High-purity FeSe1-x superconductors prepared by solid-state synthesis and liquid phase processing

Article abstract of DOI:10.1111/j.1551-2916.2010.03912.x

FeSe_1-x samples have been prepared by a solid-state reaction in the Ar13% H₂ gas flow. The samples sintered at 410°C were of the tetragonal PbO structure (β-phase), showing little trace of d-phase (hexagonal). However, the samples sintered at higher temperatures contained a fair amount of d-phase, which was unable to be eliminated completely by subsequent annealing at r450°C for a prolonged period. Both M-T and R-T measurements showed that sintered samples of β-phase were superconducting with an onset T_c above 8.1 K. Thick films of FeSe_1-x were grown on LaAlO₃ substrates by liquid phase processing with SeSn as the flux. In contrast to the sintering, the filmsgrown around 910°C from the Sn-containing liquid had a pure β-phase. Although the phase diagram showed some solid solubility of Sn in FeSe, energy-dispersive X-ray spectroscopy failed to find Sn trace in the crystallized films, which showed a steeper superconducting transition at lower T_c of 6.0 K. To clarify whether the possible Sn contamination caused the drop of T_c, FeSe _1-x:5% Sn samples were sintered, which showed little change of T _c from the pure β-FeSe_1-x. The cause for the T _c reduction of thick films remains unclear.

Full text of DOI:10.1111/j.1551-2916.2010.03912.x

J. Am. Ceram. Soc., 93 [10] 3195–3200 (2010)
DOI: 10.1111/j.1551-2916.2010.03912.x
r 2010 The American Ceramic Society
ournal
J
High-Purity FeSe1 Superconductors Prepared by Solid-State Synthesis
Àx
and Liquid Phase Processing
w
Xiaoding Qi, Jiun-Yi Wang, Chi-Jung Hung, and Jui-Chao Kuo
Department of Materials Science and Engineering, National Cheng Kung University, Tainan City 70101, Taiwan
Karen Yates and Lesley Cohen
Department of Physics, Imperial College, London SW7 2BP, U.K.
FeSe1 samples have been prepared by a solid-state reaction in
Àx
ture is very similar to d and can be regarded as the defect ordered
phase of d. Efforts to prepare samples of pure b-phase have not
been very successful so far. Most samples shown in recent pub-
7
the Ar13% H gas flow. The samples sintered at 4101C were of
2
the tetragonal PbO structure (b-phase), showing little trace of
d-phase (hexagonal). However, the samples sintered at higher
temperatures contained a fair amount of d-phase, which was
unable to be eliminated completely by subsequent annealing at
r4501C for a prolonged period. Both M–T and R–T measure-
ments showed that sintered samples of b-phase were supercon-
lications contained some amounts of d-phase and/or Fe
well as the traces of metallic a-Fe and iron oxides (Fe
The samples in these publications were usually prepared by sinte-
7 8
Se , as
2,8–12
O ).
3 4
2
,8–12
ring at over 7001C, followed by annealing at about 4001C.
6
This is not the best route suggested by the Fe–Se phase diagram,
which shows a peritectoid reaction at 4571C that will unavoidably
result in certain amount of d-phase and/or Fe Se .
ducting with an onset T above 8.1 K. Thick films of FeSe
c
were grown on LaAlO substrates by liquid phase processing
1Àx
3
7
8
with SeSn as the flux. In contrast to the sintering, the films
grown around 9101C from the Sn-containing liquid had a pure
b-phase. Although the phase diagram showed some solid solu-
bility of Sn in FeSe, energy-dispersive X-ray spectroscopy failed
to find Sn trace in the crystallized films, which showed a steeper
In this paper, we report on the low-temperature synthesis of
high-purity b-FeSe1Àx in the flowing Ar13% H atmosphere by
a solid-state reaction. In addition, liquid phase processing was
2
also used to grow thick films of b-FeSe1 on the LaAlO3
Àx
(LAO) single crystal substrates with SeSn as the flux. Other
than for the pure b-phase samples, the use of this method, which
was a process similar to single crystal growth, had another pur-
pose that was to grow high-purity Sn-doped FeSe. As shown in
superconducting transition at lower T of 6.0 K. To clarify
c
whether the possible Sn contamination caused the drop of Tc,
FeSe1 :5% Sn samples were sintered, which showed little
Àx
1
3
change of T from the pure b-FeSe . The cause for the Tc
the FeSeÀSeSn phase diagram, there exists a solid solution
c
reduction of thick films remains unclear.
1Àx
region for Fe1ÀxSeSn , which means that the single crystals or
x
polycrystalline solids precipitated from the liquid will have the
samples with Sn truly doped in the FeSe lattice, rather than
presented as the aggregates of Sn-enriched phases, as it is often
seen in many doped powder samples.
I. Introduction
ECENT discovery of superconductivity in the compounds
containing a magnetic element of iron has brought about a
R
1–5
II. Experimental Procedure
new wave of interest in this field. Intensive researches world-
wide have now driven the superconducting transition tempera-
ture to as high as 56 K. Of great interest to the research
The FeSe samples were prepared from the powders of metallic
Fe (99.9% 1–9 mm, Stern Chemicals, Newburyport, MA) and Se
99.5% 1200 mesh, Acros, Fair Lawn, NJ). The powders were
weighed according to the ratio Fe:Se 5 1:1Àx (x 5 0À0.20) and
then mixed by grinding in an agate mortar. For a better mixing
and also avoiding absorption of ambient moisture, water-insol-
uble toluene was added to the mixture during the grinding. The
dried powder mixture was then pressed into the disks of thick-
ness of 2 mm and diameter of 10 mm. The sintering was carried
out in an air-tight quartz tube furnace, which was preevacuated
3
community is that these iron-based compounds neither resem-
ble the conventional superconductors nor behave in the same
way as the high-T
(
4
cuprates, and therefore provide a new op-
c
portunity for the better understanding of unconventional super-
conductivity. In this family, FeSe has the simplest composition
and structure, but yet possesses all the electronic, magnetic, and
structural effects relevant to the superconductivity in the com-
5
plex iron pnictides. Therefore, it will play an important role for
the research into this new class of superconductors.
to 0.1 atm and flushed with the Ar13% H gas for several times.
2
There are several compounds of the composition around the
stoichiometric ratio 1:1 in the FeÀSe binary system, as indicated
To ensure extremely low oxygen partial pressure (pO ) inside the
quartz tube, the furnace was kept at 1001C for 4 h with the
Ar13% H flowing at the rate of 60 sccm. The exit gas from the
2
2
6
in the phase diagram. The superconductivity is only found for the
2
Se-deficient phase FeSe1Àx, which has the tetragonal PbO struc-
quartz tube was checked using a zirconia oxygen trace analyzer
ture with the space group P4/nmm and is termed as the b-phase.
(Rapidox 2100, Cambridge Sensotec Ltd., St. Ives, Cambs,
U.K.), which showed that the p was approaching the detect-
ing limit of the unit (10
The second phase (d-phase) is a Se-rich phase, FeSe11x, having
O2
6
the hexagonal P6 /mmc lattice isostructural to NiAs. The third
À20
ppm). The temperature was then
3
7 8
phase is actually a stoichiometric compound, Fe Se , whose struc-
raised slowly from 1001 to 3001C (501C/h) and dwelled for 2 h.
This was intended to avoid preferential loss of Se before any
chemical reaction occurred, because element selenium has a very
high vapor pressure above its melting point of 2211C. The tem-
perature was finally raised to the desired sintering temperatures
M. Parans Paranthaman—contributing editor
(
ꢀ 4101C) for the chemical reaction to complete. The last step
Manuscript No. 27287. Received December 22, 2009; approved May 4, 2010.
Author to whom correspondence should be addressed. e-mail: xqi045@mail.ncku.
edu.tw
w
was repeated at least once after the samples were crushed and
ground again.
3
195

3
196
Journal of the American Ceramic Society—Qi et al.
Vol. 93, No. 10
Liquid-phase processing was carried out using SeSn as the
flux, which was prepared from the metallic powders of Se and Sn
99.80% 325 mesh, Acros) in the same way as described above.
The final sintering temperature for SeSn was 7001C and the
SeSn powders obtained were checked by X-ray diffraction
Also, changing the nominal composition in the range x 5 0À0.2
did not have a notable result on reducing the amount of Fe Se
but it was found that the amount of a-Fe could be substantially
reduced when the Se content was increased up to FeSe0.97
7
8
,
(
.
However, efforts to eliminate the a-Fe trace completely by fine
tuning of the composition were unsuccessful. It was suspected
that the trace of a-Fe might result from the trace of some Fe O
(
exactly with a known SeSn phase in the database (PDF #75-
XRD) to be pure, whose diffraction patterns matched almost
3
4
1
4
843). A well-mixed FeSe1SeSn pellet with the weight of
1
about 2 g and composition of 10 at.% SeSn was placed on
in the starting iron powder, which was reduced back to a-Fe in
the extremely low pO2 atmosphere at a later stage of the sinte-
ring. Indeed, Fe O s were found as a tiny trace in the samples
top of a (001) LAO substrate measuring 5 mm  5 mm, which
3
4
was placed in the lower end of a tilted alumina (Al
2 3
O ) boat.
sintered at lower temperatures. Figure 1(b) is a typical XRD
pattern for the samples of the composition FeSe0.97, which were
The boat was then loaded into the quartz tube and heated up to
101C, at which the pellet was melted to form a liquid according
9
sintered at 4101C for 24 h. Except for the trace of Fe
#75-0033), the samples had a nearly pure b-phase (PDF #85-
0735) with barely observable Fe Se . Longer sintering time at
3 4
O (PDF
1
3
to the FeSeÀSeSn phase diagram. After being kept at 9101C
for 15 h, the furnace was cooled down slowly at 21C/h to 8001C
and then quickly to room temperature at 6001C/h. When taken
out from the furnace, the LAO substrate was covered with a
thick film precipitated/crystallized from the Fe–Se–Sn high-tem-
7
8
4101C appeared not to be necessary because XRD did not show
any significant change in the samples sintered for up to 7 days,
apart from that the trace of oxide was reduced to a-Fe.
6
perature solution, which also wet a large area of the Al
O
2 3
boat.
The results agreed well with the Fe–Se phase diagram. At
Finally, the LAO substrate was cleaved off from the boat by a
sharp blade.
The structure, composition, morphology, and superconduct-
ing properties of the samples obtained were characterized by a
range of techniques including powder XRD, differential scan-
ning calorimetry (DSC), energy-dispersive X-ray spectroscopy
7501C, the FeSe1Àx (x 5 0–0.2) samples were in the aÀFe1d
two phase region. Upon cooling through 4571C, although the
majority of d-phase undertook a peritectoid reaction with a-Fe
to form the b-phase, i.e. aÀFe1d2b, some d and/or a-Fe re-
mained even for those samples of the composition within the b
6
single phase range (1Àx 5 0.961–0.976), due to segregation or
(
EDX), scanning electron microscopy (SEM), electron backscat-
incomplete diffusion. Therefore, annealing at about 4001C for
the samples sintered at 7501C would not be able to get rid of the
d-phase.
ter diffraction (EBSD), SQUID magnetometry, and four-point
electric transport measurements. XRD patterns were taken us-
ing a Rigaku MultiFlex diffractometer (Tokyo, Japan). DSC
was performed using a TA Q2000 calorimeter (TA Instruments,
New Castle, DE). SEM, EDX, and EBSD were carried out us-
ing a JEOL JSM-7001F microscope (Tokyo, Japan). SQUID
magnetometry was carried with the Quantum Design MPMS
SQUID-VSM system (San Diego, CA). Four-point transport
measurements were performed with a home-made unit.
The structure of the d-phase was very similar to that of Fe Se
7
8
and it was very difficult to distinguish the powder XRD patterns
of d (PDF #75-0608) and Fe Se (PDF #72-1356). However,
7
8
DSC analysis on the samples, which showed the trace of d (or
Fe Se ) in XRD, agreed with the Fe–Se phase diagram that
7
8
III. Results and Discussion
(1) Preparation by Solid-State Reaction
Figure 1 shows the XRD y–2y scans of the samples sintered at
various temperatures. The samples of the nominal composition
FeSe0.85 and sintered at 7501C for 24 h contained a fair amount
of Fe
ing down to room temperature, as shown in Fig. 1(a). Prolonged
8
annealing at 4001C up to 72 h was unable to get rid of Fe Se .
7 8
Se (PDF #72-1356) and a-Fe (PDF #87-0721) after cool-
7
Fig. 1. X-ray diffraction yÀ2y scans of various samples: (a) powders
sintered at 7501C for 24 h and then followed by annealing at 4001C for
up to 72 h, (b) powders sintered at 4101C for 24 h, (c) large grains grown
on an alumina boat from liquid at 9101À8001C. The indexed peaks be-
Fig. 2. The differential scanning calorimetry graphs for the samples: (a)
with and (b) without the trace of Fe Se on X-ray diffraction.
long to the superconducting b-FeSe1Àx
.
7
8

October 2010
High-purity FeSe1 Superconductors
Àx
3197
6
there was an eutectoid reaction: d2b1Fe
as shown in Fig. 2(a). The heating/cooling rate for DSC was
1C/min, indicating this eutectoid reaction could occur quickly.
Therefore, we tend to think that the minor phase present in the
room temperature XRD, Fig. 1(a), was Fe Se rather than d.
7 8
Se , at 3401–3501C
5
7
8
Also observed in DSC was the peritectoid reaction:
aÀFe1d2b, at about 4601C, very close to the 4571C given in
6
the Fe–Se phase diagram. For the samples whose XRD did not
7 8
show the trace of Fe Se , only one peak corresponding to this
peritectoid reaction was observed, as shown in Fig. 2(b).
(2) Liquid-Phase Processing
Figure 3(a) is a low-magnification SEM image, showing a typ-
ical overall surface morphology of the thick films grown on
(
001) LAO. An area rich of the features are highlighted in
Fig. 3(b), in which the places where the EDX analyses have
been performed are also indicated. The EDX results, Fig. 3(c),
showed that the thick films had the Se:Fe atomic ratio approx-
imately equal to 1 within the experimental error. Both EBSD
structural analyses and the superconducting measurements,
which will be discussed later, verified the films being b-phase.
Figure 4 shows the Al
by the liquid during the growth over the temperature range of
101–8001C. The grains grown from the liquid on the rough
Al surface were large with a diameter 100–200 mm, com-
2 3
O boat surface, which was also wetted
9
2 3
O
pared with about 100 nm for the sintered powders. EDX
showed these large grains having the same composition as the
thick films on LAO. XRD of the samples, Fig. 1(c), matched
well with the tetragonal P4/nmm structure (PDF #85-0735),
2 3
except for those lines from the Al O (PDF #78-2427), con-
firming that they were pure b-phase.
The above results are interesting because the thick films on
LAO and also the large grains on rough alumina were grown
from the Sn-containing liquid at temperatures between 9101 and
8001C, well above the peritectoid temperature of 4571C, but what
was obtained was the low-temperature b-phase, instead of the
high-temperature d-phase, indicating that Sn had some kind of
effect on stabilizing the b-phase. However, EDX failed to find
any Sn trace in the thick films or the large grains, which was in-
13
consistent with the FeSe–SeSn phase diagram, where a solid
solution region was shown with the maximum Sn solubility being
about 5 at.%, for the Fe1 SeSn solid solution precipitated from
Ày
y
the liquid at the peritectoid temperature (7801C). The reason is
not clear. One possibility is the significant evaporation of Sn at
the growth temperatures, which decreases the actual solubility of
Sn to a level beyond the sensibility of the EDX. As shown in Figs.
3(b) and 4(b), some metallic iron was observed in both the thick
films and large grains. The water-drop-like shape is reminiscent
that they were solidified from the liquid. But metallic iron would
not melt at 9101C. The iron water-drops propose that they were
solidified via the evaporation of Sn and/or Se.
EBSD was carried out to check the structure and orientation
of the samples. Figure 5(a) shows the EBSD pattern from a
point on the alumina sample surface, Fig. 5(b). The surface im-
age was slightly distorted because the sample was tilted at 701
for recording the diffraction patterns, which could be indexed
only according to the tetragonal P4/nmm structure, confirming
again that they were indeed the superconducting b-phase. EBSD
mapping over an area, Fig. 5(c), showed that the sample mainly
consisted of (211) and (101) grains. This was consistent with the
XRD, Fig. 1(c), where the relative intensities of these two re-
flections were enhanced as compared with the powder samples.
To our surprise, although the thick films on (001) LAO were
pure and smooth, they did not have a single orientation, as
shown in the EBSD mapping in Fig. 6. The two observed ori-
entations, (101) and (201), looked like a twin. The thickness of
the film was measured using a profilometer to be about 1.5–2
mm, which was probably too thick to retain a good texture.
Fig. 3. (a) Scanning electron microscope image of the overall surface
morphology for the thick films grown on LaAlO (LAO). (b) A partic-
3
ular area rich of the features, also showing the places where energy-dis-
persive X-ray spectroscopy (EDX) compositional analyses were carried
out. (c) A representative EDX result for the red square marked in (b),
showing that the thick film had the Fe:Se ratio approximately equal to 1.
(
3) Superconducting Properties
Shown in Fig. 7 are the magnetization versus temperature (M–
T) curves measured by SQUID magnetometry. The samples
were cooled in zero-field from room temperature to 2 K and the
moment was then measured in the applied field of 10 Oe upon

3
198
Journal of the American Ceramic Society—Qi et al.
Vol. 93, No. 10
Fig. 4. (a) Scanning electron microscope image of the overall surface of
the alumina (Al ) boat wet by the high-temperature liquid. (b) Details
of the surface and the areas where energy-dispersive X-ray spectroscopy
EDX) analyses were performed.
2 3
O
(
heating. Rapid increase of moment was observed until about 10
K. The onset T ’s were taken at the intersection point of the two
c
extrapolated lines, which were 8.1 and 6.0 K for the sintered
powders and thick films, respectively. Four-point resistivity
measurements confirmed the transitions and showed an onset
T of 9.7 K, with zero resistance at 6.9 K, for the sintered b-
c
FeSe1Àx powder, as shown in Fig. 9(a).
Magnetization versus applied-field (M–H) measurements at 4
K showed typical hysteresis loops seen in the ferromagnets for
all the samples in Fig. 7, presumably due to one or more of the
7 8 3 4
magnetic traces of Fe Se , a-Fe, and Fe O observed in XRD
and SEM, as discussed above. A representative of such M–H
loops is shown in Fig. 8. The influence of the magnetic impurity
traces was also reflected in the M–T curves, Figs. 7(a)–(c), as the
measured moments were unable to go into negative at T . Com-
c
paratively, the thick film crystallized from the high-temperature
liquid appeared to have the best purity, as its moment was the
c
most close to zero at T . Although the measured moments were
different in Figs. 7(a)–(c) owing to the different types and/or
amounts of the residual traces, the drop of the moments from
À3
the onset T to 2 K, i.e. DꢁM(10K)ÀM(2K)ꢁ9 Â 10 emu/g,
c
was approximately the same, as it should be. For an easier com-
parison of the T
D, are plotted in Fig. 7(d). The decrease of T for the thick film is
apparent. On the other hand, its transition is steeper due to a
higher purity and structural perfection. It is not very clear what
c
’s, the normalized moments, [M(T)ÀM(10K)]/
Fig. 5. (a) A representative EBSD pattern from a point on the alumina
sample surface, which is shown in (b). The pattern matched well with the
structure of b phase. (c) Electron backscatter diffraction mapping on (b),
showing different orientations presented (the sample was tilted at 701 for re-
cording the diffraction patterns, and hence the images had some distortion).
c
c
exactly caused the drop of T . One possible cause was the small

October 2010
High-purity FeSe1 Superconductors
Àx
3199
Fig. 7. Magnetization versus temperature (M–T) curves of the samples:
(
a) sintered powder of b-FeSe1Àx, (b) thick film grown from a high-tem-
perature solution, (c) sintered powder of Sn-doped b-FeSe1Àx with the
nominal doping concentration of 5 at.%, (d) comparison of the normal-
c
ized M–T curves, showing a lower T for the thick film. The curves were
measured in the field of 10 Oe on heating up after zero-field cooling.
maining small resistance was likely to come from the segregation
of some (Se, Sn)-enriched impurity phases on the surface, which
were formed during the heating because of the high vapor pres-
sure of Se and Sn. This speculation was based on the XRD re-
sults of a number of samples sintered at high temperature for a
prolonged period of time. Unknown phases were sometimes
observed on the surface of such sintered pellets and a layer of a
few hundred micrometers thick had to be ground off to expose
the superconducting phase inside. The residuals of surface
Fig. 6. (a) Scanning electron microscope of the surface of a thick film
grown on LaAlO . (b) Electron backscatter diffraction mapping on a
3
part of (a), showing that the film did not have a single orientation.
amount of Sn doping, because the thick films were grown from
the Sn-containing high-temperature liquid. However, EDX
failed to detect any Sn trace in the films, indicating that the
Sn contamination, if there was any, must be very low, beyond
the typical EDX resolution of about 1%.
In order to see the effect of Sn doping on the FeSe1 super-
Àx
conductivity, b-FeSe:5% Sn powder samples were prepared by
sintering at various temperatures from 4101 to 7501C. All the
Sn-doped samples showed little change of the T from the pure
FeSe1 in both M–T and R–T measurements, as shown in Figs.
7
the transition for the doped sample did not go to zero. The re-
c
Fig. 8. A representative magnetization versus applied field loop of the
b-FeSe1 samples measured at 4 K. The observed moment and reman-
Àx
and 9. It is noted in Fig. 9(b) that the resistivity at the end of
Àx
ence were likely to come from the trace of the magnetic secondary phases
3 4
(e.g., Fe O , etc.) as described in the text.

3
200
Journal of the American Ceramic Society—Qi et al.
Vol. 93, No. 10
had almost a pure b-phase. However, the samples sintered at
higher temperatures contained some amount of d-phase, which
was unable to be eliminated completely by prolonged annealing
at r4501C afterward. This was due to the phase segregation
above the peritectoid reaction temperature of 4571C. The mea-
surements showed that sintered pure b-FeSe1 samples were
Àx
superconducting with the onset T ’s of 8.1 and 9.7 K determined
c
from the M–T and R–T curves, respectively. Therefore, the op-
timum conditions for sintering high-purity b-FeSe1Àx powder
samples are: temperature 4001À4501C and composition
Fe:Se 5 1:0.95–0.97.
Thick films of FeSe1Àx were grown on (001) LAO single
crystal substrates by liquid-phase processing with SeSn as the
flux. In contrast to the sintering, the films grown around 9101C
from the Sn-containing liquid had pure b-phase with a high
crystallinity, as shown by both EBSD and XRD. Although the
phase diagram showed some solid solubility of Sn in FeSe, EDX
failed to find any Sn trace in the crystallized films, which were
also superconducting with a steeper transition, but a lower T
.0 K determined by the magnetic measurements. In order to
clarify whether the possible Sn contamination caused the drop
of T , Sn-doped FeSe powder samples were synthesized,
c
of
6
c
1Àx
which showed little change of T
c
from the pure b-FeSe1Àx
.
c
The cause for the slight reduction of T of the thick films
remains unclear.
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impurity phases might cause some voltage drop across the con-
tacts, causing nonzero resistance.
(
12
S. B. Zhang, H. C. Lei, X. D. Zhu, G. Li, B. S. Wang, L. J. Li, X. B. Zhu, W.
H. Song, Z. R. Yang, and Y. P. Sun, ‘‘Divergency of SDW and Structure Tran-
sition in Fe1ÀxNi Se0.82 Superconductors,’’ Phys. C, 469 [21] 1958–61 (2009).
P. Villars, A. Prince, and H. Okamoto, Handbook of Ternary Alloy Phase Di-
agrams, Vol. 8. ASM International, Metals Park, 1995, p. 10847.
x
13
IV. Concluding Remarks
14
Bulk FeSe1Àx samples were prepared by a solid-state reaction in
the Ar13% H gas flow. The samples sintered at 4101C for 24 h
2
PDF #75-1843: Powder Diffraction Filet (PDF) #75-1843, the International
Centre for Diffraction Data (ICDD).
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