Martienssen W., Warlimont H. (Eds.). Handbook of Condensed Matter and Materials Data

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732 Part 4 Functional Materials

Table 4.2-23 Typical properties of sintered YBCO according to [2.49]

Flexural strength (MPa) 216 ±16

Young’s modulus (GPa) 141.8 ( = 87%); 180–200 ( = 100%)

Fracture toughness, K1c (MPa m

1/2

) 1.07±0.18

Critical ﬂow size (µm) 15

Thermal expansion, 50–450

◦

C(×10

−6

−1

) 11.5

Speciﬁc heat at 82% theoretical density (J kg

−1

)) 431

Thermal conductivity (Wm

−1

)) 300 K 2.67 ( = 85%)

77 K

1( = 85%)

Resistivity (µΩ cm) 300 K 670

100 K

220

(K) 92 (width 1 K)

(77 K, H =0)(Acm

−2

) 500–1000 ( = 85%)

 : bulk density

Table 4.2-24 Superconducting properties (transition temperature T

, coherence length ξ, penetration depth λ,GL

coefﬁcient κ, lower critical ﬁeld H

, upper critical ﬁeld H

) for YBCO [2.54]

(K) 92 [2.25]

Coherence length (nm) ξ



0.6 [2.50]

0.4 [2.51]

⊥

2.7 [2.50]

3.1 [2.51]

Penetration depth (nm) λ



26 [2.52]

90 [2.51]

130 [2.53]

⊥

125 [2.52]

800 [2.51]

450 [2.53]

GL coefﬁcient κ



7.6 [2.54]

⊥

37 [2.54]

(T) H



53 [2.55]

⊥

520 [2.55]

(T) H



140 [2.52]

⊥

650 [2.52]

Table 4.2-25 Basic material and critical current density relevant parameters for YBCO according to Larbalestier

et al. [2.22]

Material YBCO

Anisotropy 7

92 K

> 100 T (4 K)

∗

5–7T(77K)

In-plane coherence length ξ(0) 1.5nm

In-plane penetration depth λ(0) 150 nm

Depairing current density 3×10

Acm

−2

(4.2K)

Critical current density ≈ 10

Acm

−2

ρ(T

) ≈40–60µΩ cm

Part 4 2.2

Superconductors 2.2 Non-Metallic Superconductors 733

Oxygen content (7–δ) in YBa

7–δ

(K)

100

7.0 6.8 6.6 6.4 6.2 6.0

Fig. 4.2-38 Critical temperature T

versus oxygen content

7−δ [2.30]

Superconducting Transition Temperature.

For the T

versus oxygen-content behavior, a two plateau depen-

dence is observed (ﬁrst plateau: 6.8 < 7 −δ<7, T

not

(K)

M =Ca M =La

M =Sr

M =Zn

M =Mn

x (YBa

2–x

7–δ

) x (YBa

3–x

7–δ

)x (Y

1–x

7–δ

)

0.0 0.1 0.2 0.0 0.4 0.8 1.2 1.6 0.0 0.1 0.2 0.3

(K) T

(K)

Fig. 4.2-39 T

as a function of the substitution of the cations Y, Ba and Cu by selected cations M according to [2.56–60]

actually constant, slight decrease in T

with decreasing

(7−δ); second plateau: 6.55 < 7−δ<6.65, T

≈56 K)

(Fig. 4.2-38). Complete substitution of yttrium by rare-

earth metals do not show a signiﬁcant change for

the value of T

, with the exception of the lanthanum

= 50 K) and praseodym compounds (not supercon-

ducting) (Table 4.2-21). Other dopants on the different

atomic sizes as a general rule show decreasing T

values

with increasing substitution rates. Selected examples are

shown in Fig. 4.2-39.

Critical Current Density. The critical current den-

sity J

(B) of YBCO thin ﬁlms at three different

temperatures as a function of the magnetic ﬁeld B

⊥c

and

c

is given in Fig. 4.2-40. At temperatures T ≤ 60 K

the maximum of J

(B) occurs when the magnetic ﬁeld

is aligned along the ﬁlm plane. At T = 77 K Roas

et al. [2.61] observed higher J

values at B < 2T

and a crossover of the curves perpendicular and par-

allel to the ﬁlm plane. The crossover seems to depend

on microstructural features, e.g., defect concentration,

twin-boundary distance etc.

Hole Concentration. The oxygen content of YBCO

(YBa

7−δ

) determines the hole concentration in

the CuO

-planes. As shown by Breit et al. [2.62] the

maximum value of T

is achieved at an oxygen content

of 7 −δ =6.94 (see Fig. 4.2-41). This maximum is due

Part 4 2.2

734 Part 4 Functional Materials

42 K

60K

77 K

⏐⏐c

(θ =0°)

⊥c

(θ =90°)

8901234567

Critical current density(A/cm

–2

)

Magnetic field (T)

Fig. 4.2-40 Critical current density with B

⊥c

and B

c

various temperatures in magnetic ﬁelds up to 8 T according

to [2.61]

(K)

Oxygen content 7–δ

6.85 6.90 6.95 7.00

Fig. 4.2-41 T

as a function of the oxygen content 7 −δ

in YBa

7−δ

according to [2.62] (data from neu-

tron scattering experiments: ◦ [2.62],  [2.63]; data from

normal-state resistivity: +[2.64,65]). The line is a guide to

the eyes

to an optimum hole doping of the CuO

planes. At an

overdoped state 7 −δ>6.94, T

decreases as the holes

in the planes exceed the optimum concentration.

Lower Critical Field. Figure 4.2-42 presents the tempera-

ture dependence of the lower critical ﬁeld H

parallel

0 0.2 0.4 0.6 0.8 1

1000

800

600

400

200

t = T/T

(Oe)

||c

⊥c

Fig. 4.2-42a,b Temperature dependence of H

(a) and of

the anisotropy ratio H

⊥

(b) according to [2.66]

and perpendicular to c. The anisotropy of the H

temperature-independent. The temperature dependence

of H

is in good agreement with BCS calculations

(solid lines in Fig. 4.2-42a) as shown by Wu and Shrid-

har [2.66]. The authors obtained H

(0) = 850 ±40 Oe

and H

⊥

(0) = 250 ±20 Oe.

Upper Critical Field. The upper critical ﬁelds H

high-T

superconductors give information about mi-

croscopic parameters, e.g., the coherence length and

their anisotropies in the superconducting state. The

temperature dependence of the upper critical ﬁeld

(T ) for H

⊥c

and H

c

measured on a polydomain

sample [2.67] is given in Fig. 4.2-44. A linear depen-

dence is observed for both directions with slopes of

the critical ﬁeld of −1.9and−10.5TK

−1

for H

c

and H

⊥c

, respectively. The same dependence meas-

ured by the resistivity method is shown by the dashed

lines. Corresponding to these measurements one can

calculate: H

(T = 0K) = 122 T, H

⊥

(T = 0K) =

674 T, ξ

= 3.0Å,andξ

⊥

= 16.4Å.

In low magnetic ﬁelds a region of nonlinear depen-

dence of the upper critical ﬁeld is observed, which could

not be reproduced in monodomain crystals.

Pinning Mechanisms. In high J

value melt-processed

YBCO samples, many defects – such as twin planes,

Part 4 2.2

Superconductors 2.2 Non-Metallic Superconductors 735

T (°C)

x Bi

x Pb

850°C

800°C

750°C

Bi:Pb =1.99:0.19

Bi:Pb =1.91:0.36

Bi:Pb =1.84:0.36

Bi:Pb =1.90:0.16

Bi:Pb = 2.08:0.19

Bi:Pb = 2.13:0.18

Bi:Pb = 2.2:0.05

Bi:Pb = 2.03:0.22

Fig. 4.2-43 Schematical three-dimensional representation

of the single-phase region of the Bi2223 phase, the variables

being the temperature and the Bi and Pb contents [2.68]

stacking faults, grain boundaries, cracks, oxygen de-

fects, dislocations or non-superconducting particles –

can act as pinning centers.

Twin planes are usually introducted into YBCO

crystals at the transformation from tetragonal to or-

thorhombic symmetry. The twin planes only act as

effective pinning centers, when the ﬂuxoids are aligned

parallel to the twin plane while the Lorentz force

acts perpendicular to the plane. Twin planes are not

dominant pinning centers in melt-processed YBCO

samples.

Stacking faults. If extra CuO layers are inserted ad-

jacent to the chain plane of YBCO, this tends to a local

structure of the defects resembling that of the so-called

Y124 phase. These may act as important pinning centers

for H

⊥c

, but are not important for H

c

. From a prac-

tical point of view these are not suitable pinning centers

as it is no possible to prepare samples with or without

stacking faults.

Oxygen defects. Däumling et al. observed a peak in

the magnetization hysteresis in oxygen-deﬁcient single

crystals at an intermediate ﬁeld, which they ascribed to

ﬂux pinning. This kind of pinning center is not the most

important one.

1.0

0.8

0.6

0.4

0.2

0.0

T(K)

(T)

T(K)

91.0

92.0

85 86 87 88 89 90 91 92 93

B⏐⏐c

– 1.9T/K

B⏐⏐ab

– 10.5T/K

B⏐⏐c

B⏐⏐ab

Fig. 4.2-44 Temperature dependence of the upper critical ﬁeld

(solid line: magnetization measurements, dashed line: resistive mea-

surements) [2.67]

Cracks. As mentioned above, the tetragonal to or-

thorhombic phase transition associated with oxygen

addition causes stresses in YBCO samples leading to

cracks in the c-direction. These cracks can contribute in

ﬂux pinning as in the case of twin planes or stacking

faults.

Dislocations. In as-grown thin ﬁlms, many screw

dislocations may occur, which can act as pinning cen-

ters.

Nonsuperconducting inclusions in YBCO may be

effective in ﬂux pinning. Especially the inﬂuence of the

211-phase (or 422-phase in the case of Nd-123) is dis-

cussed. Such inclusions are stable and can be dispersed

in a controlled pattern.

Application of YBCO Superconductors. Applications

of YBCO high-T

superconductors are superconduct-

ing quantum interference devices (SQUIDS), magnetic

shielding devices, infrared sensors, analog signal pro-

cessing devices, microwave devices, and medical

imaging systems. Since melt-textured bulk YBCO ma-

terial shows magnetic pinning properties, a possible

application may be high-ﬁeld permanent magnets. For

ﬂexible tape conductors the YBCO material is coated

onto metallic carrier tapes. Conductors coated with

YBCO are studied intensively, because of the poten-

tially high J

values in high magnetic ﬁelds at T = 77 K.

Typical architectures of these coated conductors are

sequences of polycrystalline metallic substrate or

biaxially-textured metallic substrate – biaxially-textured

Part 4 2.2

736 Part 4 Functional Materials

Table 4.2-26 Present status of Y based coated conductors according to Yamada and Shiohara [2.69] (IBAD: Ion-beam-

assisted deposition, RABiTS: Rolling assisted bi-axially textured substrates, PLD: Pulsed laser deposition, TFA-MOD:

Triﬂuoroacetate–metal organic deposition)

Material YBCO

Typical Process IBAD, RABiTS, PLD, TFA-MOD

Structure Hastelloy (SUS) / oxide buffer / YBCO layer, monoﬁlament

1–2MAcm

−2

−100 A (77 K, 0 T)

Non-SC/SC ratio 100

Length achieved 46 m (reported by Fujikura, Japan AS, Jan. 2003)

Wire conﬁguration Flat tape with thin YBCO layer

Reinforcement Hastelloy, SUS substates

Application Tested –

Remarks R&D stage for 77 K, high ﬁeld

buffer layer – biaxially-textured superconducting YBCO

layer – Ag stabilizer. The texturing is achieved by rolling

and recrystallization.

Wafers of YBCO thin ﬁlms are prepared for mi-

crowave applications. Thermal coevaporation is used as

the commercial deposition method. With this method

substrates of about 20 × 20 cm

can be coated at rates

of 20–30 nm/min. A double-sided deposition can be

achieved in two deposition steps. The sol–gel method

as well as the chemical solution deposition method are

promising with respect to lower production costs. De-

tailed information on applications is given in [2.81]

(Table 4.2-26)

Table 4.2-27 Structural data of superconducting compounds in the Bi

−

O system and of Pb substituted

Bi2223 [2.15]

Compound Space group Lattice parameters

a (nm) b (nm) c (nm)

CuO

Cmmm 0.5361 0.5370 2.4369

CaCu

Fmmm 0.5408 0.5413 3.081

Fmmm 0.539 0.539 3.71

(Bi

1.6

0.4

)Sr

Fmmm 0.5413 0.5413 3.7100

Table 4.2-28 Maximum T

and a compilation of several values of the coherence lengths, London penetration depths and

upper critical ﬁelds at T = 0 K for superconducting compounds of the Bi

−

O system

Parameter Bi2201 Bi2212 Bi2223 Pb-doped Bi2223

(K) 34 [2.70]

90 [2.71]

110.4 [2.72]

109 [2.73]

(0) (nm) 3.5–4.5 [2.74]

4.0 [2.75]

≥1.09 [2.76]

–3.8 [2.77]

2.9 [2.78]

2.9 [2.73]

≈ 1 [2.79, 80]

4.2.2.3 Superconductors Based

on the Bi–Sr–Ca–Cu–O System

General Properties

Three superconducting phases exist in the Bi

−

O (BSCCO) system: Bi

(abbrevi-

ated using the stoichiometric coefﬁcients as Bi2223),

CaCu

(Bi2212), and Bi

CuO

(Bi2201).

A review about phase diagram studies in the sys-

tem Bi

−

O, particularly in view of

the superconducting compounds, is given in [2.68].

The Bi2212 phase is thermodynamically stable over

a relatively wide temperature range. By contrast, the

Part 4 2.2

Superconductors 2.2 Non-Metallic Superconductors 737

Table 4.2-28 Maximum T

and a compilation of several values of the coherence lengths, London penetration depths and

upper critical ﬁelds at T = 0 K for superconducting compounds of the Bi

−

O system, cont.

Parameter Bi2201 Bi2212 Bi2223 Pb-doped Bi2223

(0) (nm) 1.5 [2.74]

0.16 [2.77]

0.093 [2.78]

0.02 [2.80]

(0) (nm) 310 [2.82] 438 [2.75]

180 [2.83]

–310 [2.84]

165–230 [2.72]

≈ 250 [2.73]

, [2.15]

194 [2.79]

(0) (nm) 1×10

–4× 10

[2.85]

c2(ab)

(0) (T) 43 [2.74]

533 [2.77]

1210 [2.78]

c2(c)

(0) (T) 20.2 [2.75]

16–27 [2.74]

60 [2.86]

;70

39 [2.78]

297 [2.79]

Types of the investigated samples

Single crystal

Overdoped single crystal

Whisker

Thin ﬁlm

Textured Ag-sheathed tape

Bi2223 phase is stable only over a narrow tem-

perature and concentration range. Considering the

preparation of Bi2223 materials, it has to be noted

that small variations in concentration and/or tempera-

ture can result in quite different phase compositions

and in a signiﬁcant decrease of the volume content

of the Bi2223 phase. The latter exhibits an excess

of the Bi content of about Bi

2.5

. The formation of

the Bi2223 phase is signiﬁcantly promoted by partial

substitution of 20–30% Pb for Bi. Figure 4.2-43 rep-

resents the single-phase region of the BiPb2223 phase

schematically. The BiPb2223 phase contains an excess

of Bi +Pb concentration of about (Bi, Pb)

2.2

.Are-

view of the synthesis of bismuth cuprates is given

in [2.16].

As shown in Fig. 4.2-16 of Sect. 4.2.1, the Bi super-

conductors have a layered perovskite structure which

is composed of superconducting CuO

/Ca stacks (one

stack in Bi2201, two stacks in Bi2212, and three in

Bi2223) sandwiched between insulating BiO and SrO

planes. The structural data of the compounds are pre-

sented in Table 4.2-27.

Due to this layered structure, the compounds exhibit

a more or less pronounced anisotropy of the electrical

parameters (Table 4.2-28).

Here, λ

and λ

are the London penetration depths

for a magnetic ﬁeld in the c-direction and parallel to

the ab-plane, respectively. B

c2(c)

and B

c2(ab)

are ap-

proximate upper critical ﬁelds aligned in the c-direction

and parallel to the ab-plane, respectively; ξ

and ξ

are

the coherence lengths in c-direction and in the ab plane,

respectively. Considerably different values have been re-

ported for the parameters listed in Table 4.2-28. This is

true in particular for H

and also for ξ, which is usually

300

200

100

40 50 60 70 80 90 100 110 120

Temperature (K)

Resistivity (µΩ cm)

200 300

⏐⏐c

Fig. 4.2-45 Resistivity of a Bi2223 ﬁlm as a function of the mag-

netic ﬁeld aligned parallel to the c axis [2.87]

derived from H

. The scatter may bepartly due to differ-

ent chemical composition (doping) and type of samples

(e.g., thin ﬁlm, single crystal, polycrystalline textured

bulk material). Furthermore, different methods of mea-

suring, e.g., in the case of H

, magnetization M(T ) and

resistive measurements ρ(T ), respectively, were used

and various theoretical models were applied for the

determination of the parameters at T = 0 K. Different

criteria were also used for the resistive determination

of H

: ρ(T ) = 0andρ(T ) = 0.5ρ

, respectively. Here

is the normal state resistivity near T

.Sovaluesof

c2(ab)

(0) were determined to 533 T using the criterion

ρ(T ) = 0 and to 2640 T using the criterion ρ(t) = 0.5ρ

for one and the same sample [2.77]. These high values

Part 4 2.2

738 Part 4 Functional Materials

are questionable in general, as has been shown in [2.88].

A further compilation of data is given in [2.89].

Thin Films, Single and Bi-Crystals

There are essentially four methods of preparation sup-

plying high quality BSCCO ﬁlms:

•

Liquid phase epitaxy (LPE) [2.92]

•

Chemical vapor deposition (CVD) [2.93]

•

Ion sputtering [2.94]

•

Pulsed laser deposition (PLD) [2.90]

Epitaxial deposition on monocrystalline ceramic sub-

strates such as Sr

TiO

or MgO is used to prepare

monocrystalline ﬁlms. Investigation on the inﬂuence of

grain boundaries on the superconducting properties of

–1

–2

–3

–4

01234567

B (T)

(B, T)/J

(0, 2.4)

T=60K, B

⏐⏐c

T=60K, B

⊥c

T=4.2K, B

⏐⏐c

T=4.2K, B

⊥c

Fig. 4.2-46 Magnetic ﬁeld dependence of the critical cur-

rent density of an epitaxial Bi2212 ﬁlm for the B

c

and

⊥c

directions at 4.2 K and 60 K, respectively. J

(0, 4.2) is

the critical current density at 4.2 K in zero ﬁeld [2.90]

(θ)/J

(0°)

0 10203040

–1

–3

–5

Tilt angle θ (deg)

YBa

7–δ

0 10203040

Tilt angle θ (deg)

CaCu

8+x

YBa

7–δ

CaCu

8+x

(θ)/J

(0°)

4.2K

77K

Fig. 4.2-48 Ratio of the intergrain, J

(θ) and intragrain critical

current densities, J

(0), as a function of the misorientation angle θ

for bicrystalline [001]-tilt Bi2223 (circles)andYBa

7−x

ﬁlms

(squares) [2.91]

Critical current density(A/cm

)

80246

Magnetic field (T)

B⏐⏐c

20K

26K

30K

40K

45,54K

50K

60K

70K

77K

Fig. 4.2-47 Critical current density J

of an epitaxial

Bi2223 ﬁlm as a function of the magnetic ﬁeld aligned

parallel to the c axis for the temperature range of

20–77 K [2.87]

the ﬁlms’ artiﬁcial grain boundaries within the ﬁlms

have been carried out using bi-crystalline substrates.

Critical temperatures up to 91 K [2.91] and

110.4 K [2.72] were measured on Bi2212 and Bi2223

ﬁlms, respectively. Typical thickness values of such

ﬁlms are 200–400 nm. An example for the magnetic-

ﬁeld dependence of T

is given in Fig. 4.2-45 [2.87]. In

contrast to the conventional low-T

conductors, the onset

temperature of the superconducting transition practi-

cally does not depend on the ﬁeld and a characteristic

enhancement of the transition width with increasing ﬁeld

is observed.

Critical current densities on the order of 10

Acm

−2

at 77 K in zero ﬁeld were measured on BSCCO thin

ﬁlms. In Fig. 4.2-46 the ﬁeld dependence of J

a Bi2212 ﬁlm is shown for two temperatures [2.90] and

Fig. 4.2-47 demonstrates the J

(B, T ) dependence of

a Bi2223 ﬁlm [2.87]. In magnetic ﬁelds aligned parallel

to the plane of the ﬁlm, that is perpendicular to the c axis

of the crystal structure, J

is practically independent on

the ﬁeld strength, even at higher temperatures such as

60 K. The reasons for this behavior are discussed below

in connection with the J

(B, T ) correlation of wires and

tapes.

In connection withthe manufacturing of real conduc-

tors of large length the inﬂuence of grain boundaries on

the critical current density is of great importance. Inves-

tigations of BSCCO bi-crystalline thin ﬁlms have shown

that tilt boundaries act as weak links, causing an expo-

nential dependence of J

on the misorientation angle, θ,

Part 4 2.2

Superconductors 2.2 Non-Metallic Superconductors 739

similar to the J

(θ) dependence of tilt grain boundaries

in YBCO ﬁlms as shown in Fig. 4.2-48 (compare also

with Fig. 4.2-54 of Sect. 4.2.2.1).

Much effort has been expended to grow BSCCO

single crystals using several methods, such as crys-

tal growth from the melt, fused-salt reactions in

alkali-halide and alkali-carbonate ﬂuxes, respectively,

solid-state reaction, and growth in a gaseous phase. Very

high quality bulk single Bi2212 crystals with a size of

a few mm along the a axis and b axis and about 0.1mm

along the c axis were prepared by the travelling sol-

vent ﬂoating zone technique [2.95] or by slow cooling

of the melt below the melting temperature and subse-

quent removing of single crystal fragments from the

ingot by crashing [2.96]. Maximum T

values of 90 K

were measured on these crystals. The critical current

densities of such crystals were very small, on the or-

der of 100 A cm

−2

at 77 K in zero ﬁeld probably due

to the small density of crystal defects. After irradiation

of the crystals with 2.2-G eV Au ion beams directed

along the c axis at a ﬂuence of 1.0×10

−2

,anes-

sential increase of the irreversibility line were observed

at temperatures < 70 K. An enhancement of J

by a few

orders of magnitude was also measured at low tempera-

tures of about 30 K in a ﬁeld of 9 T directed parallel

to the c axis [2.97]. These effects can be attributed to

columnar tracks within the crystal lattice.

It is much more difﬁcult to grow Bi2223 than Bi2212

single crystals. Very small crystals with a phase pu-

rity of 97% Bi2223 and dimensions of 0.1 mm along

the a and b axes and 0.001–0.01 mm along the c axis

could be prepared by the fused-salt reaction method

using a KCl ﬂux [2.98]. The onset temperature of

the superconducting transition was determined by dc

magnetization measurement to 110 K and the transition

width to 5 K. Using the measured magnetization curves

and the “Bean” model, the zero ﬁeld critical shielding-

current density in the ab-plane of the crystal has been

determined to be 1× 10

Acm

−2

at 5 K.

The preparation of bulk Bi2212 bi-crystals pro-

vided the possibility to study the inﬂuence of [001]

twist boundaries (c axis twist boundaries) on the super-

conducting properties of Bi2212 conductors [2.95, 96].

Assuming the results can be applied to Bi2223 tapes too,

pieces of single crystals were cleaved. One cleave was

rotated through an angle about the c axis with respect to

the other, and placed atop it. The boundary was formed

in a controlled sintering process. Results showed that the

transition temperature of a bicrystal agreed with that of

the single crystal (the twist angle is not noted [2.96]) and

the ratio of critical current densities at T/T

≥0.9 across

the twist junction to that across the single crystal was

unity, independent of the twist angle. However, it has

to be noted that in these experiments the critical current

densities were very low and it is not clear whether the

true J

values of the grain boundaries could not be meas-

ured due to the current limiting J

of the single crystals.

Bicrystals with an enhanced critical current density after

irradiation with Au-ion beams revealed, at low tempera-

tures, a more than one order of magnitude higher critical

current across the single crystal than that across the 45

◦

twist grain boundary [2.97].

Practical Conductors

Fabrication of Wires and Tapes. The most common

approach used to fabricate practical BSCCO conduc-

tors is the so called “Powder-in-Tube” (PIT) method.

As a ﬁrst step of PIT precursor, powders of appropri-

ate chemical and phase compositions are produced. In

the case of Bi2212 conductors, usually a mixture of

, SrCaO

, CaO or CaCO

, and CuO with the ratio

Bi:Sr:Ca:Cu equal to 2:2:1:2 is calcined at 800–850

◦

to remove the residual carbon content. The mixture of

oxides and carbonates can be obtained by alternatively

applying mechanical mixing (e.g. milling), or a sol–

gel process, a co-precipitation method, or spray freeze

drying technique. Beside the oxide/carbonate mixture,

a powder produced by spray pyrolysis of a metal-

nitrate solution with subsequent calcination is often

used as an alternate precursor for BiPb2223 conduc-

tors. Typical BPSCCO(2223) precursor powders exhibit

the concentration ratio of Bi:Pb:Sr:Ca:Cu equal to

(1.7–1.8):(0.3–0.4):(1.9–2.0):(2.0–2.1):3.0, and con-

sist of the Bi2212 phase (major phase, about 70%) and

other phases such as Ca

PbO

, (Pb, Bi)

CuO

alkali earth cuprates, (Ca, Sr)

,orCuO,de-

pending on the preparation conditions (particularly on

the ﬁnal heat treatment of the precursor). The chem-

ical and phase compositions of the precursor have

a strong inﬂuence on the ﬁnal properties of the wires

and tapes [2.99].

The further steps of PIT excluding the thermal and

thermomechanical treatments, respectively, are illus-

trated in Fig. 4.2-49. The as-produced powders are ﬁlled

into an Ag tube and then swaged and drawn into wires.

Mostly the conductors are produced as multiﬁlamentary

(MF) wires due to their more uniform superconduct-

ing properties and superior behavior with respect to

the mechanical properties compared to monoﬁlamen-

tary conductors. When making MF conductors, pieces

of the monoﬁlamentary wire are bundled into a second

Ag tube and then this composite is swaged and drawn

Part 4 2.2

740 Part 4 Functional Materials

Deformation

Drawing

Precursor

Bundling

Rolling

Cross section Longitudinal section

Drawing

Fig. 4.2-49 Scheme of the “powder-in-tube” method, ex-

cluding the thermal and thermomechanical treatments

into wires again. In most cases the wires are ﬁnally

rolled into tapes since a higher critical current density

can be achieved for tape-like conductors than for round

wires.

After completion of the deformation process the

Bi2212/Ag and BiPb2223/Ag composites have to be

subjected to a thermal and thermomechanical process,

respectively, in order to transform the Ag-sheathed

precursors into superconducting phases and to real-

ize an optimal microstructure of the superconducting

phase with respect to high critical current densities.

The Bi2212 composites are heated up slightly above the

temperature of partial melting (880–900

◦

C, depending

on the oxygen partial pressure of the annealing atmo-

sphere) and then slowly cooled at a rate of 5–10 K h

−1

to about 830–860

◦

C, followed by holding the com-

posites at this temperature for a few hours. During

this process the Bi2212 phase crystallizes from the

melt as platelets with the crystal c axes aligned per-

pendicular to the wide surface of the platelets. The

microstructure reveals a pronounced degree of texture

with the c axes perpendicular to the ﬁlament/sheath

interface.

A more complicated thermomechanical process that

usually contains three steps is needed in the case of the

BiPb2223/Ag composites. At least the ﬁrst annealing

is performed in a temperature range where a transient

melt occurs, accelerating the conversion of the pre-

cursor into the superconducting phase. The optimum

temperature and dwell time strongly depend on the com-

position of the precursor and oxygen content of the

annealing atmosphere. At a lower O

-content

compared to air, lower temperatures and shorter dwell

times are suitable and less dependence of J

on the

annealing temperature has been observed. For anneal-

ing atmospheres consisting of the inert nitrogen with

an O

content of 7–21%, the optimal temperature is in

the range of 820–835

◦

C, depending on the composition

of the precursor. X-ray diffraction (XRD) investigations

of tapes quenched from different states of the anneal-

ing process as well as in-situ measurements by means

of high-energy synchrotron X-ray diffraction [2.100]

and neutron diffraction [2.101], [2.102], respectively,

have shown that during heating up and dwelling at the

reaction temperature, several reactions such as the in-

corporation of Pb into the Bi2212 phase, the formation

of transient compounds (e.g., (Ca, Sr)

and

(Sr, Bi)

CuO

) and of a liquid phase mentioned

above precede the real conversion of the Bi2212 into

Bi2223. These reactions are accompanied by the re-

crystallization of the Bi2212 phase and improvement

of its original texture degree, which was caused by the

deformation of the composite into wires and tapes.

During the reaction annealing the density of the

BiPb2223 ﬁlaments decreases. Therefore the tapes are

densiﬁed after annealing by a compressive treatment,

whereby the electrical connectivity and degree of tex-

ture of the grains are enhanced. Uniaxial pressing is the

most effective method with respect to J

. If long tapes

are produced, however, the uniaxial pressing has to be

replaced by rolling. The lower J

after rolling, compared

to that after pressing, is probably due to additional mi-

crocracks induced by rolling and can not be recovered

completely by the subsequent annealing which has to

be performed after rolling as well as pressing. Mostly,

a further increase of J

can be achieved after a second

compressive treatment and subsequent third annealing.

Figure 4.2-50 shows the longitudinal cross-section

of a BiPb2223 ﬁlament embedded in silver. The mi-

crostructure of the superconducting phase in contact

with the Ag sheath reveals a higher density and de-

gree of texture compared to areas in the middle of the

ﬁlament. Obviously, Ag inﬂuences the conversion re-

action of Bi2212 into BiPb2223. It is known that Ag

Part 4 2.2

Superconductors 2.2 Non-Metallic Superconductors 741

Fig. 4.2-50 Longitudinal cross section of a BiP2223 ﬁla-

ment embedded in silver

decreases the solidus temperature of Bi2212 [2.103] and

accelerates the BiPb2223 formation process. However,

the reaction mechanism has not yet been fully under-

stood. Two different mechanisms have been proposed:

the nucleation and growth of the BiPb2223 phase from

the melt [2.104] and the formation of the BiPb2223

phase from the Bi2212 grains by intercalation of Ca, Cu,

and O [2.105]. In connection with the ﬁrst-mentioned

mechanism beside the homogeneous nucleation, the het-

erogeneous one is discussed in [2.102]. In the frame of

the intercalation process, the occurrence of a transient

liquid phase also seems to be an important feature.

In order to improve the tensile strength of the con-

ductors, Ag can be replaced by Ag alloys, such as

−

Mg and Ag

−

Mn alloys as sheath material con-

taining a few tenths wt% of the alloying element.

Ternary alloys such as Ag

−

Mg are also used. Due

to the respective internal oxidation of Mg and Mn,

during the long-term annealing necessary to form the

superconducting phase, the tensile stress which can

be applied to the tapes without J

degradation could

be enhanced from 40 MPa for BiPb2223/Ag tapes

to 90 MPa [2.107] and 130 MPa [2.108] for AgMn-

and AgMg-sheathed tapes, respectively. A high ten-

sile strength protects the conductor against damages

by handling and hoop stresses acting on the conduc-

tor due to the Lorentz force in magnetic coils. Silver

and its alloys are indispensable since no other mater-

ials have been found so far which meet the following

requirements necessary for the construction of practical

conductors: high workability; resistance against oxida-

tion at high temperatures; oxygen permeability since an

exchange of oxygen between the precursor and ambi-

ent atmosphere during the heat treatment is necessary;

no detrimental reaction with elements of the supercon-

ducting phase at high temperatures; and sufﬁciently high

tensile strength.

Superconducting Properties. The largest efforts with

respect to the production of practical conductors in long

length have been made with BiPb2223 tapes since they

are the most promising conductors for electric power

applications at present. Figure 4.2-51 shows the crit-

ical current density, J

, of the superconducting phase

of BiPb2223 tapes, at 77 K in self-ﬁeld versus the

tape length achieved by various manufacturers [2.106].

With respect to application, the so-called engineering

critical current density, J

, deﬁned by the critical cur-

rent, I

, divided by the whole transverse cross-section

area (superconductor + sheath) is of great importance.

An example of the performance of long tapes is given

in Fig. 4.2-52, showing J

values at 77 K in self-ﬁeld

for 200-m long tapes produced in one 17 000-m man-

ufacturing run [2.109]. Much higher critical current

densities were measured on experimental short sam-

ples and still higher J

values were observed locally

in tapes, demonstrating the potential of BiPb2223 con-

ductors. In Fig. 4.2-53 the steady improvement of J

of short multiﬁlamentary samples in the 1990s is

shown [2.110]. Values of about 75 kA cm

−2

at 77 K

and 0 T were achieved [2.111]. Magnetooptical inves-

tigations on a monoﬁlamentary tape with an average

(77 K, 0T) = 35 kA cm

−2

revealed some small local

regions carrying up to 180 kA cm

−2

(77 K, 0 T) [2.22].

1000 12000 200 400 600 800

Critical current density (kA/cm

)

Length (m)

NST

ASC

Sumitomo

VA C

IGC

Fig. 4.2-51 Critical current densities J

of long mul-

tiﬁlamentary BiPb2223/Ag tapes produced by various

manufacturers as a function of the length (NST = Nordic

Superconductor Technologies, ASC = American Super-

conductor Corporation, VAC = Vacuumschmelze, IGC =

Intermagnetics General Corporation [2.106])

Part 4 2.2