Pulsed laser ablation: A new route to synthesize novel superconducting
compounds as oriented films
K. I. Gnanasekar, M. Sharon, R. Pinto, S. P. Pai, M. S. R. Rao et al.
Citation: J. Appl. Phys. 79, 1082 (1996); doi: 10.1063/1.360898
View online: http://dx.doi.org/10.1063/1.360898
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Published by the American Institute of Physics.
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Pulsed laser ablation: A new route to synthesize novel superconducting
compounds as oriented films
K. I. Gnanasekar and M. Sharon
Department of Chemistry, Indian Institute of Technology, Bombay 400 076, India
R. Pinto, S. P. Pai, M. S. R. Rao, P. R. Apte, A. S. Tamhane, S. C. Purandare,
L. C. Gupta, and R. Vijayaraghavan
Tata Institute of Fundamental Research, Homi Bhabha Road, Bombay 400 005, India
~Received 3 October 1994; accepted for publication 18 September 1995!
We report here the significance of the pulsed laser ablation technique in stabilizing strained lattices
that do not form by the conventional ceramic method and show that the technique offers unique
possibilities to probe the structure property relationship in complicated systems. One of such
systems is LuBa2Cu3O72d; a systematic investigation of structural ~in!stability of its
superconducting phase is presented here. Our analysis suggests that the system suffers from internal
strain due to lower ionic radius of Lu31; however, the structure can be stabilized only as oriented
films on ^100& LaAlO3 , ^100& SrTiO3 , and ^100& MgO, with excellent superconducting properties
~J c '5.03106 A cm22 at 77 K!. We have also investigated similar compounds having their stability
close to their crystallographic limit. The important feature of these metastable phases is that they
grow only as oriented films. Free energy of epitaxial growth of strained films are investigated and
a simple growth model is proposed based on our observation. Importance of this growth model in
explaining the superconductor–normal-metal–superconductor type of junctions, observed in
high-T c superconductors is highlighted. © 1996 American Institute of Physics.
@S0021-8979~96!01601-3#
I. INTRODUCTION
Complex metal oxides are conventionally synthesized by
the ceramic method which requires high temperature. Metastable phases are usually produced by generating a liquid or
vapor of a well-determined composition and then rapidly
quenching to prevent the second phase nucleation and
growth. A major limitation of this approach is that materials
that are thermodynamically unstable at high temperatures
cannot be prepared by this route. Significantly, several of the
oxide materials of current interest such as zeolite and other
microporous solids as well as most of the high-T c cuprates
are metastable phases. Synthesizing such solids, especially to
specification, is a challenging task.
Several chemical routes have been tried to stabilize the
metastable phases. Of these, sol-gel methods have been successful to a large extent.1 Preparation of glasses and ceramics
by this route permits the mixing of constituents at a molecular level. Owing to the intimate mixing and high surface area
of the gels, these materials can be synthesized at temperatures considerably lower than those of the equivalent compositions prepared by conventional ceramic method. However, in the case of Bi–Sr–Ca–Cu–O or Tl–Ba–Ca–Cu–O
or some doped Y–Ba–Cu–O, the volatile components ~Bi,
Tl, or the dopent! make it difficult to synthesize those compounds as the sintering temperature is only slightly reduced
~.50 K!.2– 6
In recent years, the pulsed laser deposition technique has
been receiving much attention as a viable method for the
deposition of thin films of a variety of materials. The stoichiometric ablation of constituent species from the target
simply makes the PLD technique particularly attractive for
the synthesis of complex multicomponent materials.7–10 Fur1082
J. Appl. Phys. 79 (2), 15 January 1996
thermore, the nonequilibrium nature of the ablation process
offers opportunity to synthesize interesting metastable oxides. In this context, we have illustrated through an example
the significance of PLD technique in exploring superconductivity in metastable compounds which do not form otherwise.
Since the discovery of YBa2Cu3O72d, extensive investigations have been made pertaining to substitution of Y with a
trivalent rare-earth ~RE! ion from the lanthanide series. The
results relating to Lu-1:2:3 phase have been rather controversial. Ku, Hwang, and Tai11 prepared 95% pure, superconducting polycrystalline Lu-1:2:3. Recently Zhou et al.12 reported 90 K T c for a single-crystal specimen. On the other
hand a number of other workers claimed that monophasic
Lu-1:2:3 cannot be synthesized. Hor et al.13 synthesized and
examined by x-ray diffraction most of the RE-1:2:3 compounds but could not prepare Lu-1:2:3. Tarascon et al.14 and
Moodenbough et al.15 could not reliably refine the Lu-1:2:3
lattice parameters and concluded that a second phase accounted for at least half of the powder pattern intensity. On
the other hand Hodorowicz, Hodorowicz, and Elick,16 who
investigated the phase compatibilities in air at 940 °C for the
systems Ln2O3 –BaCO3 –CuO where Ln5Yb and Lu, concluded that Lu-1:2:3 phase cannot be synthesized in air at
940 °C. Somasundaram et al.17 and Balakrishnan, Varadaraju, and Subba Rao18 observed that the difficulties encountered in isolating the Lu-1:2:3 phase are associated with
the small ionic radius of Lu13 ions rather than the preparation conditions. They have concluded that monophasic
LuBa2Cu3O72d cannot be prepared in the range of 900–
950 °C and a lower limit of the rare-earth ionic radius exists,
which the 1:2:3 structure can tolerate.
0021-8979/96/79(2)/1082/10/$6.00
© 1996 American Institute of Physics
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Formation of REBa2Cu3O72d, therefore, depends on the
size and the valence state of the rare-earth ion. The ionic
radius of the rare-earth is critical for the stabilization of the
RE-1:2:3 phase. The largest rare-earth ion that can be tolerated in this structure is La13, having ionic radius 1.14 Å,
while the smallest is Yb13 with ionic radius of 0.975 Å in
eightfold coordination.19 In fact, preparation of
YbBa2Cu3O72d in pure phase becomes difficult as the ionic
radius of Yb13 becomes nearly critical. Further, as mentioned above Lu-1:2:3 does not form in the bulk ceramic
form if one attempts to prepare by the conventional solidstate method.13–18 Although preparation of single crystals12
and thin films of Lu-1:2:3 are reported,20,21 the precise reason for the stabilization of the same is not known. The instability of LuBa2Cu3O72d phase has been suggested to be of
two different origins:
~1! higher thermodynamic stability of Lu2BaCuO5 phase
~211! than the corresponding Lu-1:2:3 phase;16
~2! ionic radius of Lu13 ~0.975 Å in eightfold
coordination19! falls below the critical radius required
for the stability of superconducting orthorhombic
phase.17,18
A systematic analysis of the factors which destabilize the
superconducting phase in the Lu–Ba–Cu–O system is discussed here. A simple model is proposed for the stabilization
of strained lattices in thin films.
FIG. 1. XRD spectra of bulk composition of ~a! LuBa2Cu3O72d, ~b!
Lu0.9Pr0.1Ba2Cu3O72d, ~c! Lu0.9Tb0.1Ba2Cu3O72d.
III. RESULTS AND DISCUSSION
A. LuBa2Cu3O72d
II. EXPERIMENT
Stoichiometric mixtures of Lu2O3 , Pr6O11 , ThO2 ,
Tb4O7 , BaCO3 , SrCO3 , CaCO3 , and CuO of the composition
of Lu12x Mx Ba2Cu3O72d ~0<x<0.2!, where M5Pr, Tb, and
Ca,
and
LuBa22x Srx Cu3O72d
~0<x<0.4!
and
Th0.5Ca0.5Ba2Cu3O72d have been prepared by the conventional solid-state method. The samples were first fired at
900 °C for 24 h, reground and pelletized, and again sintered
at 900 °C for 24 h and the process was repeated three times.
The pellets were tested for homogeneity and used as targets
for pulsed laser ablation after having annealed them in oxygen at 450 °C for 24 h followed by slow cooling to room
temperature.
Thin films of the above composition were deposited on
MgO, LaAlO3 , and SrTiO3 , all with ^100& orientation. A
Lambda Physik model 301 KrF ~248 nm! excimer laser, with
maximum energy of 1200 mJ having pulse width of 25 ns,
was used. The deposition parameters were optimized to realize in situ growth of thin films of best possible quality. The
optimized conditions were 331 mm2 laser spot, 3 J cm22
fluence, 4.5 cm target–substrate distance, and 200 mTorr
oxygen pressure. The film growth time was about 20 min.
Immediately after the run films were flushed with oxygen
and allowed to cool at the rate of 25 °C/min which is the
optimized condition for cooling for films prepared in situ.
Thickness under these conditions was in the range of 1500–
2000 Å. Films were characterized by x-ray-diffraction
~XRD! and the composition was analyzed by energydispersive x-ray ~EDX! analysis.
J. Appl. Phys., Vol. 79, No. 2, 15 January 1996
Shown in Fig. 1~a! is the XRD spectrum of the nominal
composition of Lu-1:2:3. We also have found that
LuBa2Cu3O72d does not form in the bulk. The sample is
multiphasic and insulating and contains only Lu2BaCuO5
and BaCuO2 phases along with some unreacted CuO. No
trace of LuBa2Cu3O72d phase is detected. The mechanism of
formation of RE-1:2:3, prepared by the conventional solidstate method has been well studied. Growth of RE-1:2:3
phase proceeds via RE2BaCuO5 phase ~211!.22–25 Hodorowicz and co-workers16 studied the phase diagram of the Lu–
Ba–Cu–O system and found that the intermediate phase,
namely Lu2BaCuO5 ~211 phase!, has been found to be more
stable than the corresponding Lu-1:2:3 phase for temperatures in the region of 900–950 °C. Therefore, it has been
suggested that this Lu2BaCuO5 could possibly prevent the
formation of superconducting Lu-1:2:3 phase. Hence, a systematic analysis of the relative stability of RE2BaCuO5 phase
across the lanthanides is required. Structural aspects of
RE2BaCuO5 by the neutron-diffraction technique were investigated by several authors.26 –30
The stability of numerous oxide compounds have been
studied based on the valence sum rule ~VSR!.31,32 The phenomenological relation between the bond length and the valence of a bond can be expressed as31,32
d i j 5exp@~ R 0 2R i j ! /B # ,
~1!
where B50.37. R 0 is characteristic of cation anion pair. The
valence sum rule establishes that the sum of valence bonds
around a cation must be equal to the formal valence V i of the
ith cation ~i.e., ( d i j 5V i !.
Gnanasekar et al.
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1083
FIG. 2. Variation of GII as a function of RE ion for RE2BaCuO5 after
References 26–30.
One of the main bases of the VSR is that the structure
should permit the release of the stress introduced by the coexistence of different structural units. Applying VSR to
RE2BaCuO5 by using the values obtained from Refs. 26 –30,
a clear deviation of the valence sum around each ion with
respect to the expected value is observed. This gives a clear
evidence of possible instabilities in the crystal structure. The
root mean square of the bond valence sum deviation for all
the atoms present in the asymmetric unit is a measure of the
extent to which the VSR is violated over the whole structure
and the value is called global instability index ~GII!,33,34
GII5 A( ni51 @ ( nj51 ~ d i j 2V i ! 2 # /N.
~2!
A systematic variation of structural stability of 211 phase
across the lanthanides is correlated with the size of the rareearth ion. It is observed that minimum GII values ~higher
stability! are obtained for the Er, Tm, and Yb oxides ~Fig. 2!.
When the ionic radius of the lanthanide cation increases, the
instability index goes up taking its maximum value ~less
stable! for Sm. This could justify the fact that for the
Nd2BaCuO5 and La2BaCuO5 oxides, having Nd31 and La31
with larger ionic radius than Sm31, this structure does not
exist. On the other hand, for the smallest lanthanide cation,
Lu31, with an ionic radius smaller than that of Yb31 the
instability index starts increasing which is indicative of a
larger stress in the structure.
As the instability index of Yb-2:1:1 phase is less than
~more stable! that of Lu-2:1:1 phase, it suggests that Yb1:2:3 must be relatively more difficult to form than the Lu1:2:3; but, experimentally, it has been found the other way
round. To check whether the stability of 211 phase that prevents the formation of Lu-1:2:3 phase, we have studied the
substitution effect of Pr, Tb, and Ca regarding the structural
and stability aspects of the Lu-1:2:3 system.
The structural instability of the LuBa2Cu3O72d phase has
been probed by following chemical substitutions. Although
pure Lu-1:2:3 phase does not form in the bulk, partial substitution of even 10% Pr for lutetium stabilizes the superconducting phase to a large extent. The XRD spectrum of
Lu0.9Pr0.1Ba2Cu3O72d is shown in Fig. 1~b!. Since the ionic
radius of Pr13 ~1.13 Å! is larger than that of Lu31, the effective ionic radius of the system is higher than the critical
radius and, hence, should have allowed the synthesis of su1084
J. Appl. Phys., Vol. 79, No. 2, 15 January 1996
FIG. 3. Plot of normalized resistance @R(T)/R~300 K!# as a function of
temperature for ~a! Lu0.9Pr0.1Ba2Cu3O72d, ~b! Lu0.9Tb0.1Ba2Cu3O72d ~bulk
samples!.
perconducting phase. Figure 3~a! shows the plot of normalized resistance as a function of temperature. The compound
is highly metallic and superconducts with a T c of 86 K. We
have chosen deliberately Pr because PrBa2Cu3O72d does not
superconduct but would stabilize the superconducting phase
in the Lu-1:2:3 system.
Although even partial substitution of terbium or cerium
is not possible in RE-1:2:3 phase because of the formation of
BaTbO3 and BaCeO3 , respectively, partial substitution of terbium in the Lu-1:2:3 system yielded interesting results. Figure 1~c! shows the XRD spectrum of Lu0.9Tb0.1Ba2Cu3O72d
with small amounts of impurity phases. As Tb does not form
the RE-1:2:3 phase, it cannot act as a nucleation center for
the growth of the Lu-1:2:3 phase. Figure 3~b! shows the
temperature dependence of its resistance. Therefore, it appears that the Lu-1:2:3 phase should suffer from internal lattice strain because of smaller ionic radius of Lu13 and that
terbium helps in stabilizing the phase by relieving the strain.
Formation of superconducting phase can happen only if terbium substitutes as a trivalent ion in the Lu-1:2:3 lattice
since tetravalent terbium ion has ionic radius less than the
critical radius. This conclusion is supported by the fact that
cerium could not stabilize the superconducting phase in the
Lu-1:2:3 system as cerium substitutes as a tetravalent ion and
the ionic radius of Ce41 is less than the critical radius.
One may argue that it is the ambivalent nature of cerium
which destabilizes the Lu-1:2:3 phase rather than ionic radius. Therefore, we have prepared Lu12x Cax Ba2Cu3O72d and
found that even the Ca12 ion stabilizes the Lu-1:2:3 phase
when partially substituted. In this case, the effective ionic
radius of the system increases due to the large ionic radius of
Ca12 which allows the stabilization of the Lu-1:2:3 phase in
the bulk @Fig. 4~a!#. These studies have given sufficient evidence that it is not the stability of the 211 phase which prevents the formation of the Lu-1:2:3 phase. On the contrary,
the formation of the 211 phase results due to the instability of
superconducting Lu-1:2:3 phase in the Lu–Ba–Cu–O system.
Finally, to support our conclusion, we have prepared
LuBa22x Srx Cu3O72d. Here, we have replaced the larger Ba
ions by smaller Sr ions. Even for x50.2, we could get reaGnanasekar et al.
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FIG. 6. J c as a function of temperature for Lu-1:2:3 film grown at 700 °C.
FIG. 4. XRD spectra of bulk composition of ~a! Lu0.9Ca0.1Ba2Cu3O72d, ~b!
LuBa1.8Sr0.2Cu3O72d, ~c! thin film of LuBa2Cu3O72d.
sonably pure phase ~.85%! @Fig. 4~b!#. The obvious conclusion is that the unit cell of LuBa2Cu3O72d is slightly larger so
that it gets strained to accommodate the smaller Lu13 ions;
however, if the cell volume is reduced by partial substitution
for larger Ba12 ions by smaller Sr12 ions superconducting
LuBa2Cu3O72d is stabilized.
On the other hand, the LuBa2Cu3O72d phase is stabilized
in thin films as c-axis-oriented films. Figure 4~c! shows the
XRD spectra of the Lu-1:2:3 films deposited at 700 °C on a
^100& LaAlO3 substrate. Films are highly metallic with the
ratio R~300 K!/R~100 K! close to three. They superconduct
with T c of 86 K. The deposition rate of Lu-1:2:3 film is
found to be very low and as a result it was difficult to grow
films with thickness more than 2500 Å. In contrast to Y-1:2:3
films, Lu-1:2:3 films grown at 600 °C are also found to be
highly c-axis oriented ~Fig. 5!.
Figure 6 shows the plot of J c as a function of temperature for one of the films grown at 700 °C on ^100& LaAlO3 .
FIG. 5. XRD spectrum of LuBa2Cu3O72d film grown at 600 °C.
J. Appl. Phys., Vol. 79, No. 2, 15 January 1996
The value of J c at 77 K is about 5.03106 A cm22, and at 10
K it is about 3.63107 A cm22 which compares favorably
with that of well-formed Y-1:2:3 films. However, the values
of J c for one of the best films prepared at 600 °C is about 105
A cm22 at 77 K while at 10 K it is around 9.03105 A cm22
~Fig. 7!. These values are an order of magnitude lower than
those of films prepared at 700 °C despite the fact that both
are highly c-axis oriented.
Various mechanisms that limit the current densities in
the case of conventional type-II superconductors, such as
magnetic flux creep and weak links of different nature, have
also been proposed for high-T c cuprates.35,36 Because of the
relatively small pinning energies ~'100 meV! compared to
the large superconducting temperature range ~'100 K!, it
has been suggested that the critical currents in high-T c superconductors should be limited by flux creep, as described
by the Anderson–Kim creep model.37–39 From Maxwell’s
equation the value of the field at the surface of sample at a
given current density j is given by40
H s 54 p jl/c,
~3!
where l is the magnetic penetration depth. When H s becomes larger than the lower critical magnetizing field H c1 ,
FIG. 7. J c as a function of temperature for Lu-1:2:3 film grown at 600 °C.
Gnanasekar et al.
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1085
flux will start to penetrate the sample. Thus, the critical current density is limited by the upper and lower magnetizing
fields, which gives
cH c1 /4p l,J c ,cH c2 /4p l.
Substituting the values of H c1 ~0!'250 G,
H c2 ~0!'50 000 G, and l~0!'150 nm, the above expression
becomes40
107 A cm22 ,J c ,1010 A cm22 .
In the case of epitaxial films, although the critical current
densities are already high ~J c .106 A/cm2 at 77 K!, based on
this model it might be still possible to increase the J c further
by introducing defects with large pinning energies. However,
the main difficulty is that in high-T c superconductors, because of extremely small coherence length the planar defects,
such as grain boundaries, twin boundaries, etc., that is, those
which provide the most effective pinning centers, form regions where the wave-function phase coherence is disturbed.
On the other hand, it is well known that the critical current
densities across the grain boundaries are limited by poor coupling across the boundaries and not by the pinning properties
of the boundaries.37 In fact, evidence indicates that the intragrain J c is controlled by flux creep whereas the intergrain J c
depends sensitively on properties of grain boundaries. Looking into the microstructural aspects of the films prepared at
various temperatures has given interesting results. Figure 8
shows the atomic force micrographs of Lu-1:2:3 films deposited at 600 and 675 °C. For comparison, we have taken the
atomic force microscopy ~AFM! micrograph of bare LaAlO3
substrate. It is clear that the films grown at 600 °C contain
poorly aligned oriented grains with a size of about of 200
nm. On the other hand, the grain size of films deposited at
675 °C is about 600 nm and the grains are better aligned.
Similar observation is confirmed from scanning electron microscopic studies on these films after partial etching. Figure
9 shows the scanning electron micrographs of films deposited at 600 and 700 °C. These studies give sufficient evidences that not only the films contain oriented grains but also
their enlargement and alignment improves with growth temperature. This is crucial for improving the transport properties. In fact, recent reports show that even the highly oriented
thin films also contain oriented grains.41– 44 We also have
found evidences for granular growth in the present case. The
use of higher substrate temperature also reduces the thickness of the grain boundary due to densification and better
alignment of grains. Therefore, it becomes necessary to
know about the physical properties of grain boundaries in
order to improve the transport properties of these films.
Shown in Fig. 10 is the schematic representation of
grains in films, prepared at 600 and 700 °C respectively. The
number of grain boundaries over the specified area for films
prepared at 600 °C is more than that for films prepared at
700 °C as indicated in Fig. 10, and hence, one could expect
the transport behavior of the former to be limited by larger
grain boundaries. As it has been observed that out-diffusion
of oxygen preferentially takes place via grain boundaries,45,46
it is reasonable to assume that the grain boundaries to be
composed of oxygen deficient 1:2:3 phases. That is, although
1086
J. Appl. Phys., Vol. 79, No. 2, 15 January 1996
FIG. 8. Atomic force micrographs of Lu-1:2:3 films grown at ~a! 600 °C
and ~b! 675 °C; ~c! bare LaAlO3 substrate.
the bulk of the grains superconducts with T c of 90 K, a few
layers of grains close to the grain boundary have less oxygen
content than the grains as a result of out-diffusion. T c , which
is a function of oxygen stoichiometry, also changes if one
goes from the boundary to the grain. That is, close to the
grain boundary, the grain is assumed to be composed of several layers each of thicknesses dx 1 , dx 2 , dx 3 , etc., with the
superconducting transition temperatures T 1 , T 2 , T 3 , etc.,
such that T 1 .T 2 .T 3 because of out-diffusion of oxygen
from the boundary. Such an assumption cannot be ruled out
as we have observed a tail which changes with respect to the
current in the resistance verses temperature plot. We have
Gnanasekar et al.
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FIG. 10. Schematic representation of granular nature of films prepared at ~a!
600 °C, ~b! 700 °C. ~c! Schematic representation of grain boundary with
variable oxygen content and transition temperature (T c ).
FIG. 9. Scanning electron micrographs of Lu-1:2:3 films prepared at ~a!
600 °C, ~b! 675 °C, and ~c! 700 °C.
studied high-quality films ~that is, those films having J c
greater than 53105 A cm22 at 77 K! after patterning them
into microbridges of widths 10–50 mm; however, since the
films grown at 600 °C are very poor in transport behavior,
they were studied without patterning them into small geometries. Figure 11 shows the resistance as a function of temperature for the temperature interval of 75–105 K of the film
grown at 600 °C for various current levels. This observation
is similar to that reported by Paracchini.47 The inset of Fig.
11 shows the plot for the full range. The film is metallic and
superconducts at 84 K; however, when the current that we
used to test superconductivity is changed, we have found that
the tail in the R – T plot increases with increasing current
~Fig. 12!.
Similar studies were performed on the films prepared at
700 °C. In order to study more effectively these films were
patterned into microbridges of varying widths. Figure 13
shows plots of resistance as a function of temperature for
J. Appl. Phys., Vol. 79, No. 2, 15 January 1996
various currents for the interval of 75 and 105 K for one of
the patterned films with a microbridge of 10 mm. The inset of
Fig. 13 shows one of the R – T plots for the full range. This
film is highly metallic when compared to the one prepared at
600 °C. Although this film superconducts at 86 K, we observe the tail which varies with the magnitude of the current
that we used. Figure 14 shows the expanded R – T plots of
the same film.
Figure 10~c! shows the schematic representation of a
grain with variable oxygen content and marked with corresponding T c . If such a grain is cooled, the region inside
starts superconducting at 90 K whereas the layers marked as
FIG. 11. Plots of resistance as a function of temperature for Lu-1:2:3 film
grown at 600 °C for various currents between the temperature interval of
75–105 K. Inset shows one of the plots for the full range.
Gnanasekar et al.
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1087
FIG. 14. Expanded versions of R – T plots shown in Fig. 13.
FIG. 12. Expanded version of R – T plots shown in Fig. 11.
T 1 ,T 2 ,T 3 do not superconduct at 90 K ~T 1 .T 2 .T 3 and so
on!; however, they are still metallic. As the temperature is
further lowered, dx 1 starts superconducting whereas dx 2 remains metallic. Further lowering of temperature would make
dx 2 superconducting. In each case, the outer layer exhibits
metallic properties when the previous layer turns out to be
superconducting. Therefore, near T c these layers should exhibit metal-like junctions until all such layers become superconducting. Figure 15 shows the plots of J c as a function of
temperature near T c for the films grown at 600 and 700 °C
which give evidence for the metal-like junctions in these
films. Degennes has developed an expression for the supercurrent that can flow in superconductors having metallic
regions;48 however, the simplified equation for
superconductor–normal–metal–superconductor ~SNS! has
been derived by Clarke.49 The maximum current that can
flow across the SNS junction is given by
I 0 } u F 0 u 2 @ j n ~ T ! / j 2 ~ T !# exp@ 2a/ j n ~ T !# ,
~4!
where jn is the normal-metal coherence length, j is coherence length of the superconductor, F 0 (T) is the bulk value of
FIG. 13. Plots of resistance as a function of temperature for Lu-1:2:3 film
grown at 700 °C for various currents between the temperature interval of
75–105 K. Inset shows one of the plots for the full range.
1088
J. Appl. Phys., Vol. 79, No. 2, 15 January 1996
the density of the superconducting electron pairs and a is the
thickness of the normal metal barrier. The above expression
can be reduced to
I 0 } ~ 12t ! 2 exp@ 2a At/ j n ~ T !# ,
~5!
where t5T/T c . As jn has the weak temperature dependence,
near T c the above equation further reduces to
I 0 } ~ T c 2T ! 2 .
~6!
A plot of AJ c as a function of temperature (T c 2T) is
shown in Fig. 16. Near T c , i.e., between T c and ~T c 215!, all
the films show linear behavior, consistent with SNS junctions. In this model, the width of the grain boundary a is
assumed to be constant throughout. This is valid for conventional and intermetallic superconductors; however, on the basis of these studies, for high T c oxides we postulate here that
near T c the width of the grain boundary varies with temperature.
In this context, we would like to comment on addition of
Ag to YBa2Cu3O72d which has been reported to improve the
transport behavior of both thin films and bulk. It has been
shown that silver goes into the grain boundary and remains
in elemental form and hence behaves as metal-like
junctions.50 On the other hand, although Sb2O3 is also reported to increase the current density of bulk samples, in thin
FIG. 15. The plot of J c as a function of temperature near T c for films grown
at 600 and 700 °C.
Gnanasekar et al.
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FIG. 16. Plot of AJ c as a function of (T c 2T) for the films grown at 600 and
700 °C.
films J c decreases gradually as the percentage of Sb2O3 increases. It has been shown that at higher concentrations
Sb2O3 also goes into grain boundaries and remains as an
insulating oxide.51,52 In small concentrations, however, due
to improved grain alignment the plot of AJ c as a function of
temperature is linear thus giving evidence for the presence of
SNS-type junctions. All these experimental observations are
consistent with our model.
Finally, we briefly discuss the stabilization of these type
of strained lattices ~metastable phases! in thin-film form.
Since the film growth begins with the generation of threedimensional islands, the total free energy of the substrate
system DG is given by the sum of the concentrations of
various size islands G i times their formation energy and an
entropy of mixing term DS mix arising from the distribution of
islands on the available n 0 substrate sites per unit area,53,54
DG5
( n i G i 2TDS mix .
~7!
The free energy of formation of the island depends on
the surface energy given by the relation
s s 5 s i 1 s 0 cos u ,
~8!
where ss is surface energy of the substrate, si is surface
energy of overgrowth, and s0 is the surface energy of the
interface.
The energies of the free surfaces of the epitaxial and
misaligned nuclei will be very nearly equal; however, the
energy of the interface between a well-aligned nucleus and
the substrate will be smaller than that of the energy of the
interface between the poorly aligned nucleus and the substrate. The fact that Lu-1:2:3 is stabilized only as oriented
films implies that it is this interfacial energy that provides the
stability and helps the Lu-1:2:3 grains grow epitaxially.
There exists an elastic strain which operates across the interface of epitaxially aligned films to reduce the difference between its normal lattice parameter and that of its substrate.
For strong interfacial bonding, this elastic strain accommodates all of the misfit up to the thickness of several hundred
angstroms thickness. Because of strong interfacial bonding
observed in the case of Y-1:2:3–type films, we suggest that
J. Appl. Phys., Vol. 79, No. 2, 15 January 1996
FIG. 17. XRD spectra of thin film of ~a! Th0.5Ca0.5Ba2Cu3O72d, ~b!
Lu0.5Tb0.5Ba2Cu3O72d.
this elastic strain should also accommodate the strains due to
inherent instability of the LuBa2Cu3O72d lattice and allow
the stabilization of the Lu-1:2:3 phase.
B. Related materials
It is well known that all the rare earths except Pr, Tb, and
Ce stabilize in superconducting RE-1:2:3 phase. Pr forms
1:2:3 phase but it is semiconducting and antiferromagnetic
and, hence, sheds light on the mechanism of high-T c
superconductivity.55 Even partial substitution of Pr in
REBa2Cu3O72d results in systematic depression of
superconductivity.56 – 65 Several mechanisms have been proposed to account for the systematic depression of superconductivity in RE12x Prx Ba2Cu3O72d and most of them centered around the valence state of Pr which is under
controversy.66 –76 In this context, observation of superconductivity in Pr0.5Ca0.5Ba2Cu3O72d ~.35 K! has been interpreted in favor of a possible higher valence state for Pr,75
although most of the spectroscopic measurements showed a
trivalent state. With this in view we explored the superconductivity in Th0.5Ca0.5Ba2Cu3O72d, where thorium is expected to be in tetravalent state with much extended 5 f orbitals for hybridization with the Cu–O band.
Bulk samples were insulating and there is no trace for
superconducting phase; however in thin films they form as
highly a-axis-oriented films @Fig. 17~a!#. These films showed
T c onset at 44 K, followed by a sharp transition at 39 K and
a tail extending to 16 K before showing zero resistance ~Fig.
18!.
Similarly we have investigated superconducting properties of Lu12x Tbx Ba2Cu3O72d. The bulk samples were multiphasic; however, they form highly c-axis-oriented films
@Fig. 17~b!#. These films superconduct at about 88 K. They
have very high current densities, J c .33106 A cm22 at 77 K.
Although both Tb and Lu do not form respective 1:2:3
phases in bulk, they are stabilized as oriented films. These
films are important because Tb, just as Pr also exhibits mixed
valence characteristics in simple oxides. X-ray photoelectron
spectroscopic analysis of Tb ~4d level! in these films has
Gnanasekar et al.
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FIG. 18. Plot of normalized resistance @R(T)/R~300 K!# as a function of
temperature for ~a! Th0.5Ca0.5Ba2Cu3O72d, ~b! Lu0.5Tb0.5Ba2Cu3O72d.
indicated that Tb is in trivalent state.76 Further, as the redox
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which the latter depresses T c must be of trivalent origin.
IV. CONCLUSION
LuBa2Cu3O72d does not stabilize in the bulk ceramic
form due to smaller ionic radius of Lu13 which just falls
short of the critical radius required to crystallize in
REBa2Cu3O72d. On the other hand, stabilization of the superconducting phase in thin films is provided by
~a!
~b!
free energy of formation at the interface between the
film and the substrate and
the elastic strain at the substrate–film interface which
accommodates lattice misfit up to 10%, as well as the
strain due to the inherent instability of the Lu-1:2:3
lattice.
Thus pulsed laser ablation technique offers unique possibilities in terms of stabilizing metastable phases and hence exploring the structure-property relationship in complicated
systems.
ACKNOWLEDGMENTS
The authors would like to thank C. P. D’Souza for his
assistance in experiments. One of the authors ~K.I.G.! would
like to thank the Council of Scientific and Industrial Research ~CSIR!, New Delhi, for providing financial assistance.
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1091