalline boundary plane in the case of YBCO [43]; and (v) oscillatory

changes in misorientation of the neighboring grains due to twinning in YBCO

[43, 1079, 1080]. Diﬀerent faceted structures are shown in Fig. 3.12.

The superconductor crystals that nucleate on one substrate crystal can

often grow past the substrate boundary and over the “other” crystal for ap-

preciable distances before impinging on a crystal growing in the “correct”

orientation. This overgrowth, which may be more prevalent for certain mis-

orientation angles, results in “wavy” boundary topography, causing boundary

morphology and properties [505, 606, 1080]. On the other hand, the twinning

in each YBCO single crystal alone results in discrete, systematic changes in the

local misorientation relationship along the boundary, which should produce

Strong link behavior can be observed in melt-processed YBCO bicrystals [280,

795] and in ﬂux-grown YBCO bicrystals [43] at angle θ up to 20

◦

3.2 Intergranular Boundaries in HTSC 121

∼ξ

Grain 1

Grain 2

n n n

Fig. 3.11. Schematic illustration of the heterogeneity, proposed for weak-linked

grain boundaries in YBCO. The current ﬂows through microbridges (labeled by m)

that exist between non-superconducting regions (labeled n). The electrical char-

acter of the microbridges is assumed to be normal, superconducting or insulating

depending on the model used

concomitant oscillations of PGDBs, as observed for a high-angle boundary in

sintered material [305].

There are numerous examples of PGDBs at the low- and high-angle YBCO

intercrystalline boundaries. Some high-angle boundaries include regularly po-

sitioned amorphous regions [697, 887]. Nevertheless, usual observations show

high degree of the intercrystalline structure localization, ranging only on

1–2 elementary cells to neighboring grains [149, 305, 606]. These observations

agreed with the width of dislocation cores, which are observed on representa-

tive intercrystalline boundaries in ceramics [697].

For the framework of description of the high-angle boundary structure in

terms of PGDB [41], concentrated picture of the “good/bad” regions may be

used, according to defect distribution. Intercrystalline dislocations are local-

ized suﬃciently in order to create periodic ﬁeld of the elastic physically signed

strains. In this case, the accommodation dislocations of the lattice mismatch

at [100] boundary are partial dislocations and can dispose suﬃciently far from

the boundary into one from grains [125, 305, 754]. These partial dislocations

are joined with boundary by stacking faults, which rather introduce copper

excess in a boundary region. Nevertheless, this boundary can act as strong

link despite its extended structure and possible stoichiometry [255, 586, 619].

Nanofacets are observed at diﬀerent intercrystalline boundaries [128, 697].

In this case, the facet planes are often dictated by a crystalline structure.

Facets are observed on boundaries, possessing relatively simple crystallogra-

phy and low-coincidence index (Σ)values,whereΣ is the fraction of lat-

tice sites of one crystal that are coincident with the other [754], and also on

boundaries with complex geometry [606]. Facets aspire to align parallel to

planes with lowest index of one from crystals and are observed, as rule, in

(100), (010) and (001) planes. The facet formation introduces considerable

122 3 Experimental Investigations of HTSC

(a)

(b)

Grain 1

Grain 2

Grain 1

Grain 2

50 nm

010

Fig. 3.12. Faceted structures in HTSC: (a)ag|| [010] diﬀraction-contrast TEM im-

age of nominally pure [001] tilt boundary in 10

◦

[001] bicrystal of Bi-2212 [1082]; (b)

perspective and [001] plan view schematics of the “stepped” boundary topography,

composed of pure symmetric tilt facets and pure twist facets, that was observed in

many sections of the [001] tilt 2212 bicrystals [1082]

3.2 Intergranular Boundaries in HTSC 123

(c)

130

221

GB2

20 nm

110

010

221

130

50 nm

(d)

Fig. 3.12. (c) diﬀraction-contrast TEM image showing saw-tooth faceting onto

(130) and (22

1) plane in GB2 [1079];(d) enlargement of a portion of (c), showing

that the (130) facets further facet onto (010) and (110) subfacets; the dot-like strain

contrast along the subfacets is produced by the grain boundary dislocations [1079]

structure heterogeneity in the boundary. Due to limited size of facets, macro-

scopic boundaries with diﬀerent planes may be considered as mixture of

some types of the boundary facets of ﬁxed structure. The intersection lines

of facets possess their own atomic structure, and secondary dislocations are

often associated with them [43]. In this structure level, heterogeneities, which

are important for superconducting properties, include saw-tooth structure of

facets, additional faceting of facets, initially appearing plane, intensive strains

at facet junctions, inhomogeneously distributed dislocations at boundaries in

the range of individual facets and favorable conditions for dislocation-division

with Burgers vector b = 100 in partial dislocations at the center of some

facets (see Fig. 3.12c and d) [1079].

The facet junctions can inﬂuence stoichiometry of intercrystalline bound-

aries; in particular, it is possible for the concentrated excess of copper to

be connected with facet distribution in volume bicrystals with a low mis-

orientation. Some more oscillations of stoichiometry on copper with period of

124 3 Experimental Investigations of HTSC

about 100–200 nm are observed at high-angle boundaries [42]. Contrast strains

take place in the regions with copper excess, but the facet morphology and

the strain origin have not exactly been stated. These copper redistributions

together with facet topography and/or strains in facet junctions can cause

intercrystalline electromagnetic properties.

Dissociation of PGDBs, forming pairs of partial dislocations, can con-

siderably inﬂuence the value of the critical angle at transition from weak

links to strong ones due to the accompanying decreasing of dislocation spac-

ing. During PGDBs dissociation, superconducting channels narrow and can

lead analogously to microbridges, demonstrating Josephson eﬀect, that is ,

weak link behavior [630]. This transition could be expected by decreasing

below the coherence length (about 1 nm in ab-plane). These narrow supercon-

ducting channels, caused by the dislocation dissociation in partial ones, form

weak links at intercrystalline boundary (one could be strongly connected,

if consists of PGDBs only). Thus, a transition to lower energy structure of

boundaries (at dislocation dissociation in partial ones) will suppress an inter-

crystalline supercurrent owing to mechanism of the superconducting channel

narrowing.

Strain ﬁelds near tilt boundaries must considerably diﬀer from ﬁelds near

twist boundaries due to diﬀerences between parallel rows of edge dislocations

and cross-shaped lattices of screw dislocations, covering their structures, re-

spectively. While the boundaries along [001] direction have, as a rule, one row

of dislocation, in the case of mixed orientation of grains, there are at least two

intersecting rows of dislocations. This alters a form of channel with strong

coupling owing to inclusion of both rectangular sites, located between paral-

lel rows of dislocations, and almost point contacts of high-angle boundaries

with mixed orientation (Fig. 3.13). These boundaries retain a good connected

Fig. 3.13. Schematic presentation of possible structures for intercrystalline bound-

aries in general case. The horizontals, verticals and diagonals show dislocation net-

work, grey regions present barriers for supercurrent, white sites are the channels of

strong links. Contrast to the model of boundaries directed along [001] direction (see

Fig. 3.11), the strongly connected channels at these boundaries to rather point, than

linear contact

3.2 Intergranular Boundaries in HTSC 125

system of ﬂaw pattern of the supercurrent (see test results [280]), despite

far smaller number of “strong” channels predicted by the PGDBs model

[205, 206].

Faceted structures in YBCO and Bi-2212 diﬀer in the type of facets, as well

as in regularity of their distribution. At nanoscale in Bi-2212, there is non-

regular distribution of symmetric tilt and twist [001] facets. At the same time,

very regular saw-tooth distributions of tilt facets dominate in YBCO [1079,

1107]. In this case, the strains observed in YBCO bicrystals are always in ac-

cordance with saw-tooth facet intercrystalline structure [1079, 1080]. Elevated

strains could be caused by incomplete elimination of long-range component

because of small number of dislocations that are contained at the facets with

limited size.

Despite diﬀerences in the dislocation topography and distribution at the

boundaries in Bi-2212 and YBCO, there is one common feature concerning the

Burgers vectors of these dislocations, which may be correlative with supercon-

ductivity of boundaries. It is known that existence of the partial dislocations

is common for intercrystalline boundaries in Bi-2212 and YBCO [1079, 1080,

1107]. Moreover, in both materials, Burgers vector value and direction corre-

spond to distance between two nearest atoms of oxygen in the CuO

planes

of unit cells. Then, the strain caused by these dislocations can initiate local

superconductivity. In particular, the shear strain in {110} planes is associated

with the beginning of superconductivity in YBCO on the basis of analysis of

the anomaly in thermal expansion factors in a-axis and b-axis directions [110,

111, 439].

It is known that a layer of material with diﬀerent superconducting proper-

ties exists at YBCO grain boundaries. This layer appears to be hole-deﬁcient

YBCO [39, 94, 1202]. However, key questions about the nature of this layer

remain unanswered: (i) What is the average width of the hole-depleted zone,

and does it depend on the form of the material (bulk or thin ﬁlm) or process-

ing method? (ii) Are the magnitude of depletion and width of the depleted

zone uniform along the boundary? (iii) What is the microstructural origin of

the holes’ depletion?

There is correlation between the value of hole depletion and initiation of

eﬀects, demonstrating the weak link behavior [39]. The hole-concentration

proﬁle studies are based on the concept that the hole depletion zone width

exceeds both the superconducting coherence length for YBCO and the ap-

parent structural width of the boundary. The widths reported vary from one

material form to another: ∼ 8 nm wide aﬀected zone at polycrystalline thin-

ﬁlm boundary [94]; the width of the depletion zone in FGBS is ∼ 60 nm [39]

and, in sintered materials, a width the order of 10–20nm [1202]. All of these

values exceed the width of the structurally disordered region (∼ 1nm)andthe

metal cation non-stoichiometry (∼ 5 nm) [42], and also the coherence length ξ.

Unfortunately, it is not possible to associate the reduced density of holes with

a speciﬁc atomic defect such as oxygen non-stoichiometry or oxygen disorder,

these being the most likely candidates. In this case, there may be important

126 3 Experimental Investigations of HTSC

diﬀerences in the structure, composition and electronic properties of grain

boundaries in thin ﬁlms as compared to their bulk counterparts. Therefore, it

is expected that hole density variations might occur on a number of diﬀerent

length scales in both strongly and weakly coupled boundaries, and also in

superconductors with diﬀerent forms and sizes.

Thus, the main features of grain boundaries in the considered HTSC sys-

tems (Fig. 3.14 illustrates some of them) include:

– the cores of the grain boundary dislocations, where the nearest neighbor

conﬁgurations are disrupted and the strains are very large;

– the ribbons of stacking fault, disposing between dislocations that dissociate

to form pairs of partial dislocations;

– the variety of changes in boundary plane that occur as facets form, for

example, due to twinning;

– the diﬀerent dislocation distributions in the various facets and the uneven

arrangement of dislocations that occur in smaller facets;

– the compositional and structural changes, including heterogeneous but pe-

riodic copper excess (XS Cu), that are caused by local diminution and

disorder of oxygen, and also cation modulations;

– the electron–hole depletion that occurs at boundaries and also appears to

be non-uniform;

– stress–strain states of diﬀerent length scales that arise as a result of the

dislocation network, dislocation dissociation, facet junctions, phase tran-

sitions, twinning, initiation and growth of pores and microcracks.

These features comprise a boundary that is heterogeneous on a variety

of length scales with respect to atomic distributions, electronic structure,

composition and stress–strain state. The above heterogeneities inﬂuence all

structure-sensitive properties of bicrystals and polycrystalline materials.

•

•••••

•

••••

•

•••• ••

Different

Strain Scales

Dislocation Cores

Stacking Faults

Missing Oxygen,

e–holes

Missing Oxygen,

e–holes

XS Cu

Fig. 3.14. Schematic illustration of various types and scales of structural and

compositional heterogeneities at [001] tilt boundaries in YBCO and in other HTSC

3.3 Superconducting Composites, Based on BSCCO 127

3.3 Superconducting Composites, Based on BSCCO

3.3.1 BSCCO/Ag Tapes

Investigations of BSCCO/Ag tapes show that the c-axis alignment reaches

maximum near the “superconductor–silver sheath” interface, but is often ob-

served in the tape center [273, 600]. This microstructure assumes that an

oxide core of Bi-2223 tapes consists of the plate-late grain colonies, in which

the grains are divided by twist boundaries. The boundaries of colonies could

be divided into four types, namely [14, 411, 638] (i) the boundary in the c-

axis direction approximately parallel to the c-axis of colony; (ii) the colony

boundary adjacent to basis plane at one side of the boundary; (iii) the colony

boundary, neighboring to basis planes at both sides of the boundary; and

(iv) three-point junctions of colonies. Generally, in the above cases, the twist

boundaries are low-angle ones existing at the screw dislocations. Despite high

texture, high-angle boundaries are often observed in the tapes. Majority of

colony boundaries in the c-axis direction are the mixed type, including tilt

and twist components and also dislocation network. The value of J

could be

controlled by residual layers of Bi-2212 phase at the twist [001] boundaries

[1096, 1097] and also by Bi-2201 inclusions [1132], if their content is suﬃciently

great.

Amorphous buﬀer layers can exist at all four types of boundaries and

considerably inﬂuence conductive properties, because their thickness usually

exceeds considerably the superconducting coherence length [14, 1186]. The

triple point junctions of the colonies in the form of triangles (Fig. 3.15a) are

usually formed in tapes owing to plate-late growth of the colonies and misori-

entation between them. In these triple point junctions, the amorphous phase

and admixtures are often observed in the tapes even with high J

. Obviously,

the triple point junctions of colonies, exceeding on some orders of magnitude

the coherence length, render negative eﬀect on the transport current in the

tape. Secondary phases, consisting of amorphous structure and admixtures,

form into Bi-2223 colonies separated particles of (Ca, Sr)

CuO

and other ox-

ides (see Fig. 3.15b) [828]. Large particles of secondary phase lead to arbitrary

orientation of Bi-2223 colonies and serious distortion of crystalline lattice, cre-

ating numerous defects in these regions. Mutual intergrowths of Bi-2212 and

Bi-2223 phases that can be usually observed in colonies in the oxide core of

the tape eliminate near silver sheath [1097]. At speciﬁc thermal treatment,

the intergrowth of superconducting phase is possible in silver [887]. The prop-

agation of these intergrowths increases Ag/BSCCO interface and assumes an

energetic advantage of their growth along concrete crystallographic planes of

Ag. Then, the silver texture is able to eﬀectively inﬂuence superconducting

grain alignment.

There are other microstructure features of Bi-2223/Ag tapes, namely [001]

edge dislocations, high density of stacking faults, bending and interrupting of

(001) planes in local small regions, which can destroy perfection and continuity