Effects of MgO impurities and micro-cracks on the critical current density of Ti-sheathed MgB2 wires
Department of Mechanical Engineering, University of Houston, Houston, Texas, United StatesPhysica C Superconductivity (Impact Factor: 0.94). 06/2007; 457(1-2):47-54. DOI: 10.1016/j.physc.2007.02.013
Ti-sheathed monocore MgB2 wires with improved magnetic critical current density (Jc) have been fabricated by in situ powder-in-tube (PIT) method and characterized by magnetization, X-ray diffraction, scanning electron microscopy and electrical resistivity measurements. For the best wire, the magnetic Jc values at 5 K and fields of 2 T, 5 T, and 8 T are 4.1 × 105 A/cm2, 7.8 × 104 A/cm2, and 1.4 × 104 A/cm2, respectively. At 20 K and fields of 0.5 T and 3 T, the Jc values are about 3.6 × 105 A/cm2 and 3.1 × 104 A/cm2, respectively, which are much higher than those of the Fe-sheathed mono-core MgB2 wires fabricated with the same in situ PIT process and under the same fabricating conditions. It appears that the overall Jc for the average Ti-sheathed wires is comparable to that of the Fe-sheathed wires. Our X-ray diffraction and scanning electron microscopy analysis indicates that Jc in the Ti-sheathed MgB2 wires can be strongly suppressed by MgO impurities and micro-cracks.
Eﬀects of MgO impurities and micro-cracks on the critical
current density of Ti-sheathed MgB
, M. Alessandrini
, F. Yen
, M. Hanna
, H. Fang
, C. Hoyt
, J. Zeng
, K. Salama
Department of Physics, Sam Houston State University, 1908 Ave. J, Huntsville, TX 77341, USA
Department of Mechanical Engineering, University of Houston, 4800 Calhoun Road, Houston, TX 77204, USA
Texas Center for Superconductivity, University of Houston, Houston, TX 77204, USA
Department of Chemistry, University of Houston, Houston, TX 77204, USA
Received 27 November 2006; received in revised form 5 February 2007; accepted 21 February 2007
Available online 3 March 2007
Ti-sheathed monocore MgB
wires with improved magnetic critical current density (J
) have been fabricated by in situ powder-in-tube
(PIT) method and characterized by magnetization, X-ray diﬀraction, scanning electron microscopy and electrical resistivity measure-
ments. For the best wire, the magnetic J
values at 5 K and ﬁelds of 2 T, 5 T, and 8 T are 4.1 · 10
, 7.8 · 10
1.4 · 10
, respectively. At 20 K and ﬁelds of 0.5 T and 3 T, the J
values are about 3.6 · 10
and 3.1 · 10
tively, which are much higher than those of the Fe-sheathed mono-core MgB
wires fabricated with the same in situ PIT process and
under the same fabricating conditions. It appears that the overall J
for the average Ti-sheathed wires is comparable to that of the
Fe-sheathed wires. Our X-ray diﬀraction and scanning electron microscopy analysis indicates that J
in the Ti-sheathed MgB
can be strongly suppressed by MgO impurities and micro-cracks.
2007 Elsevier B.V. All rights reserved.
PACS: 74.70.Ad; 84.71.Mn; 74.25.Sv; 75.60.Ej; 61.10.Nz
superconductor; Magnetic J
; Hysteresis loop; X-ray diﬀraction; SEM image
For certain type of applications, such as in the growing
sector of electric space propulsion, lightweight supercon-
ducting magnets are preferred. Lightweight magnets
require the use of light metals as the sheath materials in
the powder-in-tube (PIT) process to make MgB
[1,2]. The most favorable sheath metal should at least meet
the following requirements: low mass density, excellent
chemical compatibility (i.e., non-reactive or nearly non-
reactive with MgB
, B, or Mg during sintering process),
non-magnetic, and appropriate mechanical strength for
cold work. In the last ﬁve years, PIT MgB
have been fabricated with the use of diﬀerent sheath metals,
such as iron (Fe), nickel (Ni), copper (Cu), silver (Ag), nio-
bium (Nb), and tantalum (Ta) [3–10]. For un-doped MgB
wires sheathed with some of these metals such as Fe, J
ues as high as 3 · 10
have been achieved at 5 K and
in zero ﬁeld [4,9–11]. However, none of these sheath metals
are suitable for making lightweight MgB
their mass densities are all greater than 7.8 g/cm
(Ti) seems to be a very promising sheath material for light-
weight and high J
wires due to its very low mass
density (4.57 g/cm
), excellent chemical stability, non-mag-
netic property [12,13], and high mechanical strength (50%
harder than Fe). Recently, we have successfully fabricated
0921-4534/$ - see front matter 2007 Elsevier B.V. All rights reserved.
Corresponding author. Tel.: +1 936 294 1608; fax: +1 936 294 1585.
E-mail address: firstname.lastname@example.org (G. Liang).
Physica C 457 (2007) 47–54
wires with J
values as high as 4 · 10
at 5 K and zero ﬁeld [14,15]. Currently, much eﬀo rt has
been made by our groups to improve the J
. In this paper,
we present the best results on the magnetic J
for the newly
developed un-doped Ti-sheathed MgB
with the results from electrical resistivity, X-ray diﬀraction
(XRD), and scanning electron microscopy (SEM) measure-
ments. This study also provides the ﬁrst analysis on the
micro-cracks and MgO impurities in Ti-sheathed MgB
wires to address the observed degradation of J
of these wires.
The Ti-sheathed monocore MgB
wires were fabricated
by an in situ powder-in-tube (PIT) method, which has been
described in detail elsewhere [14,15]. In this paper, we only
report results on un-doped MgB
wires sintered at 800 C
for 30 min. The milled Mg + 2B powder particles are very
uniform in size and smaller than 2 lm, as checked by SEM
. The cross-sectional areas for the wires and cores are
about 1 mm
(=1 mm · 1 mm) and 0.126 mm
tively. The XRD patterns were obtained using a Rigaku
X-ray diﬀractometer with Cu K
radiation. All of the pat-
terns were calibrated using Si powder as the standard. The
ﬁtting of the XRD peaks was done using software Jade 6.1
provided by Material Data Inc. (MDI). The SEM images
were taken using a JEOL JSM-6330F Field Emission Scan-
ning Electron Microscope. The temperature (T) dependent
resistivity was measured by a standard four-probe dc tech-
nique in a temperature range of 10–100 K. The hysteresis
loops of the magnetization (M) were measured using an
Oxford Instrument ’s Maglab 9-Tesla Vibrating Sample
Magnetometer (VSM). The temperature dependent magne-
tization was measured in both zero-ﬁeld-cooled (ZFC) and
ﬁeld cooled (FC) modes using a magnetic properties mea-
surement system (MPMS) magnetometer from quantum
3. Results a nd discussion
Fig. 1 displays the magnetization hy steresis loops,
M(H), measured at 5 K, 20 K, and 30 K for two wire sam-
ples: sample A and B. The diﬀerence between the wire sam-
ple A and B is: wire A was made with relatively slower
rolling speed and is micro-crack free; whereas wire B was
made with faster rolling, resulting in producing micro-
cracks (see the SEM result below) in the wire. In the
M(H) hysteresis loop measurements, each piece of wire
sample was prepared with almost identical dimension to
avoid the sample-size eﬀect on J
. The sample A con-
sisted of six pieces of short wires, each has a dimension
of a · b · c = 0.033 cm · 0.038 cm · 0.73 cm for the MgB
core. The sample B had a dimension of a · b · c =
0.034 cm · 0.037 cm · 0.75 cm for its MgB
core. The wire
samples were oriented with the c-axis parallel to the direc-
tion of the applied magnetic ﬁeld. In Fig. 1, the solid curves
are the M(H ) data measured with the VSM at a very small
step of 33 Oe. For sample A, it is observed from Fig. 1a
that the M(H) loop measured at 5 K and low ﬁelds
(between 1.7 T and 1.7 T) displays a plateau like shape
with some spikes. This behavior can be attributed to the
ﬂux jumps caused by local ized warming due to very high
critical current density and very low speciﬁc heat .
For a superconductor, when the ratio of J
to speciﬁc heat
(C) is greater than a critical value, ﬂux jumping occurs .
For M gB
, since J
increases and C decreases  with the
decrease of the applied ﬁeld, ﬂux jumping could occur at
low ﬁelds. Such ﬂux jumps can cause the ﬂuctuation of
the magnetization and magnetic critical current density,
as observed previously by some groups [16,19,20] for the
wires. To verify this ﬂux jumping, we
measured the M(H) loop at 5 K and up to ±20 kOe (or
H = ±2 T) using the MPMS magnetometer. The data
points are shown in Fig. 1a by the open circles. Indeed,
Fig. 1. The solid curves are the M(H) curves measured for samples A and
B at diﬀerent temperatures using VSM. The open circles in (a) are the data
points measured at 5 K using MPMS.
48 G. Liang et al. / Physica C 457 (2007) 47–54
we see that the data points are ﬂuctuating in the plateau
region and match well with the solid M(H) curve in the
regions beyond the plateau. In contrast with the loop mea-
sured at 5 K, the M(H) hysteresis half loops measured at
20 K and 30 K (see Fig. 1b) do not show any ﬂux-jump-
caused ﬂuctuation due to both the substantial decrease in
and increase in speciﬁc heat at these temperatures.
Fig. 1c shows that for sample B, the ﬂux-jump-caused mag-
netization ﬂuctuation in the 5 K M(H) curve below 1 T is
much smaller than the ﬂuctuation for the sample A, sug-
gesting a much lower J
for sample B.
Shown in Fig. 2 are the ﬁeld dependent magnetic J
curves for these two samples. The magnetic J
lated by J
=20DM/[a(1 a/3b)] [15,16,21] from the Bean
critical state model , where DM (in emu/cm
) is the dif-
ference between the upper and lower branches of the M(H)
curve. There is no contribution to DM from the paramag-
netic Ti-sheath [12,13] because its magnetization is revers-
ible with applied ﬁeld . Fig. 2 shows that for sample
A, the magnetic J
at 5 K and ﬁelds of 2 T, 5 T, and 8 T
are about 4.1 · 10
, 7.8 · 10
, and 1.4 ·
, respectively. At 20 K and ﬁelds of 0.5 T and
3 T, the J
values are about 3.6 · 10
3.1 · 10
.InTable 1 , we summarize the values of
the magnetic J
at diﬀerent temperatures and ﬁelds for sam-
ple A and sample B. Fig. 2 shows that at 5 K and 20 K, the
for sample A at any ﬁeld is much higher than that for
sample B. At 30 K, J
of sample A is higher than that of
sample B for l
H < 0.94 T and lower than that of sample
B for l
H > 0.94 T.
For sample A, since the J
at 5 K and in the low ﬁeld
regions (for l
H 6 1.7 T) is extremely high, the Bean model
does not apply due to signiﬁcant ﬂux jumps. For sample B,
Fig. 2 shows that its J
at 5 K is only about half of J
sample A, the ﬂux jump caused ﬂuctuation in the J
is very small, below 1 T. Thus, if we use the J
values of the
sample B below 1.7 T and assume that the variation rates
below 1.7 T are the same for both sample A and B,
then the J
at 5 K and 0 T for sample A is estimated to
be about 1.0 · 10
. The dashed curve section in
Fig. 2 represents the estimated values of J
at 5 K and
below 1.7 T for sample A.
To evaluate the performance of the Ti-sheath on J
would like to compare the magnetic J
between the Ti-
sheathed and Fe-sheathed MgB
wires. Since magnetic J
could be aﬀected by certain factors such as the fabricating
condition, doping content (such as SiC ), and the qual-
ity of the precursor powder, the most appropriate
Fe-sheathed wires for such comparison should be those
fabricated the same way as that used for the Ti-sheathed
wires. Unfortunately, no other groups, except ours, have
fabricated Fe-sheathed, un-doped MgB
wires with the
same in situ PIT process, fabricating equipment (such as
the groove rolling mill), and fabri cating conditions in terms
of Mg and B powders, milling time (2 h), sintering temper-
ature (800 C), and size of the wires (1 mm · 1 mm cross
section)  . For these Fe-sheathed wires, the magnetic J
was measured at 20 K and in a ﬁeld range from 0 to 3 T
, and the data are presented in Fig. 2 by the open circles
and listed in Table 1. It can be seen from Fig. 2 that for
ﬁelds in the range of l
H < 2.5 T, the J
values for the
Fe-sheathed wires are between the values for samples A
and B; for l
H > 2.5 T, the J
values for the Fe-sheathed
wire are lower than values of both sample A and B. This
comparison indicates that even though the best Ti-she athed
Fig. 2. The ﬁeld dependent magnetic J
curves measured at temperatures
5 K, 20 K, and 30 K for wire sample A (thicker curves) and sample B
(thinner curves). The dashed curve section below 1.7 T for the 5 K J
of sample A represents the estimated values (see text). The open circles are
the data points of magnetic J
measured at 20 K for a Fe-sheathed MgB
wires (Ref. ).
The values of the magnetic J
for sample A and B at two temperatures, 5 K and 20 K, and at diﬀerent applied magnetic ﬁelds
0 (T) 2 (T) 5 (T) 8 (T) 0.5 (T) 2 (T) 3 (T) 4 (T)
wire sample A 10.0
4.1 0.78 0.14 3.58 0.87 0.31 0.09
wire sample B 4.2 1.7 0.36 0.10 1.41 0.34 0.14 0.05
2.50 0.49 0.09
For comparison, the J
values for the Fe-sheathed MgB
wires, fabricated previously by us (Ref. ) with the same fabricating conditions, are also listed.
Estimated (see text).
From Ref. .
G. Liang et al. / Physica C 457 (2007) 47–54 49
wire (wire A) has a much higher J
than the Fe-sheathed
wires, the overall performance of the Ti-sheath on J
the average Ti-sheathed wires (including wire B) is compa-
rable to the performance of the Fe-sheath.
To explain the observed big diﬀerence in J
sample A and sample B, detail analysis on the crystal
phases and micro-structure of the samples is needed.
Fig. 3 shows the powder XRD patterns of the sample A
and B, together with the patterns of three reference com-
, MgO, and Ti (all 325 mesh, from Alfa
Aesar). All of the major peaks in the patterns of sample
A and B can be indexed with the MgB
hexagonal struc -
ture, indicating that the core material of the wires is in
nearly pure MgB
phase. For the pattern of sample A, a
weak impurity peak due to the MgB
phase is observed
located at 2h = 35.51, which is not seen from the pattern
for sample B. The broad peak located at 2h = 62.33 can
be attributed to the second strongest peak of the cubic
MgO phase, i.e., the (220) peak (also see Fig. 4). The stron-
gest peak of the MgO phase is located at 2h 42.9 and
overlaps with the neighboring MgB
(101) peak located
at 2h 42.4, causing some additional broadening eﬀect
on this peak. Such line broadening was conﬁrmed by our
detail analysis on the full width at half-maximum (FWHM)
of the peaks. In the presence of oxygen, MgO impurities
could be formed by the reaction 2Mg + O
! 2MgO dur-
ing the sintering process. The main source of the oxygen
could be from the air trapped in the cores during the
short-time crimping sealing of the end of the Ti tube (with
Mg + 2B mixture packed in) in air. The residual O
tained in commercial argon gas could be the secondary
source. Compared with the pattern of the Ti powder, the
two weak impurity peaks locat ed at 2h = 38.40 and
40.12 in the patterns of sample A and B can be identiﬁed
as the (0 0 2) and (1 0 1) peaks of the a-phase Ti. The inten-
sities of these two peaks are less than 2% of the intensity of
(101) peak in each pattern. We believe that the
Ti impurities were introduced into the XRD samples dur-
ing the preparation of the XRD powder slide and thus they
are extrinsic. When the wires were peeled open by a knife,
Fig. 3. Powder XRD patterns for the core materials of the Ti-sheathed
wire samples A and B. For comparison, the XRD patterns for the
reference compounds MgB
, MgO, and Ti are also shown.
Fig. 4. XRD patterns in the 2h range of 59 6 2h 6 65 for sample A,
sample B, and a mixture of MgB
and MgO with 1:1 molar ratio. The
patterns are normalized to the intensity of the MgB
(110) peak in each
pattern. Note that the scale for the intensity of the pattern of the mixture
and MgO is reduced by a multiplying factor of 0.25.
Fig. 5. EDS spectrum taken for the entire cross-sectional area of the
core of sample B.
50 G. Liang et al. / Physica C 457 (2007) 47–54
some very small Ti particles were stripped away from the
Ti-sheath and fell into the core material. Fig. 5 shows an
energy dispersive spectrometry (EDS) spectrum for the
core of the wire sample B, and it indeed conﬁrms the
absence of the Ti impurities in the cores of the as-sintered
It can be seen from the expanded patterns in Fig. 4 that
the intensity of the MgO (2 2 0) peak for sample B is stron-
ger than that for sample A, indicating a relatively larger
content of MgO in sample B. To better understand how
the concentration of the MgO impurities correlates to the
variation of J
in the MgB
wires, it is very necessary to
have a good estimate about the MgO concentration in
these two samples. For this purpose, we performed a quan-
titative analysis on the MgO (220) and the MgB
peaks by a similar technique which we recently used to ana-
lyze the concentra tion of MgCu
impurities in Cu-sheathed
wires . Fig. 4 shows the XRD patterns of the two
samples in comparison with a pattern measured on a mix-
ture of MgB
and MgO, which has a molar (mol) ratio of
1:1. The pattern for the mixture is used to obtain the cali-
bration line . This pattern for the mixture shows that
the intensity of the MgO (22 0) peak is much stronger
(about ten times stronger) than that of the MgB
peak, even though the molar content of these two compo-
and MgO, are equal. Due to the particle size
eﬀect , the peaks for the sample A and B (with much
ﬁner size) are much wider than that for the powders of
the mixture (particle size 325 mesh). Note that the pattern
for the Ti powder (in Fig. 3) is not included in Fig. 4,
because for the patterns of sample A and B, the intensity
of the Ti (1 1 0) peak located at 2h 63 is estimated (using
the intensities of the relevant peaks of the patterns in
Fig. 3) to be only about 3% of the intensity of the MgB
(102) pe ak. The percentage weight of MgO in the mixture
of MgO and MgB
can be expressed as W
where the relative intensity I
is deﬁned as the ratio of
the intensity of the MgO (2 2 0) peak to that of the MgB
(102) peak, i.e., I
); K is a con-
stant which depends only on the Miller indices of the two
selected peaks and can be determined experimentally by
the slope of the W
calibration line . The rela-
tionship between the molar (mol)%, M
, and weight
, of MgO in a mixture of MgO and MgB
is given by formula M
/(r + W
where r = 0.8776 is the molar mass ratio of MgO to MgB
To determine the intensities and peak positions (2h)
(110), MgO (2 20), and MgB
peaks in the patterns of Fig. 4, these peaks wer e ﬁtted with
the pseudo-Voigt proﬁle shape functions using the Jade 6.1
XRD analyzing software package. The ﬁtting results are
listed in Table 2 and the values of the I
are calculated from the ﬁtting result and summa-
rized in Table 3. Table 3 shows that the molar MgO con-
centration for sample B is about 4.8%, which is three
times of the MgO concentration in sample A.
Fig. 6 shows the SEM images for these two samples. The
images clearly show that there exist a large number of
cracks in sample B and no cracks in sample A. A close
inspection of the outer surface of the Ti-sheath of the sam-
ples also revealed that some cracks were developed in cer-
tain portion of the Ti-sheath of wire sample B but not in
the wire sample A. These micro-cracks could be produced
due to the relat ively faster feeding of the wire sample B
during the wire rolling process. The SEM result supports
the following explanation for the big diﬀerence in the
MgO concentration between sample B and sample A: with
the existence of substantial micro-cracks in both the Ti-
sheath and the core region of sample B, the oxygen in
the air could percolate or diﬀuse along the cracks into
the core region from outside of the wire and then react with
Mg to form extra MgO during the sintering process. Such
extra MgO could not be produced in the crack-free sample
A. Correlating to the J
results shown in Fig. 2, it appears
that the decrease of J
from sample A to sample B could be
due to the substantial increase of MgO content from 1.6%
in sample A to 4.8% in sample B. Similar correlation
and MgO content was also seen from the recent
results of Chen et al. [25,26] that J
is decreased substan-
tially (by a factor of 3 or more) with slight increase of
the MgO content in their MgB
samples. It was shown
The peak ﬁtting result for the MgB
(110), MgO (2 2 0), and MgB
(10 2) peaks in the XRD patterns shown in Fig. 4
(11 0) peak MgO (2 2 0) peak MgB
(10 2) peak
FWHM I 2h FWHM I 2h FWHM I
MgO + MgB
59.86 0.35 1 62.26 0.27 3.7 63.16 0.40 0.33
Sample A 59.88 0.64 1 62.19 0.94 0.13 62.91 1.08 0.38
Sample B 59.88 0.66 1 62.07 0.94 0.61 63.10 1.17 0.61
Fitting parameters: (2h)
is the centroid position of the peak, FWHM is the full width at half maximum, and I is the peak intensity measured by the area
under the ﬁtting curve for each peak. The values of intensity are normalized to the intensity of the MgB
(11 0) peak for each XRD pattern.
The values of the relative intensity I
, and M
, which are
deﬁned respectively as I
), MgO wt.%, and MgO
mol%, for the following three samples: the mixture of MgO and MgB
powders with 1:1 molar ratio, sample A, and sample B
MgO + MgB
11.2 46.7 50
Sample A 0.34 1.4 1.6
Sample B 1.00 4.2 4.8
G. Liang et al. / Physica C 457 (2007) 47–54 51
previously that the MgO grains can be formed along the
boundaries of or even inside the MgB
grains [27,28]. Such
MgO impurities at the grain boundaries could make the
inter-grain connectivity worsen and thus cause sub-
stantial degradation in J
In addition to MgO, MgB
impurities could also be
formed during the sintering process by either the reaction
interesting question related to the production of MgO
is whether the formation of MgB
the concentration of the MgO and vice versa. Previously,
Li et al.  observed that MgB
impurities were formed
in their MgB
samples sintered at 800 C but not in samples
sintered at 660 C. This observation indicates that MgB
may be formed only at high temperatures, at least above
660 C. On the other hand, since the autoignition tempera-
ture of Mg is 473 C, Mg can react easily with O
MgO at temperatures between 473 C and 660 C. Thus,
we propose the following picture: In the sintering process
of our Ti-sheathed wires, it is very possible that prior to
the formation of MgB
, some starting Mg particles had
already reacted with all (or most) of the O
the cores of the wires to form MgO. After the temperature
reached 800 C, small amount of MgB
could be produced
under certain conditions by the decomposition of MgB
At present, the detail conditions for the formation of
in the wires are not clear and further study on this
issue is needed. Since all of the O
trapped in the core of
the wire had been consumed earlier in forming MgO, no
oxygen would be available for the newly produced Mg
+ Mg) to be oxidized to form new
MgO, instead, the Mg could either remain in the core
(may not be detected by XRD due to low concentration)
or escape from the open ends of the wires. In such pro-
posed picture, the concentration of the MgO in the samples
should be predominantly (or solely) determined by the ini-
tial content of the oxygen trapped in the core before the
sintering process. This means that the MgO concentration
should be either independent of or ne arly unaﬀected by the
formation or concentra tion of MgB
impurities. In another
word, less production of MgO does not necessarily mean
less production of MgB
and vice versa, as observed in
the XRD patterns (see Fig. 3) of the samples. This explana-
tion is also consistent with the observation that for many
-based materials, MgO impurities usually present
without the accompanying formation of MgB
The SEM images in Fig. 6 show that there exists a large
amount of spherical holes or voids in the cores of sample A
and B. Most of the holes/voids are about 1–2 lm in diam-
eter, which is close to the size of the Mg powder particles in
the milled Mg + 2B powder precursor. These voids could
be produced by the volume reduction in the Mg + 2B !
reaction, it also could be partially attributed to the
evaporation of the Mg particles during the sintering of
the wires. It can be seen from Fig. 6 that the density of
the voids for the sample B is higher than that for sample
A. This is understandable becau se the micro-cracks (which
do not exist in sample A) in the core and certain portion of
the Ti-sheath of sample B provided more ‘‘doors’’ for the
evaporated Mg particles to escape during the sintering,
leaving behind, in the core, more voids/holes. Previously,
the existence of voids/holes was also observed in the cores
of the Fe-sheathed MgB
wires [6,32]. Thus, besides more
content of MgO impurities, the existence of micro-cracks
and more voids in sample B in contrast with those in sam-
ple A could be another contributing factor for the depres-
sion of J
due to the worsening connectivity between larger
Fig. 7 shows the 4pv(T) curves for the wire samples,
where v (= M/H) is the dc magnetic susceptibility. The crit-
ical transition temperature, T
, deﬁned as the onset of the
diamagnetism, is about 36 K for both the two samples.
The width of the transition (10–90% of the full drop in v)
is about 1.5 K. This T
value is comparable to the T
(35–37 K) reported for some Fe-, Nb-, and Cu-sheathed
wires [7,8,33,34]. Shown in the inset of Fig. 7 are
the electrical resistivity q( T) curves for the samples, which
give the same onset T
and transition width as the values
determined from the 4pv( T) curves. Above T
, the q(T)
curve displays a meta llic behavior similar to the q(T) curves
for the MgB
pellet samples . Fig. 7 shows that the val-
Fig. 6. SEM images of the cores of the Ti-sheathed MgB
wires: (a) for
sample A, and (b) for sample B. The surfaces of the cores were polished
before taking the images. These SEM images show that some micro-cracks
exist in sample B but not in sample A.
52 G. Liang et al. / Physica C 457 (2007) 47–54
ues of 4pv at 5 K in the ZFC condition are about 0.75
and 0.63 for sample A and B, respectively, corresponding
to 25% non-superconducting volume fraction in the core
of sample A and 37% in sample B. Since the non-supercon-
ducting volume fraction represents the sum of the volume
of all open holes, voids, cracks, and non-superconducting
materials in the cores, this result is consistent with the
SEM result that more voids (including cracks) were formed
in sample B when compared those formed in sample A. For
both sample A and B, the 4pv values in the FC mode is
very small (only ab out 3%). The large diﬀerence between
the 4pv vales in the ZFC mode and FC mode indicates
that there exists a strong ﬂux pining force in both sample
A and B, which holds most of the magnetic ﬂux in the cores
under the FC condition.
In summary, we have successfully fabricated Ti-
sheathed, undoped monocore MgB
wires by in situ PIT
method with J
improved substantially from pr eviously
fabricated wires. Particularly, for our best wire (sample
A), the J
is much higher than that of the Fe-sheathed
wires fabricated with the same in situ
PIT process and under the same fabricating conditions.
However, our J
results also suggest that the overall J
for the average Ti-sheathed wires including wire sample B
is comparable to the J
of the Fe-sheathed wires. Our
XRD and SEM an alysis reveals that MgO impurities and
micro-cracks could be the possible factors for the degrada-
tion of critical c urrent density in these Ti-sheathed MgB
wires. The results from the XRD, magnetic J
EDS, M(T), and q(T) measurements are mutually consis-
tent and well correlated.
The authors thank Dr. Z. Tang, Dr. J. Meen, A. Scotti,
M. Crush, and Dr. J. Horvat for either their assistance in
measurements or helpful discussion. This work was sup-
ported by Sam Houston State Un iversity’s Faculty Re-
search Grant and 2006 EGR grant, an award from
Research Corporation, and by the State of Texas through
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