Japanese Journal of Applied Physics
Vol. 47, No. 2, 2008, pp. 1412–1415#2008 The Japan Society of Applied Physics

Synthesis of Polymerized C60 Bulk and Film in Three Dimensions
by Irradiating Free-Electron Laser
Nobuyuki I WATA, Shingo A NDO, Ryo NOKARIYA, and Hiroshi YAMAMOTO
Department of Electronics and Computer Science, College of Science and Technology, Nihon University,
7-24-1 Narashinodai, Funabashi, Chiba 274-8501, Japan(Received April 16, 2007; revised July 14, 2007; accepted July 31, 2007; published online February 15, 2008)

Pressed C60 powder was polymerized with free-electron laser (FEL) irradiation. Brightened grains were synthesized in part at
the sample surface, the Raman spectrum of which indicated the polymerization. It was noted that polymerization occurred as aresult of irradiating a third harmonic FEL at 500 nm with an extremely low energy density of 2.6 mJ/(pulsecm2 ). The
polymerization is considered to be due to the micropulse width less than one ps forming the 20 ms macropulse of the FEL.
Irradiating FEL during ﬁlm growth, at 200  C in particular, promoted the polymerization, as demonstrated by the Raman
spectrum. The X-ray diﬀraction showed an oriented hcp structure with the relationship that the [0001] direction of the ﬁlm
was normal to the surface. However, Bragg reﬂection peaks did not shift in the ﬁlm grown with FEL irradiation. These results
indicate that the polymerized structure is a two-dimensional (2D) rhombohedron in the ﬁlm plane. We expect that the (0001)
surface of the oriented ﬁlm will be suitable for a 2D polymerized rhombohedron and that migration energy will contribute topolymerization at an elevated temperature. [DOI: 10.1143/JJAP.47.1412]
KEYWORDS: C60 , free-electron laser, FEL, polymerization, Ag(2) mode, Raman spectrum, rhombohedral structure

1.

Introduction

Polymerization processes for C60 molecules have beendeveloped for photo-, electron-beam-, pressure-, and plasma-induced methods. Among them, a crystalline threedimensional (3D) C60 polymer can be obtained at 15 GPa
and approximately 650  C from a two-dimensional (2D)
polymer; however, the 2D polymer remains in part andincludes a graphite like amorphous phase.1) The aim of ourstudy is to develop a novel photopolymerization process for
C60 molecules and synthesize an amorphous 3D C60 polymer
on the bulk scale for various applications. We expect thefeatures of the polymer to be a hardness level higher thanthat of diamond, ﬂexibility, and lightness.2) We chose a freeelectron laser (FEL) as a light source, which has uniquefeatures: a tunable wavelength in the infrared range, and a
20 ms macropulse containing several hundred micropulses
with a pulse width less than 1 ps.3) The FEL was generated at
the Laboratory for Electron Beam Research and Application(LEBRA) Institute of Quantum Science, Nihon University.
In previous reports, the wavelength dependence of polymerization degree was discussed for a pressed C60 powder, inwhich a 450 – 500 nm wavelength is most eﬃcient forpolymerization.4,5) However, the polymerization takes placeat the sample surface owing to the absorption. In this report,
because the FEL was irradiated onto the pressed C60 powder
surface directly, apparently polymerized grains were visible
in optical microscopy. In addition, ﬁlms were deposited with
FEL irradiation during growth to achieve bulk polymerization. Interestingly, polymerized ﬁlms were observed.
2.

Experimental Procedure

C60 powder was pressed at 600 MPa in a vacuum, ca.
106 Torr. The pressed C60 powder, called the bulk sample,
was removed, immediately set on a tungsten heater, andannealed at 130  C for 30 min in a vacuum. In Table I, thesample names and preparation conditions are summarized.
The FEL with a frequency of 2 Hz was irradiated onto thebulk sample surface through a quartz window after annealing without exposure to air. Normally, the FEL includes

Table I. Sample names and experimental conditions for synthesizing bulk
materials.
Wavelengthof irradiatedFEL (nm)Irradiation time(h)Sample B0

As-pressed——

Sample B1

500

1

6.8

Sample B2

500

2

8.4

Sample B3

500

4

10

Sample B4

500

4

578

Energy density[mJ/(pulsecm2 )]Table II. Sample names and experimental conditions for thin ﬁlms.
Substrate
temperature( C)

Wavelengthof irradiatedFEL (nm)

Energy density[mJ/(pulsecm2 )]Sample F1

40

—

—

Sample F2

40

1350

4.36

Sample F3

200

—

—

Sample F4

200

1350

4.18

harmonics; therefore, only the third harmonic at 500 nm was
obtained through a band-pass ﬁlter. The energy densities
listed in the table are of the fundamental wave, not of
harmonics. In sample B4, the FEL energy density wasincreased with a condenser lens. The energy density of thethird harmonics was less than 2.6 mJ/(pulsecm2 ) owing to
the lower limitation of the power meter.
In Table II, the growth conditions of C60 ﬁlms are
summarized. C60 powder was evaporated onto a cleaved
mica substrate, which was attached to a Cu foil to maintain athermal equivalent on the substrate surface. Film growth was
carried out for 30 min with the FEL irradiation perpendicular
to the substrate surface at ca. 107 Torr. The irradiated1350 nm FEL includes a third harmonic at 450 nm. It isnoted that no signiﬁcant diﬀerence has been observed whenirradiating pressed C60 specimens with a fundamental FEL
including harmonics and with only the third harmonic.6)1412

Jpn. J. Appl. Phys., Vol. 47, No. 2 (2008)

(b)

1.0

Intensity (arb. units)(a)N. IWATA et al.

10μm

(c)

(d)

0.8
0.6
0.4
0.2
0.0
1420

1440

1460

1480

1500

-1

Raman Shift (cm )Fig. 1. Optical microscopy images of (a) sample B0, as-pressed specimen; (b) sample B1, 500 nm FEL irradiation for 1 h; (c) sample B3,
irradiation for 4 h; (d) sample B4, 500 nm condensed FEL irradiation for
4 h. A scale bar is shown in the bottom right of (a). The circle in each
image is the area used for micro-Raman experiments. Polymerized whitegrains can clearly be seen in (b)–(d).

All the specimens were evaluated using micro-Raman
scattering (Kaiser HoloLab5000R Modular Research Microprobe System) and a second harmonic excitation yttrium
aluminum garnet (YAG) laser, 532 nm. The image of thearea for micro-Raman analysis was obtained using an optical
microscope. The excitation power was from 0.18 to
0.40 mW to prevent polymerization. The surface of ﬁlms
was investigated by scanning probe microscopy (SIISPI3800N-SPA400) in a dynamic force microscope (DFM)
mode, and X-ray diﬀraction (XRD; Rigaku RAD-C) analysis
was carried out.
3.

Results and Discussion

Optical microscopy images of the pressed C60 powder are
shown in Figs. 1(a) sample B0, 1(b) sample B1, 1(c)sample B4, and 1(d) sample B5. The area indicated by acircle in the ﬁgures is for detecting the micro-Raman
spectrum. All of the surfaces were packed with the tensof-mm grains. After irradiating the FEL, as shown in
Figs. 1(b)–1(d), white grains of approximately 5 mm appeared. The white grains brightened as a result of reﬂecting
the light of the optical microscope. The white grain density
increased as the energy density of the FEL increased, as
shown in Fig. 1(d). Figure 2 shows the Raman spectra ofsample B3 detected from the white grains (solid line) andanother area approximately 12 mm away from the whitegrain (dotted line), indicated by the circles in Fig. 1(c).
Intensity is normalized with the peak of an Ag(2) vibrationalmode. In addition to the intrinsic Ag(2) mode, Raman
spectrum detected from the white grain showed softened
Ag(2)-derived modes typically at 1460, 1454, and 1447
cm1 , indicated by arrows. The asymmetric peak at approximately 1470 cm1 in the dotted line was the sum of thepeaks at 1470 and 1464 cm1 . The Raman spectrum-like
dotted line in Fig. 2, showing a tail toward the low-energy

Fig. 2. The Raman spectra of sample B3 are shown. The spectra of thesolid and dotted lines, respectively, were detected at the white area andca. 12 mm away from the area in Fig. 1(c). In the solid line, in addition to
the intrinsic Ag(2) mode, typical peaks at 1460, 1454, and 1447 cm1
appeared. In the dotted line, a tail toward a lower wavenumber wasobserved.

side, was observed for all of the surfaces except for thewhite grain. In both specimens, the intrinsic Ag(2) moderemained.
Irradiating the FEL resulted in the formation of C60
polymers, demonstrated by the appearance of additional
peaks at 1460, 1454, and 1447 cm1 , in particular in thewhite grains. The peaks at 1460 and 1454 cm1 are typicalpeaks for photopolymers and the peak at 1447 cm1 isobserved in photopolymers and rhombohedral-type polymers.7) However, the polymerization took place around thesample surface owing to the remnant of the intrinsic Ag(2)peak at approximately 1470 cm1 . The penetration depth of
the FEL was expected to be shallower than that of theexcitation laser in the Raman experiment.
In Fig. 3 all the Raman spectra are shown. The excitation
laser energy was calibrated to be 1467 cm1 for the Ag(2)mode using the as-pressed bulk specimen. A symmetricAg(2) mode peak was observed in an as-pressed specimen,(a) sample B0. The Raman spectra illustrated in Figs. 3(b)–3(e) were all obtained from the white grains found in
Figs. 1(b)–1(d). As shown in Fig. 3(b), sample B1, andFig. 3(c), sample B2, diﬀerent Raman spectra were observed, even though the spectra were detected for similar
white grains. Reproducibility was not conﬁrmed. By
increasing irradiation time and/or energy density, the same
Raman spectrum with a typical photopolymerized peak at
1460 cm1 was obtained at all white grains, indicated by thedotted line.
Figure 4 shows the XRD spectra of sample B0 (dottedline) and sample B3 (solid line). Apparently all the Bragg
reﬂections shifted toward a value higher than 2 insample B3 compared with those of sample B0, indicating
the compression of lattice spacing.
From the results mentioned above, polymerization took
place at the sample surface, particularly at the white grains.
Polymerization was promoted by increasing irradiation timeand energy density. In general, photopolymerization occurs
at a moderate optical ﬂux higher than 5 W/cm2 and thedegree of polymerization depends on the irradiation time on1413

Jpn. J. Appl. Phys., Vol. 47, No. 2 (2008)

N. IWATA et al.

3

Intensity (arb. units)

Intensity (arb. units)

43 (e) Sample B4−(d) Sample −B32(c) Sample − B2

1 (b) Sample −B1(a) Sample − B0

0
1420

2
(c) Sample−F3

1 (b) Sample−F2(a) Sample−F1

1440

1460

1480

0
1420

1500

Raman Shift (cm-1)

(220)

(111)

1440

1450

1460

1470

1480

1490

1500

Fig. 5. Raman spectra of C60 ﬁlms grown at (a) 40  C, sample F1; (b)40  C with 1350 nm FEL irradiation during growth, sample F2; (c)200  C, sample F3; and (d) 200  C with 1350 nm FEL, sample F4. Even at
the low substrate temperature, the peak at 1460 cm1 appeared in (a). Asgrowing ﬁlms with the FEL irradiation polymerization phase were
dominant in (b) and (d). At 200  C with 1350 nm FEL irradiation, a peakat 1455 cm1 can clearly be seen.

(311)

1.4
1.2

1430

Raman Shift (cm-1)

Fig. 3. The Raman spectra of all the specimens are shown. In (b),sample B1, and (c), sample B2, although the spectra were detected fromsimilar white grains, reproducibility was not obtained. In (d), sample B3,and (e), sample B4, at all white grains, the same Raman spectrum wasobserved with the additional peak at 1460 cm1 , which is typical for C60
photopolymers.

(a)

(b)

1.0
0.8

(222)

Intensity (arb.units)(d) Sample−F4

0.6
0.4
0.2
0.0

6

8

10

12

14

16

18

20

22

24

0.0
nm

2θ (deg)
Fig. 4. The XRD results of sample B0 (dotted line) and sample B3 (solidline) are shown. Bragg reﬂections for the fcc structure can clearly beobserved and indexed. With FEL irradiation, it was found that the
specimen was compressed.

the order of hundreds of hours.7–9) Considering the macropulse width (20 ms) of the FEL, the third FEL harmonic is0.13 W/cm2 at most. It is expected that one of the featuresof FEL, namely, the extremely small micropulse width,
prevents the laser power from diﬀusing to the underlayer asphonons. Therefore, the power concentrates and is used topromote a phototransformation. In addition, all the areas,
except for the white one, showed a tail toward the lower
wavenumber in Fig. 2 (dotted line), which indicate that theareas act as a precursor to white grains.
In Fig. 5, the Raman spectra of the ﬁlms are shown.
The excitation energy was calibrated to be 1469 cm1 forthe intrinsic Ag(2) mode. At 40  C, for (a) sample F1, an
additional peak at approximately 1460 cm1 appeared.
Because a thin ﬁlm was deposited with FEL irradiation at
40  C for (b) sample F2, the peak value at 1460 cm1
became higher than that of the original Ag(2) peak, which
seldom occurs in bulk samples. At 200  C, for (c) sample F3,
peaks at not only 1460 cm1 but also 1455 cm1 appeared.

9.2
nm

0.0
nm

12.9
nm

Fig. 6. (Color online) Surface images of (a) sample F3 and (b) sample F4
are shown. Although the number of nucleation sites was the same, thegrain height in (b) was higher than that in (a), giving rise to a rougher
surface. FEL irradiation is expected to enhance the polymerization, and,as a result, the migration length in the ﬁlm (b) became shorter than that in(a). In (a), the surface showed the presence of grains on a ﬂat surface,which was composed of coalesced grains, whereas in (b), grainboundaries can clearly be observed.

With FEL irradiation at 200  C for (d) sample F4, thedominant peak was not 1469 cm1 and a robust downshift
was observed at approximately 1455 cm1 .
Figure 6 shows the 5 mm2 surface images of the ﬁlms
grown at 200  C (a) without FEL irradiation and (b) with
irradiation. Steps and a terrace structure can clearly be seenin Fig. 6(a) with facets. The shapes of the grains were
triangular and/or hexagonal. Grains growing around thenucleation site, the density of which was ca. 400  400
mm2 , coalesced with adjacent grains and then became a ﬂat
surface with an average surface roughness (Ra ) of 1.03 nm.
With irradiation in Fig. 6(b) for sample F4, the density of
the nucleation site was almost the same, but coalescence did
not occur to a large extent compared with that in the case ofsample F3, giving rise to a larger grain height and rougher1414

Jpn. J. Appl. Phys., Vol. 47, No. 2 (2008)

N. IWATA et al.

surface with an Ra of 1.59 nm. Grain boundaries were clearlyobserved. The FEL-irradiated ﬁlm surface is approximately
the same as that of a ﬁlm grown at an ca. 50 lower substratetemperature without FEL irradiation.10) The XRD spectrumof the ﬁlm grown at 200  C showed (000n) (n ¼ 2, 4, 6)
Bragg reﬂections and the result indicated a hexagonal closed
packed (hcp) structure, which was the same as that shown ina previous report.10) The [0001] direction was normal to thesubstrate surface. However, no peak shift was observed in
the ﬁlm with FEL irradiation, indicating that polymerization
did not occur along the direction normal to the surface.
The Raman spectra showed the polymerization, whereas
the XRD result did not, in the ﬁlm. A plausible interpretationof this inconsistency is the synthesis of 2D rhombohedralpolymerization in the ﬁlm plane at 200  C. As shownin Figs. 5 and 6, elevated substrate temperature and FEL
irradiation promoted the polymerization. By increasing thesubstrate temperature, the oriented ﬁlm with the hcp
structure showed the (0001) plane on the ﬁlm surface during
growth. In high-temperature high-pressure (HTHP) synthesis
processes for polymers, rhombohedral (R), tetragonal (T),and orthorhombic (O) polymerized phases are reported.7)
The (0001) surface is expected to be suitable for thesynthesis of a R polymerized phase owing to the packing
structure of C60 molecules. Considering the results ofpolymerization in the ﬁlm grown at 40 and 200  C without
FEL irradiation and that growth was promoted at 200  C, as
shown in Figs. 5(a) and 5(c), it is reasonable to conclude that
migration energy contributes to polymerization. When the
migrating molecules incorporate in a kink or a step, themigration energy is expected to transfer to the energy forpolymerization. Polymerized molecules were not easily
released and showed limited migration; consequently, thesurface became rough. From the results of the Raman
spectra, polymerization was promoted in the ﬁlm by
irradiating the FEL during growth.
4.

Conclusions

We propose a novel photopolymerization process with the
FEL irradiation of pressed C60 powder and ﬁlms tosynthesize 3D amorphous C60 polymers for application toa novel hard, light, and tough material. White grains in
optical microscopy showed the polymerized Raman spec-

trum with additional peaks at 1460, 1454, and 1447 cm1 .
Reproducibility for the grains was improved by increasing
the irradiation time, and the white-grain density wasincreased with condensed FEL irradiation. The mechanismof the polymerization with low power FEL irradiation isconsidered to be an extremely sharp micropulse. As
irradiating the FEL during ﬁlm growth caused downshifting,
the Ag(2)-derived mode became dominant, in particular,when grown at 200  C. No changes in Bragg reﬂections
along the perpendicular to the surface exhibited nonpolymerization along the direction. The expected polymer
structure was a 2D rhombohedron in the plane. The peak
appearance at 1455 cm1 in the ﬁlm grown at 200  C with
FEL irradiation was lower than the typical photopolymer
peak at 1460 cm1 , indicating the strong correlation between
adjacent C60 molecules in the plane.
Acknowledgments
This work was partially supported by the ‘‘Academic
Frontier’’ Project for Private Universities: matching fundsubsidy from Ministry of Education, Culture, Sports, Scienceand Technology (MEXT), 2000 – 2007. The authors are
grateful for the ﬁnancial support of an Interdisciplinary
General Joint Research Grant (Nihon University).

1) S. Yamanaka and A. Syouji: Kotai Butsuri 41 (2006) 407[in Japanese].
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253 (2007) 7977.
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published in Trans. Mater. Res. Soc. Jpn. (2007).
6) N. Iwata and H. Yamamoto: private communication.
7) P. C. Eklund and A. M. Rao eds.: Fullerene Polymers and Fullerene
Polymer Composites (Springer, Berlin, 1999) Chaps. 1 and 4.
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Wang, W. T. Lee, X.-X. Bi, P. C. Eklund, D. S. Cornett, M. A.
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Applied Surface Science 253 (2007) 7977–7980
www.elsevier.com/locate/apsusc

Photon-assisted synthesis of C60 polymers by laser irradiation
Hiroshi Yamamoto *, Nobuyuki Iwata, Ryoji Hashimoto, Shingo Ando
College of Science and Technology, Nihon University, 7-24-1 Narashinodai, Funabashi-shi, Chiba 274-8501, Japan
Available online 2 March 2007

Abstract
The harmonics of a free electron laser (FEL) were irradiated in vacuum to surfaces of compressed C60 and a mixture of C60 and I2. The power
and frequency of the fundamental FEL macro-pulse were ca. 0.5 mJ/pulse and 2 Hz, respectively. The irradiation time was 120–180 min. After
irradiation of FEL with a typical wavelength of 450 or 345 nm, the Raman peak of Ag(2)-derived vibration mode of C60 shifted to the lower-energy
side. The Raman peak shift of the mixture powder sample was greater than that of pure C60. Furthermore, changes of the crystalline structure
indicated that various intermolecular combinations occurred by irradiation. These results strongly suggest that three-dimensional polymerizationof C60 was promoted by laser irradiation and the effect of photon-assisted hole-doping from iodine atoms to C60 molecules.
# 2007 Elsevier B.V. All rights reserved.
Keywords: C60; Polymerization; Free electron laser (FEL); Hole-doping; Raman spectroscopy

1. Introduction
Since the polymerization of C60 films was first accomplished using phototransformation [1], many studies haveexamined C60 polymer synthesized using various excitation
processes [2] such as electron-beam irradiation [3], highpressure [4], plasma-induced excitation [5], and so on.
Because of those adequate chemical or physical excitations, acarbon double bond in a hexagon of a C60 cage is broken; ittransforms into an intermolecular four-fold ring with aneighboring molecule. Consequently, two C60 molecules are
combined and change into a dimer or some higher order ofpolymer. This process is the fundamental mechanism for C60
polymerization.
The polymerized fullerite structures based on a ‘‘2 + 2
cycloaddition’’ bonding [6] were observed experimentally andtheoretically [7–9]. One-dimensional (1D) chains revealed anorthorhombic phase; two-dimensional (2D) planes include atetragonal and a rhombohedral phase. Few experimental data
[4] have been reported about the possibility of a threedimensional (3D) polymerized phase, but 3D C60 polymers are
expected to be extremely hard, with lower density and lessbrittleness than diamond.

* Corresponding author. Tel.: +81 47 469 5457; fax: +81 47 469 5457.
E-mail address: hyama@ecs.cst.nihon-u.ac.jp (H. Yamamoto).
0169-4332/$ – see front matter # 2007 Elsevier B.V. All rights reserved.
doi:10.1016/j.apsusc.2007.02.155

The polymerizations usually proceed by compression at theorder of gigapascals or heating to greater than a thousand
degrees [4]. The C60 cages partially collapse under stringent
pressure conditions. These polymerization processes are notwell-controlled and are unsuitable to supply sufficiently large
3D C60 polymers for applications.
This work is intended to develop a novel photon-assisted
process for synthesis of C60 polymers by laser irradiation undermoderately high pressure. We have also described a hole doping
effect, as demonstrated in the synthesis of diamond fromgraphite [10]. We expect that this effect is applicable to thepolymerization reaction of C60 molecules because of themutually similar structures of carbon networks in graphite and
C60.
Hereinafter, the newly developed photon-assisted process isintroduced and the possibility of synthesis of 3D C60 polymers
by free electron laser (FEL) irradiation will be demonstrated.
2. Experimental
Fig. 1 shows the schematic apparatus developed forsynthesis of C60 polymers. The vacuum chamber was evacuatedto ca. 10 4 Pa using a turbo molecular pump. The pristine
powder was set in an anvil placed in vacuum. The powder wascompressed via a bellows using a hydraulic press. Before laser
irradiation, the sample was pre-annealed at ca. 500 K for 20 h
in a vacuum chamber. Pressure, usually ca. 470 MPa, was7978

H. Yamamoto et al. / Applied Surface Science 253 (2007) 7977–7980

The obtained samples were 5-mm-diameter disks withthickness of less than 0.5 mm. They were packed in a vacuum
box to prevent contamination from air. Raman microscopic
spectroscopy (SYSTEM1000; Renishaw plc.) was used to study
molecular vibration spectra. The light source used for
excitation was a 514.5 nm Ar laser. Crystalline structures
were studied using reflective X-ray diffraction (XRD) (RAD-C;
Rigaku Corp.) with Cu Ka1 X-rays.
3. Results and discussion

Fig. 1. Schematic apparatus constructed for synthesis of C60 polymers. The
compressed surfaces of the sample were irradiated using a laser by the BK7
window. Aluminum cold mirrors were used for high reflectance of visible laser(wavelength less than ca. 700 nm) and for cutting near-infrared light. The quartz
BK7 ensures high transmittance of greater than 95% of visible light.

applied to the sample powder using a quartz BK7 window.
Visible laser light was introduced to the surface of the pressed
sample using Al cold mirrors that reflect only light withwavelength of less than ca. 700 nm.
Since the 1980s, mid-infrared (wavelength of a few tens of
micrometers) FELs have been developed and applied tonumerous studies of laser-solid interaction [11]. We adopted an
FEL system that generates a tunable near-infrared laser andhigher harmonics. Details of the FEL system used there haveappeared elsewhere [12]. A macro-pulse of a few tens of
microseconds, which included micro-pulses of a few-hundredpicoseonds, constituted the FEL pulse, which was repeated in
2 Hz. The power of the fundamental macro-pulse was ca.
0.5 mJ/pulse. The irradiation time was typically 120–180 min.
Because the FEL can supply a laser with a tunable wavelength,
the fundamental wavelength was chosen arbitrarily for thiswork. The laser used for irradiation was the third harmonic of
FEL with a fundamental wavelength of ca. 1350 nm. The
wavelength that was mainly noticed was ca. 450 nm, of which
the value was evaluated from the difference of the bonding
energy of the C C double bond and the C–C bond.
Furthermore, an ultraviolet laser (345 nm) was also takenout as the fifth harmonic of ca. 1700 nm fundamental FEL.
Two kinds of pristine powder were prepared. Pure C60(99.95%) was a commercial product. The C60 was also mixed
with I2 at the molar ratio of 1:4. Iodine was used forenhancement of polymerization reactions expected from thehole-doping effect. Because the electron affinity of the C60 ishigh, we aimed at promotion of electron transfer from C60 toiodine atoms, assisted by photon-excitation.
For a C60 molecule and its polymer, the optimized molecular
structure and/or the molecular orbital was semi-empiricallycalculated using WinMOPAC (3.9 Professional; Fujitsu Co.
Ltd.). Then molecular vibration modes were evaluated forinvestigation of the observed Raman spectra.

Our calculation of vibration modes indicated that the Raman
spectrum broadened and shifted to a lower energy because of
the polymerization reactions. A Raman Ag(2) mode of a C60
molecule is defined as a contracting vibration of pentagon rings.
Fig. 2 shows results of schematic vibration analysis of theAg(2)-derived mode in a C60 monomer, a 2 + 2 dimer, and atrimer. When a C60 molecule is combined with anothermolecule, the Ag(2) modes changed into other vibration modes
because of lower symmetry or vibration reduction. The
polymerization of the C60 monomer engendered the Raman
peak shift of the Ag(2)-derived mode to the lower energy. Our
calculations indicate that the amounts of the Raman peak shift
obtained in the dimer and the trimer, as shown respectively in
Fig. 2, were ca. 9 and 16 cm 1, thereby confirming
qualitatively that the Raman peak broadens and shifts to thelower-energy side according to the progress of polymerization.
The Ag(2)-derived mode is therefore noted as an index of thepolymerization reaction progress.
Fig. 3 shows typical Raman peaks of the Ag(2)-derived
mode in the case of 450 nm FEL irradiation. The wave numberof the peak of the non-irradiated pristine C60 was 1463 cm 1,
whereas that of the irradiated C60 changed to 1460 cm 1 with a

Fig. 2. Schematic molecular structures and Ag(2)-derived vibration modes
calculated in: (a) a C60 monomer, (b) a dimer, and (c) a trimer. The dimer ortrimer is formed by a ‘‘2 + 2’’ cycloaddition reaction between parallel carbon
double bonds on two adjacent molecules. The amplitudes of the vibration of
constituent C atoms are represented schematically as thick lines.

H. Yamamoto et al. / Applied Surface Science 253 (2007) 7977–7980

7979

Fig. 3. Raman peaks of the Ag(2)-derived vibration modes observed in thethree samples: non-irradiated pristine C60, C60, and C60 + I2 irradiated by
450 nm FEL.

Fig. 5. XRD patterns of the three samples: non-irradiated pristine C60, C60, andC60 + I2 irradiated by 450 nm FEL. The inset index corresponds to each fcc planeof C60. The (1 1 1) peak intensity of the irradiated C60 + I2 decreased remarkably.

decrement by ca. 3 cm 1. On the other hand, the irradiatedmixture of C60 + I2 revealed the peak wave number of1455 cm 1. The Raman shift was as great as ca. 8 cm 1.
The half-width of the Raman peak also increased incorrespondence to the amount of the peak shift.
Nearly identical results were obtained for samples irradiatedby FEL with various wavelengths of 400–500 nm. No distinct
difference was observed in the Raman spectra, even bychanging the wavelength.
However, the ultraviolet rays caused somewhat differentchanges of the Raman spectrum. Fig. 4 shows Raman peaks of
the Ag(2)-derived mode in the case of 345 nm FEL irradiation.
Double splitting of the peak by the irradiation was observed:
C60 molecules are known to reveal large optical absorption at
wavelengths of 221, 271, and 347 nm [13]. Because the opticaldensity at 345 nm is about three times higher than that at
450 nm, polymerization reactions might proceed better in thethinner region of sample surfaces by 345 nm irradiation than by

450 nm irradiation. Consequently, the Raman peak of nonpolymerized C60 also remained in the case of 345 nm
irradiation. The spectrum shown in Fig. 4 revealed the splitting
peak. The amounts of the Raman shift observed were ca.
7 cm 1 in the C60 sample and ca. 10 cm 1 in the C60 + I2
sample. The peak decreased, broadened, and came to resemble
that of 450-nm irradiation.
It is noteworthy that the change of the Raman peak of theC60 + I2 sample was definitely larger than that of pure C60. This
result implies that some additional mechanisms work for
promotion of polymerization by addition of I2. It can be inferredthat photo-excited electrons of C60 molecules transfer to iodine
atoms: hole-doping into C60 molecules occurs.
The hole-doping effect was also confirmed through crystalline structure analyses. Fig. 5 shows typical XRD patterns ofsamples irradiated using 450 nm FEL. The d values of (1 1 1)
planes of the irradiated C60 and C60 + I2 sample were smaller by
7  10 3 nm and 8  10 3 nm, respectively, than that of the
non-irradiated C60. The decrement of the lattice parameter was
about 1%. Furthermore, the main peak (1 1 1) markedly
decreased and broadened in the case of the C60 + I2 sample.
That result suggests that binding between C60 molecules took
place and that the decline of the crystalline structure progressed
according to polymerization reactions induced by FEL
irradiation and/or addition of iodine.
4. Conclusion

Fig. 4. Raman peaks of the Ag(2)-derived vibration modes observed in thethree samples: non-irradiated pristine C60, C60, and C60 + I2 irradiated by
345 nm FEL.

The FEL, usually of 450 or 345 nm wavelength, wasirradiated to surfaces of compressed C60 or C60 + I2 powder invacuum. Simulation of molecular vibration modes revealed that
the Raman peak of the Ag(2)-derived mode shifts to the lowerenergy side and/or the half-width of the peak broadensconcomitant with progress of the polymerization reaction. A
mixture sample C60 + I2 irradiated by 450 nm FEL showed a
Raman shift of ca. 8 cm 1, of which the value was distinctly
larger than that of the pure C60 sample. The XRD results
revealed that the average distance between C60 molecules

7980

H. Yamamoto et al. / Applied Surface Science 253 (2007) 7977–7980

shortened by about 1% and the crystalline structure wasdegraded by the FEL irradiation. The obtained results suggestthat 3D polymerization of C60 was promoted using laser
irradiation and/or the effect of the photon-assisted hole-dopingfrom iodine to C60 molecules. Detailed characterization of theobtained samples is a subject for future work.
Acknowledgements
This work was partly supported by the ‘‘Academic Frontier’’
Project for Private Universities: a matching fund subsidy from
MEXT, 2000–2004. The authors would like to thank the staff ofLaboratory for Electron Beam Research and Application of
Nihon University for assistance with FEL operations.
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[3] Y.B. Zhao, D.M. Poirier, R.J. Pechman, J.H. Weaver, Appl. Phys. Lett. 64(1994) 577.
[4] H. Yamawaki, M. Yoshida, Y. Kakudate, S. Usuba, H. Yokoi, S. Fujiwara,
K. Aoki, R. Ruoff, R. Malhotra, D.C. Lorents, J. Phys. Chem. 97 (1993)
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[7] A.M. Rao, P.C. Eklund, J.-L. Hodeau, L. Marques, M. Nunez-Regueiro,
Phys. Rev. B 55 (1997) 4766.
[8] V.A. Davydov, L.S. Kashevarova, A.V. Rakhmanina, V. Agafonov, H.
Allouchi, R. Ceorin, A.V. Dzyabchenko, V.M. Senyavin, H. Szwarc, Phys.
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[9] X. Chen, S. Yamanaka, K. Sako, Y. Inoue, M. Yasukawa, Chem. Phys.
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[10] H. Nakayama, H. Katayama, H. Yoshida, Jpn. J. Appl. Phys. 41 (2002)
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[11] H.P. Freund, J.T.M. Antonsen, Principles of Free-Electron Lasers, Chapman & Hall, London, 1992.
[12] Y. Hayakawa, I. Sato, K. Hayakawa, T. Tanaka, K. Yokoyama, T. Sakai, K.
Kanno, K. Ishiwata, E. Hashimoto, Nucl. Instrum. Methods A507 (2003)404.
[13] S. Kazaoui, R. Ross, N. Minami, Solid State Commun. 90 (1994) 623.

Japanese Journal of Applied Physics
Vol. 47, No. 1, 2008, pp. 546–549#2008 The Japan Society of Applied Physics

Crystal Growth of Magnetoelectric Cr2 O3 Thin Film on Sapphire and SrTiO3
Takeshi A SADA, Kenjirou NAGASE, Nobuyuki I WATA, and Hiroshi Y AMAMOTO
College of Science and Technology, Nihon University, Funabashi, Chiba 274-8501, Japan(Received July 12, 2007; accepted September 18, 2007; published online January 22, 2008)

Representative magnetoelectric (ME) material Cr2 O3 ﬁlms were prepared on R-cut sapphire (R-Al2 O3 ) and SrTiO3 (STO)
substrates with the aim of applying them in perovskite oxide electronic devices. Cr2 O3 thin ﬁlms were deposited using the
DC–RF hybrid magnetron sputtering system at diﬀerent substrate temperatures. When the R-Al2 O3 substrate temperature was
550  C, Bragg peaks ðn; n; 0; 2nÞ were observed. The full width at half maximum (FWHM) of the rocking curve was smallest
at Ts ¼ 550  C with a value of 0.52 . The grain size of Cr2 O3 increased with substrate temperature in the 400 – 600  C range.
The lengths of rectangular Cr2 O3 grains were in the ranges of 200 – 600 nm along the ½1 101 direction and 100 – 300 nm alongthe ½112 0 direction. Furthermore, Cr2 O3 ﬁlms were deposited on SrTiO3 (100) with CeO2 buﬀer layers. The Cr2 O3 thin ﬁlm
deposited on the smooth CeO2 buﬀer layers with the average roughness ðRa Þ ¼ 0:41 nm showed a c-axis orientation;
moreover, a Cr2 O3 (0001)/CeO2 (100) == STO(100) crystal relationship was obtained. [DOI: 10.1143/JJAP.47.546]
KEYWORDS: magnetoelectric effect, sputtering, Cr2 O3 , sapphire, CeO2 buffer, SrTiO3

magnetron sputtering. The sputtering conditions for CeO2
buﬀer in sample A were Ts , RF power, and sputteringpressure of 700  C, 80 W and 9 Pa (Ar : O2 ¼ 1 : 1), respectively. The sputtering conditions for CeO2 in sample B were
Ts , RF power, and sputtering pressure of 380  C, 50 W, and6 Pa (Ar : O2 ¼ 1 : 1), respectively. The CeO2 == STO ﬁlmof sample B was annealed at 1000  C for 3 h in air. The
surface average roughness (Ra ) of CeO2 buﬀer layers wasinvestigated before the sputtering of Cr2 O3 ﬁlms.
The crystalline structures of the ﬁlms were investigated by
X-ray diﬀraction analysis (XRD; Riguku RAD-C system)with the –2 scan and rocking curve at 2. The surface
morphology was observed by scanning probe microscopy inthe dynamic force microscopy (DFM) mode (Nanonavistation SPI3800N and SPA400 unit; Seiko Instruments).

Introduction

The magnetoelectric (ME) eﬀect is deﬁned as theappearance of magnetic ﬁelds induced by electric ﬁelds.
When electric ﬁelds are applied, the changes in the relative
position between ions result in changes in magnetic
superexchange interactions and then magnetic ﬁelds appear
outside.1–3) As a representative ME material, antiferromagnetic insulator Cr2 O3 has been studied.4) Because magnetic
ﬁelds are induced along the direction of electric ﬁelds
applied, the ME coeﬃcients are diﬀerent in the directions of
the a-, b-, and c-axes.1,3) The coeﬃcient in the c direction ishighest at room temperature as compared with those in theother axes. In contrast, the ME coeﬃcient along the adirection is higher than that along the c-direction at a lowtemperature. Therefore, investigation of the crystal orientation of Cr2 O3 is important for application of the ME eﬀect.
We expected that ME materials can be applied to novel
electronic devices, particularly in combination with a seriesof perovskite oxides showing various functional properties,
such as superconductivity, ferroelectricity, and ferromagnetism. Most perovskite oxide ﬁlms have been prepared on
cubic perovskite crystal substrates as SrTiO3 (STO). However, lattice matching is not good between cubic STO andrhombohedral Cr2 O3 . In order to combine Cr2 O3 with aseries of perovskite oxides devices, Cr2 O3 ﬁlms should be
deposited on STO with cubic CeO2 buﬀer layers.5–7) The
purposes of this work are to optimize substrate temperaturefor sputtered Cr2 O3 ﬁlms on R-Al2 O3 and to investigate ﬁlm
growth on STO with CeO2 buﬀer layers.

Results and Discussion

10

3

10

2

10

1

10

20

30

40

50

60

70

80

Al2O3 (3306)

4

Al2O3 (2204)

10

Cr2O3 (3306)

Cr2 O3 ﬁlms were deposited on R-Al2 O3 (11 02) by oﬀ-axis
DC–RF hybrid magnetron sputtering. Substrate temperature(Ts ) was changed from 300 to 600  C. Other typical growth
conditions were DC current, RF power, and sputteringpressure of 0.04 A, 100 W, and 0.26 Pa (Ar : O2 ¼ 4 : 1),
respectively.
As a substrate, a R-Al2 O3 single crystal (5  10 mm2 ) was
etched by HF (HF : H2 O ¼ 1 : 7) for 10 s, and subsequently
annealed at 1000  C for 12 h in air.
Two samples (samples A and B) of Cr2 O3 /CeO2 bilayers
were deposited on the STO substrates by oﬀ-axis RF5

Cr2O3 (2204)

Experimental Procedure

10

Al2O3 (1102)

3.1 Cr2 O3 on R-Al2 O3
The XRD pattern of Cr2 O3 /R-Al2 O3 at Ts ¼ 550  C isshown in Fig. 1. The Cr2 O3 (11 02) and Al2 O3 (11 02) Bragg
diﬀraction peaks were observed and the crystal relationshipbetween Cr2 O3 and Al2 O3 was Cr2 O3 ð11 02Þ == Al2 O3
ð11 02Þ. Similar XRD patterns were observed in the ﬁlms
prepared at Ts ¼ 300 { 600  C.

Intensity (arb. unit)

2.

3.

Cr2O3 (1102)

1.

90

2θ (deg)
Fig. 1. XRD pattern of Cr2 O3 deposited at Ts ¼ 550  C on R-Al2 O3 . The
crystal relationship between Cr2 O3 and R-Al2 O3 was Cr2 O3 ð11 02Þ ==Al2 O3 ð11 02Þ.

546

Jpn. J. Appl. Phys., Vol. 47, No. 1 (2008)

T. ASADA et al.

Full width at
half maximum ∆ θ

1.00

3.635
Lattice constant∆ θ (deg)

0.90
0.80

3.630

0.70

3.625

0.60
3.620
0.50
3.615

0.40
300

350

400

450

500

550

600

The small value of  at higher Ts indicates that high Ts
induces crystal orientation along the normal direction for
substrates. Crystal degradation might proceed at Ts above
550  C.
Two typical DFM images (2  2 mm2 ) of R-Al2 O3
substrate surfaces are shown in Figs. 3(a) and 3(b). All thesubstrates showed a step and terrace structure. From the line
proﬁle shown in Figs. 3(a) and 3(b), the widths of eachterrace were in the 180 – 200 nm range and 400 nm,
respectively.
Figures 3(c)–3(f) show the 2  2 mm2 DFM surface
images and line proﬁles of Cr2 O3 == R-Al2 O3 deposited
at Ts ¼ 300 { 600  C. Table I shows grain size (along½112 0  along ½1 101) as a function of Ts . Figure 3(c)shows the ﬁlm deposited at Ts ¼ 400  C on the substrateshown in Fig. 3(a). Rectangular grains appeared with sizes
in the 100  ð100 { 200Þ nm2 range. The diﬀerence in height
between grains was approximately 3.5 nm, as determinedfrom the line proﬁle shown in Fig. 3(c). Figure 4(d) shows
the ﬁlm deposited at Ts ¼ 500  C on the substrate shown in
Fig. 3(b). The rectangular grains became larger along the½112 0 direction and coalesced. The lengths of grains were in
the 200 – 300 nm range along the ½1 101 direction the sameas the widths of substrate terraces shown in Fig. 3(b). On theother hand, the lengths of grains were in the 200 – 300 nm
range along the ½112 0 direction. Grain boundaries wereobserved and their lengths were longer along the ½112 0

Lattice constant of Cr2O3(1102) (Å)

3.640

1.10

Substrate temperature Ts (°C)
Fig. 2. Ts dependence of Cr2 O3 [11 02] lattice constant and  (FWHM)
of rocking curve at Cr2 O3 (11 02) reﬂection. The left-side index and theright-side index show FWHM of rocking curve reﬂection and lattice
constant of Cr2 O3 (101 2) plane, respectively.

Figure 2 shows the Ts dependence of Cr2 O3 (11 02) lattice
constant and full width at half maximum (FWHM) of
rocking curve reﬂection (). The lattice constant wasalmost constant from Ts ¼ 300 to 400  C and lattice constantincreased from Ts ¼ 400 to 500  C. The value of 
decreased monotonically with increasing Ts from Ts ¼300  C. The smallest value was 0.52 at Ts ¼ 550  C.

(a)

(c)

(b)

181nm

[1120]

401nm

0.33
532nm

(d)

19.4
nm

225nm

0.54
647nm

(e)

232nm

9.7nm
nm

1.14
nm

0.88
nm

4.6nm
515nm

(f)

0.33~0.38nm

3.7
637nm

727nm

21.0
nm
18.0

156nm

14.6
nm
13.4

[1101]
Fig. 3. DFM surface images (2  2 mm2 ) of R-Al2 O3 substrates are shown in (a) and (b), as well as DFM surface images of Cr2 O3 ﬁlms
deposited at (c) Ts ¼ 400  C, (d) Ts ¼ 500  C, (e) Ts ¼ 550  C, and (f) Ts ¼ 600  C. The ﬁgures below the surface images are line
proﬁles. The arrows from the bottom to the top and the left to the right show ½112 0 and ½1 101 crystal directions, respectively. The
arrows in the line proﬁle show the width of terrace, step height and/or grain length. The grain size of Cr2 O3 increased with Ts in the
400 – 600  C range. The lengths of Cr2 O3 grains were in the range of 200 – 600 nm along the ½1 101 direction and 100 – 300 nm alongthe ½112 0 direction.

547

Jpn. J. Appl. Phys., Vol. 47, No. 1 (2008)

100  200
300  300

550

300  600

600

400  400

1000

direction. About 30 grain boundaries were observed in the
DFM image of 2  2 mm2 . Figure 3(e) shows the ﬁlm
deposited at Ts ¼ 550  C on the substrate shown in Fig. 3(b).
The grains were clearly larger along the ½1 101 direction.
The lengths of grains were in the 200 – 600 nm range alongthe ½1 101 direction. Namely, the grain size was larger than
the width of terraces observed in Fig. 3(b). As shown in theline proﬁle, several steps along the ½1 101 direction wereobserved in the one-grain plane. The step heights were in the
0.3 – 0.4 nm range, the same as that of 1/3 unit cell of Cr2 O3
with the R-plane (about 0.363 nm). About 21 grain boundaries were observed in the DFM image of 2  2 mm2 . In theﬁlm grown at 600  C, grains of various sizes grew along the½112 0 and ½1 101 directions. The grain size was 400 
400 nm2 . Bunching steps, the height of which was approximately 5 nm, were observed, as shown in Fig. 3(f). The
surface smoothness was degraded.
Grain size is expected to increase with Ts because surface
energy is decreased by large islands with decreasing island
boundaries and increasing migration length. The rectangulargrain sizes at Ts ¼ 400 { 550  C were restricted by steps andterraces of R-Al2 O3 shown in the Figs. 3(a) and 3(b).
Therefore, at Ts ¼ 400  C, coalescence did not occur and thesurface was covered by many grains of 200 – 400 nm sizes
the same as the width of terraces. At Ts ¼ 500  C, grain size
along the ½112 0 direction was increased owing to a large
migration length. Therefore, coalescence occurred along the½112 0 direction. However, the length of grains along the½1 101 direction was the same as the width of terraces (about
400 nm), as shown by the line proﬁle in Fig. 3(b), because
sputtered ﬂux atoms were not able to move over the steps of
R-Al2 O3 . At Ts ¼ 550  C, migration length increased andatoms moved over the steps of R-Al2 O3 . Therefore, larger
grains along the ½112 0 and ½1 101 directions were obtained.
As a result, the density of grain boundaries decreased. At
Ts ¼ 550  C, the terrace in one grain was composed ofseveral steps of 1/3 unit cell at most, showing a twodimensional island growth. At Ts ¼ 600  C migration length
along the ½112 0 and ½1 101 directions became isotropic,
because of suﬃciently migration energy for moving over thesteps. Therefore, grains of various sizes were observed at
Ts ¼ 600  C.
The surface of the ﬁlm prepared at Ts ¼ 550  C wassmooth and revealed a two-dimensional growth. The resultsof XRD analysis and AFM indicated that about Ts ¼ 550  C
was the optimum for obtaining the highest quality of a
Cr2 O3 (11 02) crystal, and a large ME eﬀect is expected.
3.2 Cr2 O3 on STO with CeO2 buﬀer
Figure 4 shows the XRD pattern of Cr2 O3 on the STO
substrate with two kinds of CeO2 buﬀer layer (samples A
548

800

Sample B

600

Cr 2O 3 (00012)

400
500

Cr 2O 3 (0006)

Grain size
(nm2 )

CeO 2 (200)

Ts
( C)

STO diffraction peaks

1200

CeO 2 (400)

Ts depending on grain size.

Intensity (arb. unit)

Table I.

T. ASADA et al.

400
200
Sample A

0
10

20

30

40

50

60

70

80

90

2θ (deg)
Fig. 4. XRD pattern of Cr2 O3 /CeO2 == STO bilayered samples A and B.
Open circles indicate peaks from STO (100). Sample A did not show any
diﬀraction peaks from Cr2 O3 . Sample B revealed a trigonal c-axis Cr2 O3
growth.

and B). Cr2 O3 ﬁlms were grown at Ts ¼ 550  C, at which ahighly oriented ﬁlm growth was obtained on R-Al2 O3 . No
Bragg diﬀraction peak from Cr2 O3 was observed forsample A. In contrast, sample B revealed a c-axis oriented
growth and the crystal relationship was Cr2 O3 (0001)=
CeO2 (100) == STO(100).
The lattice constant of Cr2 O3 is almost the same as that of
Al2 O3 . The lattice matching between CeO2 and Al2 O3 hasbeen studied.5,6) The lattice mismatchings in CeO2 (100) vs
Al2 O3 (112 0) and CeO2 (100) vs Al2 O3 (0001) were approximately 24 and 8%. Namely, the mismatching of Cr2 O3(0001)/CeO2 (100) is smaller than that of Cr2 O3 (112 0)/
CeO2 (100). Therefore, we consider that a c-axis-oriented
growth was obtained in sample B. The results of XRD
analysis indicated that Cr2 O3 grew on CeO2 (100) ==STO(100) with a c-axis orientation.
The Ra of the CeO2 layers was ﬁrst investigated using a
DFM image with a size of 1  1 mm2 . The Ra values of the
CeO2 buﬀer in samples A and B were 4.2 and 0.41 nm,
respectively.
The DFM surface images of samples A and B are shownin Figs. 5(a) and 5(b). In Fig. 5(a) showing sample A
deposited on the rough CeO2 buﬀer with Ra ¼ 4:2 nm,
various sizes of grains and deep valleys of grain boundaries
were observed and Ra was 17.0 nm. As shown in Fig. 5(b),smooth Cr2 O3 surfaces with Ra of 4.2 nm were observed onsmooth CeO2 with Ra ¼ 0:41 nm.
The surfaces of sample A were not smooth and no Bragg
diﬀraction peak was observed. In sample B, valleys disappeared and grains with clear facets appeared.
It was thought that the smooth part of sample B corresponded to the c-planes of Cr2 O3 . The crystal smoothnessof the CeO2 layer strongly aﬀected Cr2 O3 growth as animportant factor. The crystal quality of Cr2 O3 ﬁlm will be
improved more by improving the crystalline alignment ofCeO2 layers. From the obtained results on the crystalline
structure, Cr2 O3 ﬁlms are expected to be applied to devices,for example, as a magnetic control gate layer for novel oxide
electronic functional devices fabricated on STO substrates.

Jpn. J. Appl. Phys., Vol. 47, No. 1 (2008)

T. ASADA et al.

(a)

Ra=17.0nm(b)Ra=4.2nm

0nm

106.2nm

0nm

39.4nmCr2 O3 was aﬀected by the steps and/or width of terraces in
Al2 O3 substrates. An adequate substrate temperature Ts ¼550  C resulted in a highly oriented Cr2 O3 growth and theﬁlm was composed of coalesced large rectangular grains.
The STO substrate with CeO2 buﬀer showed the growth of
Cr2 O3 with a c-axis orientation. It was also conﬁrmed thatsuperior Cr2 O3 growth required smooth CeO2 buﬀers withsmall Ra . In conclusion, the Cr2 O3 ﬁlms are applicable tonovel perovskite oxide crystals fabricated on cubic STO
substrates.

Fig. 5. DFM surface images (1  1 mm2 ) of Cr2 O3 ﬁlms, (a) sample Aand (b) sample B. Sample A was deposited on the CeO2 buﬀer layer withRa ¼ 4:2 nm. Sample B was deposited on the comparatively smooth
CeO2 buﬀer layer with Ra ¼ 0:41 nm.

4.

Conclusions

The ME Cr2 O3 ﬁlms were prepared by sputtering on R-cut
sapphire and STO with CeO2 buﬀers. The grain shape of5491) K. Siratori, K. Kohn, and E. Kita: Acta Phys. Pol. A 81 (1992) 431.
2) P. J. Brown, J. B. Forsyth, and F. Tasset: Solid State Sci. 7 (2005) 682.
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Schmid: Ferroelectrics 162 (1994) 141.
4) M. Foebig: J. Phys. D 38 (2005) 123.
5) M. Shirakawa, M. Miura, T. Ohazama, Y. Shingai, A. Saito, M.
Mukaida, and S. Ohshima: Physica C 412–414 (2004) 1277.
6) J. C. Nie, H. Yamasaki, Y. Nakagawa, K. Develos-Bagarinao, M.
Murugesan, H. Obara, and Y. Mawatari: J. Cryst. Growth 284 (2005)
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7) H. Y. Lee, S. I. Kim, Y. P. Hong, Y. C. Lee, Y. H. Park, and K. H. Ko:
Surf. Coat. Technol. 173 (2003) 224.

Mol. Cryst. Liq. Cryst., Vol. 472, pp. 209=[599]–216=[606], 2007
Copyright # Taylor & Francis Group, LLC
ISSN: 1542-1406 print=1563-5287 online
DOI: 10.1080/15421400701547359

Growth of Vertically Aligned Carbon Nanotubes
Depending on Thickness of Catalyst Films by
Plasma-Enhanced Chemical Vapor Deposition
Hiroki Okuyama
Nobuyuki Iwata
Hiroshi Yamamoto
College of Science and Technology, Nihon University,
Funabashi-shi, Chiba, Japan

We attempted to control the diameter and density of vertically aligned carbon
nanotubes (VACNTs) using plasma-enhanced chemical vapor deposition. The
VACNT diameter was reduced by decreasing the Ni catalyst thickness. Introducing the Mo layer between Ni catalyst and quartz substrate was effective for thin
VACNT growth from the thin Ni catalyst film. This result demonstrated that themetal films had to be highly electric-conductive for inducing plasma on the Ni
catalyst to grow VACNTs. The maximum density of VACNTs was obtained
through the use of optimum plasma pretreatment time.
Keywords: density control; diameter control; Ni catalyst; plasma-enhanced chemical
vapor deposition; vertically aligned carbon nanotubes

1. INTRODUCTION
Carbon nanotubes (CNTs) have been studied intensively because of
their many characteristics such as high electric conductivity [1], highpermissible current density [2], and high mechanical strength in spiteof their high flexibility [3]. Furthermore, the growth direction of CNTs
is controllable by applying an electric field during chemical vapor
deposition (CVD) growth [4]. These characteristics are beneficial to
The authors are grateful for the financial support of a Grant-in-Aid from the Futaba
Foundation. This work was supported by the Ministry of Education, Culture, Sports,
Science and Technology through a Scientific Grant-in-Aid (No. 15360173).
Address correspondence to Hiroki Okuyama, College of Science and Technology,
Nihon University, 7-24-1-401, Narashinodai, Funabashi-shi, Chiba 274-8501, Japan.
E-mail: okuyama@yamanoya.ecs.cst.nihon-u.ac.jp
209=[599]

210=[600]

H. Okuyama et al.

prepare various nanostructures for use in scanning probe microscopy(SPM) tips [5–7], field emitters [8–10], wire of ultra large-scale integrated circuits [11], and so on. The required properties for CNTs aredifferent for each purpose. Requirements pertain for CNTs regarding
the diameter and density: high density of thin CNTs is necessary for
the wire to satisfy high permissible current density and low resistance[11]. In the case of field emitters, low density of thin CNTs is preferred
to enhance the electric field at CNT tips [12]. An individual and a thin
CNT are necessary for application in an SPM tip.
In this article, we report vertically aligned (VA) CNT growth using
plasma-enhanced (PE) CVD. We attempted to control the diameterand density of VACNTs intended to support electronics applications.
Diameter control was investigated using various thicknesses of
Ni catalyst and introducing a Mo underlayer. Density control was
performed by varying the plasma pretreatment time.

2. EXPERIMENTAL
Patterned metals of Ni catalysts and Mo underlayers were depositedon quartz substrates (5  5 mm2, 0.5 mm thickness) through a metalmask (30 mm thickness, 200 mm pattern width, 100 mm gap) using radio
frequency magnetron sputtering. The chamber background pressure
was 3.0  10 3 Pa, and the pressure during deposition was fixed at
2.0 Pa by feeding Ar gas. The growth conditions were 120C substratetemperature and 50 W sputtering power for Mo underlayers, and 70Cand 20 W, respectively, for Ni catalysts. The films’ deposition rates
were 10 nm=min for Mo and 3.5 nm=min for Ni.
The CNTs were grown using DC PECVD method. The reactor
tube of the CVD equipment was heated to 500C. The backgroundpressure was 0.1 Pa. The flow of H2:Ar (50:50 ccm) was introduced
into the reactor tube during heating. Plasma pretreatment was performed by inducing DC glow discharge from immediately before the
CNT growth. The DC glow discharge was induced by applying a bias
voltage of
250 V between the substrate holder and grounded
anodes separated from the substrates by approximately 5 mm. The
metal pattern was connected electrically to the cathode of substrateholder by applying Ag paste. The CNT growth was performed byadding 5 ccm of ethylene (C2H4) gas. The total pressure was maintained at 2 kPa.
The obtained CNTs were observed using a field emission scanningelectron microscope (SEM; S-4500; Hitachi Ltd.) operated at 15 kV anda field emission transmission electron microscope (TEM; HF-2000;
Hitachi Ltd.) operated at 200 kV.

Growth of Vertically Aligned Carbon Nanotubes

211=[601]

3. RESULTS AND DISCUSSION
Figure 1 shows the effect of Ag paste, which electrically connects
the Ni=Mo patterns to the cathodes. The respective thicknesses of Ni
catalyst and Mo underlayer were 3.5 nm and 10 nm. The specimen of
Figure 1(a) was prepared by putting the substrate on the cathodewithout applying Ag paste. The specimen shown in Figure 1(b) was
prepared by electrically connecting the Ni=Mo pattern to the cathodeby applying Ag paste. Some particles were observed on the patterns,
which were not connected electrically to the cathode. Figure 1(b)FIGURE 1 SEM images of the surface of the Ni=Mo pattern after CVD. (a)The Ni=Mo pattern was not connected electrically to the cathode. (b) The
Ni=Mo pattern was connected electrically to the cathode using Ag paste.

212=[602]

H. Okuyama et al.

shows that VACNTs were prepared on the patterns by applying Ag
paste. Figure 2 shows the VACNTs as observed using TEM. The
hollow center and graphitic sidewalls were visible. In this case, thenumber of walls was 25, and the CNT internal and external diameters
were ca. 4.6 nm and ca. 26 nm, respectively.
In the case of PECVD method, a cathode dark space is used mainly for
preparing CNTs [13]. The cathode dark space contains many positive
ions and radicals. The active species reduce the activation energy for acatalytic reaction. In addition, the voltage drop takes place mostly inthe cathode dark space because of its high impedance. The positive ionsare accelerated using an electric field; some of them bombard the cathodeof catalyst films. These impacts generate catalyst particles from continuous catalyst films [14–16]. The CNTs are aligned using an applied strong
electric field. The thickness of the cathode dark space was estimatedusing the Child-Langmuir law [17]. The estimated thickness of the cathode dark space was 0.21 mm. Therefore, the DC glow discharge was not
induced on a quartz substrate with 0.5 mm thickness although the

FIGURE 2 Typical TEM image of a VACNT.

Growth of Vertically Aligned Carbon Nanotubes

213=[603]

substrate was placed on the cathode. The catalysts should be connected
electrically to the cathode to induce plasma on the catalysts.
Figure 3 depicts SEM images of VACNTs with various Ni thickness.
The Ni thickness, the pattern resistance from the Ag paste, andthe diameter and the density of the grown VACNTs are listed inTable 1. The pretreatment was maintained for 1 min. The VACNT
diameter was decreased by reducing the Ni catalyst thickness.
Figure 3(a) shows that no CNTs grew from 3.5 nm of Ni thickness.
Although the Ni thickness was equal to that shown in Figure 3(a),the smallest VACNT diameter was obtained by introducing the Mo
underlayer, as shown in Figure 3(d).
The VACNT diameter was reduced by reducing the Ni catalyst
thickness. The continuous Ni catalyst films were broken into fine particles using thermal and plasma pretreatment. The CNT diameter isdetermined by the catalyst particle diameter. In addition, the catalyst
particle diameter depends on the initial thickness of the catalyst films[18,19]. However, VACNTs did not grow when the catalyst films were
too thin, i.e., 3.5 nm. The pattern conductivity was reduced because

FIGURE 3 SEM images of VACNTs grown from various Ni catalyst thickness. The Ni catalyst film thicknesses were (a) 3.5 nm, (b) 35 nm, (c) 350 nm,
and (d) 3.5 nm with a 10 nm Mo underlayer.

214=[604]

H. Okuyama et al.

TABLE 1 Summary of the Obtained VACNTs with Various Ni Catalyst
Thicknesses

Thickness of Ni [nm]Pattern resistance from Ag paste [X]VACNT diameter [nm]VACNT density [mm 2]

(a)

(b)

(c)

(d)

3.5
> 60 M



35
360 k
40
100

350
12 k
40–15030

3.5 with 10 nm Mo
300 k
30
90

the Ni catalysts formed fine particles. The resistance of the 3.5 nm Ni
patterns was greater than 60 MX. Therefore, the plasma was not
induced on the 3.5 nm Ni patterns during CNT growth because of theirhigh resistance. By introducing the Mo underlayer, VACNTs withsmallest diameters were obtained from 3.5 nm Ni patterns. The densities of VACNTs grown from 3.5 nm Ni patterns were slightly less thanthose of 35 nm Ni. Although the Ni catalysts were thin, plasma was
induced on Ni catalysts because of the conductive Mo underlayer

FIGURE 4 SEM images of the VACNTs at different plasma pretreatment
times of (a) 0.5 min, (b) 1 min, and (c) 5 min.

Growth of Vertically Aligned Carbon Nanotubes

215=[605]

and Ag paste. Therefore, introducing an Mo underlayer is effective
for preparing VACNTs with small diameter and low density.
Figure 4 shows SEM images of samples with different plasma
pretreatment times. The Ni catalyst and Mo underlayer were 3.5 nm
and 10 nm thick, respectively. The pretreatment times were (a) 0.5 min,(b) 1 min, and (c) 5 min. In fact, VACNTs with ca. 30 nm of diameter grewfrom the films, which had been pretreated for 0.5 and 1 min. For pretreatment time of 5 min, poorly aligned CNTs with 20–35 nm diameter and
80–450 nm of length grew sparsely. The respective densities of the
VACNTs are shown in Figure 5: they were (a) 50 mm 2, (b) 90 mm 2,and (c) 20 mm 2.
The CNT density was increased by extending the pretreatment
time from 0.5 min to 1 min. The plasma induces catalyst films toform fine particles, as discussed above. Many catalyst particles were
produced during the extended pretreatment time. However, theCNT density was decreased by extending the pretreatment time to
5 min. The Ni catalyst and Mo underlayer were removed throughthe use of longer pretreatment times. Sparse CNTs were preparedbecause a small amount of Ni catalyst remained at the pattern.
Short and the poorly aligned CNTs were prepared because theresistance of the Mo underlayer was increased as a result of the
reduced thickness.

FIGURE 5 Density of the VACNTs as a function of plasma pretreatment
time.

216=[606]

H. Okuyama et al.

4. CONCLUSIONS
We attempted to control the density and the VACNT diameter using
DC PECVD. It was important for VACNT growth that the DC plasma
be induced onto the Ni catalyst. The VACNT diameter was reduced bydecreasing the Ni catalyst thickness. Introducing the Mo underlayer
was effective for thin VACNT growth from the thin Ni catalyst film.
The VACNT density was maximized by plasma pretreatment time of
1 min. Longer or shorter pretreatment time than 1 min yielded a lower
VACNT density.

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2767.

Japanese Journal of Applied Physics
Vol. 47, No. 1, 2008, pp. 700–702#2008 The Japan Society of Applied Physics

Fabrication of Carbon Nanoﬁber Emitters for Excitation of Organic Phosphor Thin Films
Shinji I DE, Masato Y OSHIKUNI, Nobuyuki I WATA, and Hiroshi YAMAMOTO
College of Science and Technology, Nihon University, 7-24-1 Narashinodai, Funabashi, Chiba 274-8501, Japan(Received July 12, 2007; revised September 27, 2007; accepted October 11, 2007; published online January 22, 2008)

For applications to a novel low-voltage acceleration ﬁeld emission display, carbon nanoﬁber (CNF) emitters were fabricatedusing a screen printing method. The two types of CNF of diﬀerent lengths, normal-length CNFs, and cut CNFs were
examined. A current density of 1.4 mA/cm2 was achieved at 10 V/mm in CNF emitters prepared from the paste composed ofnormal-length CNFs (1 wt %), glass paste (6 wt %), and gold nanometal ink (30 wt %). The ﬁeld enhancement factor  in the
Fowler–Nordheim equation was ca. 5800. CNF emitters without gold nanometal ink showed low performance. Emitter
performance was determined on the basis of the results of scanning electron microscopy (SEM) and ﬁeld emission
characteristics. [DOI: 10.1143/JJAP.47.700]
KEYWORDS: carbon nanofiber, field emission display, screen printing, CNF emitter, emission current density

1.

Introduction

Recently, organic light-emitting devices (OLEDs) andﬁeld emission displays (FEDs) have been investigatedaggressively as next generation displays. However, eachdisplay has its own problems: limited lifetime for OLEDs orhigh acceleration voltage for FEDs. We have proposed anovel low-voltage acceleration ﬁeld emission display (LVFED) composed of simply fabricated carbon nanoﬁber(CNF) emitters and organic electroluminescent (EL) thinﬁlms as shown in Fig. 1.1) The speciﬁc structural features ofhigh luminescence eﬃciency and long lifetime are due to thediﬀerence in reﬂectance between vacuum and emissive
layers and/or the absence of cathode layers. In our previous
study, a stacking structure of tris(8-hydroxyquinolate)aluminum (Alq3 )/N,N0 -di-[(1-naphthyl)-N,N0 -diphenyl]-(1,10 biphenyl)-4,40 -diamine (-NPD)/CuPc on indium tin oxide(ITO) substrates was irradiated by thermally activatedelectrons with an emission current density of ca. 1 mA/
cm2 .2,3) However, the observed luminescence disappeared
quickly in 10 s. It was concluded that the deterioration of
luminescence was caused by the accumulation of electrons
at the interfaces of Alq3 /-NPD because of an insuﬃcientelectric ﬁeld applied in the organic stacking layers. This
problem will be overcome by placing electron emitters close
to luminescent layers. For example, with 50 mm and 500 V
between the cathode and the anode, an electric ﬁeld of 10 V/
mm is applied.4)
We aimed to simply fabricate CNF emitters applicable tothe above-proposed LV-FED. Two kinds of CNF were
investigated. The purpose of comparing a normal-lengthﬁber with a shortened ﬁber is to study the dependence of
CNF emitter performance on length and/or shape. Normallength tangled and short-length CNFs easily came apart in
the paste. Therefore, two types of emitter site are expected.
In one type of site, the side of tangled CNFs are notsurrounded by paste and not exposed on paste surfaces. In
other type of emitter site, the ends of CNFs protrude frompaste surfaces. The latter type is expected for cut-CNFs
rather than normal-length CNFs. Moreover, binder materials
and conducting materials in CNF paste were studied toobtain excellent ﬁeld emission properties.5) Glass paste wasused as the binder material to increase adhesive strength andgold nanometal ink as the conducting material to increaseelectric conductivity between CNFs and the substrate. The

Fig. 1. Schematic conﬁguration of the novel LV-FED composed ofsimply fabricated CNF emitters and EL thin ﬁlms.

CNF emitter fabrication technique used was a screen
printing method. The method using CNF paste has beenrecognized as a promising technique for FED fabrication,
which is applicable to large-area displays or amenable tomass production.6) We intended here to prepare CNF
emitters that can achieve an emission current density higher
than 1 mA/cm2 .
2.

Experimental Procedure

CNF emitters were fabricated on an ITO substrate using ascreen printing method. The CNFs used were multiwall
carbon nanoﬁbers (MWCNFs) produced by Showa Denko.
The CNFs were massively synthesized from a mixed
solution of ferrocene, thiophene, and benzene by thermal
chemical vapor deposition at ca. 1000  C. The obtained
CNFs were baked above 2500  C for 30 min to eliminate
amorphous carbon and catalysts.7) The CNFs used were
multiwalled with a diameter of about 0.1 mm (hereafter,
normal-length CNFs). The shorted CNFs (hereafter, cutCNFs), which were prepared from the normal-length CNFs
by a fragmentation process, were also used. The scanningelectron microscopy (SEM) images of normal-length CNFs
and cut-CNFs are shown in Fig. 2. The normal-length CNFs
were tangled and the cut-CNFs were straight. The lengths ofnormal-length CNFs and cut-CNFs were ca. 20 mm and lessthan 20 mm, respectively. The insulating glass paste (Asahi
Glass AP5346) and conducting gold nanometal ink (ULVAC
Materials Au1T) were adopted as CNF binding materials.
700Jpn. J. Appl. Phys., Vol. 47, No. 1 (2008)(a)S. IDE et al.

(a)

(b)

(b)

3µm

1µm

1µm

3µm

(c)

(d)

Fig. 2. SEM images of (a) normal-length and (b) cut-CNFs.

Table I. Preparation conditions of CNF emitters (unit: wt %).
C1G6Au0

C1G6Au30

Normal

Cut

Cut

1

1

1

1

Gold nanometal ink

0

30

0

30

-Terpineol

93

63

93

63

Glass paste

6

6

6

6

CNF

3µm

Fig. 3. SEM images of CNF emitters (a) N1G6Au0, (b) N1G6Au30, (c)C1G6Au0, and (d) C1G6Au30.

N1G6Au0
N1G6Au30
C1G6Au0
C1G6Au30

1400

In Table I, the preparation conditions of CNF emitters are
summarized. The normal-length or cut-CNFs (1 wt %) weredispersed in -terpineol and sonicated for 24 h. The solution
with glass paste (6 wt %) and gold nanometal ink (0 or
30 wt %) was mixed and stirred for 30 min. CNFs were
presonicated in gold nanometal ink for 1 h. The obtained
solution was CNF paste. Regarding the CNF emitter sample
name, the ﬁrst initial N or C means normal-length or cutCNF, respectively. The second initial G means glass paste
and the third initial Au means gold nanometal ink. The
number appearing with the name is the wt % of eachcomponent.
ITO substrates were cleaned by rinsing in acetone andethanol, and then dried in blown compressed nitrogen gas.
The substrate size was 10  20 mm2 . Squeezed CNF paste
was dried at 150  C for 10 min and baked at 510  C for 5 min
to remove organic solution in the paste.
The CNF emitter surface morphology was observed by
SEM. Emission current was measured in 3  106 Torr atroom temperature. The distance between an anode copper
plate and the cathode CNF emitter was ca. 50 mm. The
emission sites should be distributed uniformly. There were,
however, inequalities on screen-printed surfaces. Therefore,
the gap length of electrodes is a critical parameter for emitter
performance. In this work, a gap of ca. 50 mm wasinvestigated because the length of CNFs and the roughnessof paste surfaces were typically in a few micrometer order.
The emission area was 5  5 mm2 .
3.

3µm

1200

–2.5
–3.0

1000

–3.5
2

N1G6Au30

NormalLog(I/V )N1G6Au0

Current Density [ µ A/cm2 ]Sample

800
600

–4.0
–4.5
–5.0
–5.5

400

–6.0
0.002

0.003

200

0.004

1/V

0
0

2

4

6

Electric field [V/µ m]

810Fig. 4. Emission current density vs. applied electric ﬁeld. The solid lines
are theoretical ﬁtting curves obtained using the Fowler–Nordheim
equation with each evaluated  value. The inset shows logðI=V 2 Þ vs 1=V.

particles. N1G6Au30 showed a cluster of glass and gold but
CNFs were not clearly observed. In C1G6Au30, thediameter and the length of CNFs were ca. 200 nm and over
6 mm, respectively.
The bent normal-length CNFs protruded from the glass
paste surface, whereas cut-CNFs were buried in the glass
paste. Thus, the considered suﬃcient ﬁeld emission was
expected in the normal-length CNF emitter. The SEM
images suggest that glass paste is a suitable binder material
and normal-length CNFs are suitable for emitter materials.
Figure 4 shows the obtained emission current density as afunction of applied electric ﬁeld. N1G6Au30 revealed amuch higher emission current density than other samples. A
current density of 1.4 mA/cm2 was achieved at 10 V/mm.
The emission currents obtained from C1G6Au0, N1G6Au0,
and C1G6Au30 were smaller, 29, 58, and 230 mA/cm2 , at
10 V/mm, respectively. It was noted that the current density
obtained from normal-length CNF emitter N1G6Au30 was
about 6 times as large as that from cut-CNF emitter
C1G6Au30. Deﬁning the turn-on electric ﬁeld Eto as the

Results and Discussion

The SEM images of N1G6Au0 and C1G6Au0 are shownin Figs. 3(a) and 3(b). Furthermore, those of N1G6Au30 and
C1G6Au30 with gold nanometal ink are shown in Figs. 3(c)and 3(d). In N1G6Au0, CNFs and the glass paste wereintermingled and CNFs stood against glass particles. The
glass particles were round. The diameter and the length ofprotruded CNFs were 20 – 200 nm and 5 mm, respectively.
The glass paste held CNFs and increased adhesive strength.
In C1G6Au0, the emitter CNF bunches were ﬁlled with glass
701

Jpn. J. Appl. Phys., Vol. 47, No. 1 (2008)

S. IDE et al.

value at which a current density of 1 mA/cm2 was observed,
Eto was 6 V/mm in N1G6Au30. Other emitters showed highEto , 8 V/mm.
The emitters containing gold nanometal ink showed
higher current densities. Gold nanometal ink was suitablefor increasing emission current density because electric
conductivity between CNFs and substrates increased.
The insert in Fig. 4 shows the Fowler–Nordheim8) plots.
The ﬁeld enhancement factor  was evaluated by ﬁtting theexperimental data, assuming that the work function of CNF
is 5 eV.9) The solid curves and lines are ﬁtting results
obtained using the  value of each emitter. The  value of
N1G6Au30 was larger, ca. 5800, than that of other samples.
From the observation of SEM images, it was found that
normal-length CNFs leaned against each other and cut-CNFs
were lying on paste surfaces. The applied electric ﬁeld iseasily concentrated at the ends of the leaning CNFs, whichfunctioned as eﬀective emitter sites. Thus, an inferior
emitter performance was obtained in the case of using thecut-CNFs. Gold nanometal ink included in the glass binder is
expected to induce electric activity for CNFs noncontactingwith substrates. On the other hand, small air spaces intangled CNFs are not ﬁlled by the glass binder. Even in thiscase, very small gold nanometal ink particles with adiameter of ca. 7 nm adhere to CNFs and might improvethe electric contact between CNFs. Thus, the total electric
conductivity of CNFs on substrates increases and the densityof emission sites increases. As a result, the sample with gold
nanometal ink showed better emission performance.
Conclusively it was demonstrated that as emitter materials, normal-length CNFs and gold nanometal ink are suitablefor obtaining a high performance of a CNF emitter preparedby a screen printing method.

4.

Conclusions

CNF paste mixed with glass paste and gold nanometal ink
was adopted to prepare CNF emitters with high performance. The glass paste was used to increase adhesive strengthwith substrates. The gold nanometal ink was used to increaseelectric conductivity between CNFs and substrates. Also, thetwo types of CNF with diﬀerent lengths, normal-length
CNFs and cut-CNFs were examined. The highest emission
current density obtained was 1.4 mA/cm2 at 10 V/mm in theemitter composed of 1 wt % normal-length CNFs and
30 wt % gold nanometal ink. The ﬁeld enhancement factor
 in the Fowler–Nordheim equation was ca. 5800. Since
current density above 1 mA/cm2 was successfully obtained
from the simply prepared CNF emitters, they will be applied
to novel LV-FEDs.

1) N. Iwata, Y. Hata, M. Yoshikuni, and H. Ymamoto: Mater. Sci. Eng. C
27 (2007) 1174.
2) K. Onozuka, N. Iwata, and H. Yamamoto: Int. J. Mod. Phys. B 16
(2002) 4445.
3) Y. Hata, T. Ichinose, N. Iwata, and H. Yamamoto: Trans. Mater. Res.
Soc. Jpn. 30 (2005) 175.
4) B. J. Chen, W. Y. Lai, Z. Q. Gao, C. S. Lee, S. T. Lee, and W. A.
Gambling: Appl. Phys. Lett. 75 (1999) 4010.
5) H. Y. Shin, W. S. Chung, K. H. Kim, Y. R. Cho, and B. C. Shin: J. Vac.
Sci. Technol. B 23 (2005) 2369.
6) H. Machida, S. Honda, S. Ohkura, K. Oura, H. Inakura, and M.
Katayama: Jpn. J. Appl. Phys. 45 (2006) 1044.
7) M. Yoshikuni, S. Ide, N. Iwata, and H. Yamamoto: Trans. Mater. Res.
Soc. Jpn. 32 (2006) 329.
8) R. H. Fowler and L. W. Nordheim: Proc. R. Soc. London, Ser. A 119(1928) 173.
9) O. Gröning, O. M. Küttel, C. Emmenegger, P. Gröning, and L.
Schlapbach: J. Vac. Sci. Technol. B 18 (2000) 665.

702

Physica C 463–465 (2007) 1369–1371
www.elsevier.com/locate/physc

Development of high capacity Stirling type pulse tube cryocooler
J. Imura a,*, S. Shinoki a, T. Sato a, N. Iwata a, H. Yamamoto a, K. Yasohama a, Y. Ohashi b,
H. Nomachi b, N. Okumura b, S. Nagaya c, T. Tamada c, N. Hirano caCollege of Science and Technology, Nihon University, 7-24-1, Narashinodai, Funabashi-shi, Chiba 274-8501, Japanb
Aisin Seiki Co., Ltd., 2-1, Asahi-machi, Kariya, Aichi 448-8650, Japanc
Chubu Electric Power Co., Inc., 1, Toshin-cho, Higashi-ku, Nagoya-shi, Aichi 261-8680, Japan
Accepted 5 April 2007
Available online 5 June 2007

Abstract
We have been developing a Stirling type pulse tube cryocooler, aiming for a cooling capacity of 200 W at 80 K for a superconducting
magnetic energy storage system. In this work, we adopted stainless steel meshes for the regenerator of the cryocooler, and studied theinﬂuences of the mesh number on the cooling capacity. The prepared mesh numbers were #150, 200, 250, 350 and 400. Using #250 mesh,and at a frequency of 45 Hz and power consumption of 3.1 kW, the achievable lowest temperature and cooling capacity at 80 K was
46.2 K and 123 W, respectively. Furthermore, in order to optimize the performance, some regenerators were made by stacking several
kinds of meshes with diﬀerent stacking orders. Using these regenerators, we have obtained a high cooling capacity of 169 W at 80 K withpower consumption of 4 kW.
Ó 2007 Elsevier B.V. All rights reserved.
PACS: 07.20.Mc
Keywords: Pulse tube cryocooler; Regenerator; Cooling capacity

1. Introduction
Usually there is a method of using liquid gases such as
liquid helium or liquid nitrogen, for cooling of superconducting apparatus. However, considering running cost
and troubles of handling liquid gases, usage of a cryocooler
is much preferable. Typical cryocoolers are Stirling [1]cryocooler, Giﬀord-McMahon (GM) cryocooler [2] andpulse tube cryocooler [3–5]. The pulse tube cryocooler isdivided into two types of Stirling and GM according to
the method of pressure wave generations. Stirling andGM cryocoolers have a displacer for phase shifting at
low temperature part. This moving displacer induces undesirable vibrations in these cryocoolers. However, the pulse
tube cryocooler does not have such a displacer, so we can
expect it has low vibration and high reliability.
*

Corresponding author. Tel./fax: +81 47 469 5457.
E-mail address: junnosuke_imura@yahoo.co.jp (J. Imura).

0921-4534/$ - see front matter Ó 2007 Elsevier B.V. All rights reserved.
doi:10.1016/j.physc.2007.04.309

We have been developing a Stirling type pulse tube cryocooler, aiming at a cooling capacity of 200 W at 80 K for a
superconducting magnetic energy storage system. In orderto achieve the purpose, we investigated the characteristics
of a regenerator that is one of the most important components of a cryocooler.

2. Experimental
Fig. 1 shows a schematic drawing of Stirling type pulse
tube cryocooler used in this experiment. As shown in
Fig. 1, the cryocooler consists of a pressure wave generator(PWG) [6], heat exchangers, a regenerator, a cold head, apulse tube, and a phase shifting system made up with aninertance tube [7] and a buﬀer tank.
The linear compressor of CFIC Inc. was used as a PWG.
The pistons of the linear compressor were diameter of

1370

J. Imura et al. / Physica C 463–465 (2007) 1369–1371

Fig. 1. Schematics of Stirling type pulse tube cryocooler.

115 mm and length of 50 mm. Displacement of the piston
was 20 mm of peak to peak (P-P) at the maximum.
Pulse tube was diameter of 40 mm and length of
200 mm.
The regenerator was diameter of 100 mm and length of
60 or 80 mm. We have prepared diﬀerent meshes for theregenerator; these numbers are #150, 200, 250, 350 and
400 mesh. These meshes had almost the same porosity.
An inertance tube, which is suitable for high frequency
operation of 50 Hz around, was adopted as a phase shifting
system. The external diameter of inertance tube was
12.7 mm, thickness 1 mm, and length 1.2 m. The capacityof the buﬀer tank was 3.78 l.
Working ﬂuid was helium gas and its charged pressure is
2 MPa.
Temperature at the cold head and regenerator was measured by using Pt–Co and Pt resistance thermometers,
respectively.
Pressure amplitudes of helium gas were measured at theoutput of the PWG (P1) and at the inlet of the inertancetube (P2).
3. Results and discussion
At ﬁrst, we show the results for the regenerators madeby stacking one kind of meshes. Experiments were carriedout at driving frequency of 45 Hz and piston displacementsof 14.5 mm P-P. The lowest temperature of the cold headFig. 2. The lowest temperature of the cold head vs. mesh number.

Fig. 3. Cooling capacity at 80 K vs. mesh number.

Fig. 4. Pressure ratio vs. mesh number.

and cooling capacity at 80 K vs. mesh number are shownin Figs. 2 and 3, respectively. With #250 mesh, the lower
temperature of cold head and highest cooling capacity at
80 K were obtained in this series. Furthermore, as for thelength of regenerator, 80 mm is better than 60 mm. With
#250 mesh and regenerator length of 80 mm, the lowest
temperature of the cold head and the highest cooling
capacity at 80 K was 46.2 K and 123 W, respectively. In
this case, power consumption was 3.1 kW. Fig. 4 showspressure ratio P2/P1 vs. mesh size. Pressure wave is transmitted in the direction of the pulse tube from the PWG.
When P2 is smaller than P1, pressure loss takes place inthe cryocooler. That is, the decrease of the pressure ratiomeans the increase of pressure loss. This loss causes temperature rise and reduction of cooling power. We thinkthat the regenerator causes the pressure loss.
The temperature distribution of a regenerator is shownin Fig. 5. The origin of the horizontal axis of the ﬁgure corresponds to the position of the higher temperature side ofregenerator, and 80 mm its lower temperature side. For#150, #200 and #250 meshes, a gradient of the temperature was almost constant as a function of a position. But
for #350 and #400, temperature was almost constant up
to the middle of the regenerator, and then fell down.
The following results are for the regenerators made bystacking mixed meshes. Ideally, the capability of the regen-J. Imura et al. / Physica C 463–465 (2007) 1369–1371

1371

number mesh in the lower temperature region. The performance of cooling capacity with 3 kinds of meshes is shownin Fig. 6. The result for #250 mesh regenerator is also
shown for comparison. By stacking mixed meshes thehigher cooling capacity was obtained rather than by stacking one kind of meshes. The highest performance wasobtained in the regenerator with #200, 250 and 400 meshes;
the cooling capacity of 148 W at 80 K. Furthermore, when
the condition was changed into driving frequency of 47 Hz
and the piston displacement of 16 mm P-P, the higher performance was obtained; the cooling capacity of 169 W at
80 K under power consumption of 4 kW.
Fig. 5. Temperature distribution of the regenerator.

4. Conclusion
In this work, the eﬀects of the mesh size on cooling performance have been studied. To optimize the performance,
several kinds of mesh number were stacked in the regenerator. Then the small number mesh was placed for the hightemperature region of the regenerator and the larger number mesh for the low temperature. With one of these regenerators we have obtained the cooling capacity of 169 W at
80 K under electric input power of 4 kW.
References

Fig. 6. Cooling capacity vs. temperature of the cold head.

erator constructed with the larger number mesh is higher.
But, pressure loss becomes large as increasing the meshes,
as shown in Fig. 4. Therefore, the small number mesh
was placed in the higher temperature region and the large[1] C.M. Hargreaves, Elsevier Science Publishers, 1991, p. 94.
[2] H.O. McMahon, Adv. Cryo. Eng 5 (1960) 354.
[3] W.E. Giﬀord, R.C. Longsworth, J. Eng. Ind 63 (1964) 264.
[4] W.E. Giﬀord, R.C. Longsworth, Adv. Cryo. Eng. 10 (1965) 69.
[5] K. Oshima, Y. Matsubara, T. Kubo, Cryo. Eng, Japan 3 (1968) 72.
[6] T.W. Bradshaw, J. Delderﬁeld, S.T. Werreu, G. Davey, Adv. Cryo.
Eng. 31 (1986) 801.
[7] S.W. Zhu, S.L. Zhou, N. Yoshimura, Y. Matsubara, Cryo 9 (1997)
269.

Physica C 463–465 (2007) 1005–1008
www.elsevier.com/locate/physc

Preparations of magnetoelectric thin ﬁlms for superconducting devices
N. Iwata *, T. Asada, K. Nagase, T. Yamada, H. Yamamoto
Department of Electronics and Computer Science, College of Science and Technology, Nihon University, 7-24-1 Narashinodai,
Funabashi-shi, Chiba 274-8501, Japan
Received 30 October 2006; accepted 9 February 2007
Available online 26 May 2007

Abstract
Representative magnetoelectric (ME) Cr2O3 thin ﬁlms were grown on R-cut sapphire and SrTiO3(1 0 0) substrate in order to integratewith high temperature superconductor oxides. The replacement of a base current and stabilization of IC spread for single ﬂux quantum
device is expected. Highly oriented Cr2O3 thin ﬁlm grew on sapphire at 550 C by oﬀ-axis DC–RF hybrid magnetron sputtering methodwith two-in. Cr metal target. The surface was composed of coalesced rectangular grains, the surface roughness Ra of which was 1.9 nm.
Estimated induced magnetic ﬁeld with electric ﬁeld applied of 0.24 MV/cm was 1.0 G around 40 K.
 2007 Elsevier B.V. All rights reserved.
PACS: 68.55. a; 81.15. z; 81.15.Cd
Keywords: Magnetoelectric eﬀect; Cr2O3; Single ﬂux quantum; Thin ﬁlm

1. Introduction
We pay attention to apply the magnetoelectric (ME)materials of single phase compounds to single ﬂux quantum (SFQ) device integrated by high critical temperature(high-Tc) superconductors [1–4]. The ME eﬀect is deﬁned
as the dielectric polarization of a material in an appliedmagnetic ﬁeld or an induced magnetization in an external
electric ﬁeld [1–4]. The performance of the SFQ device isfeatured as a high-speed (over 100 GHz) and low consumption (less than 100 nW/gate) following Si digital devices.
However, increasing integration degree, large amount ofbase current is required, which give rise to uncontrollable
operation due to accompanying magnetic ﬁeld. In addition,
to reduce the ﬂuctuation of critical current (IC) and to stabilize the device operation, working temperature is limited
less than 40 K.

*

Corresponding author. Tel./fax: +81 47 469 5457.
E-mail address: iwata@ecs.cst.nihon-u.ac.jp (N. Iwata).

0921-4534/$ - see front matter  2007 Elsevier B.V. All rights reserved.
doi:10.1016/j.physc.2007.02.036

As shown in the Fig. 1, we propose a novel element how
the ME material is embedded in a SFQ device. Since theME material is grown to generate a base current, the basecurrent is localized and then error operations could be
reduced. It is expected that a base current in a SFQ device
is replaced by the ME materials. The electric power loss
from a current source is excluded. Furthermore, considering critical current IC of the Josephson junction is modulated by the magnetic ﬁeld penetrating the junction area,
IC spread of the junctions could be reduced by the induced
magnetic ﬁeld from the ME ﬁlm. We expect that a largerscale circuit is achievable by the induced and localized magnetic ﬁeld from ME materials integrated with SFQ devices.
In this paper, we have focused on an antiferromagnetic
insulator Cr2O3 which is a representative oxide among ME
single phase compounds. Cr2O3 has the linear ME susceptibility of axx, ayy and azz and the others are zero [1–4].
Magnetic ﬁeld is induced along the direction of electric ﬁeld
applied. Since the ME eﬀect is derived from a magnetic
superexchange interaction, highly oriented ME ﬁlms areclaimed on SrTiO3 (STO), MgO, sapphire, etc., to integratewith high-Tc superconductors.

1006

N. Iwata et al. / Physica C 463–465 (2007) 1005–1008

Junction

5Fig. 1. Schematic image of a part of the SFQ device with ME material we
propose. Induced magnetic ﬁeld from the ME ﬁlm generates the shielding
current along the high-Tc superconductor. Shielding current could be abase current for operation of SFQ device.

3

10

Cr2O3(3306)

insulator

Cr2O3(2204)

ME

4

10

Cr2O3(1102)

YBCO

Intensity (arbitrary units)

10

2

10

1

10

0

10

20

30

40

50

60

70

80

90

2θ (degrees)

2. Experimental

Fig. 2a shows a XRD pattern of Cr2O3 ﬁlm grown on
R-cut sapphire substrate at 500 C and Fig. 2b is the magniﬁed pattern around Cr2O3(1 1 0 2) Bragg reﬂection at various substrate temperatures from 300 C to 600 C. Bragg
reﬂections from R-plane of Cr2O3 were clearly observed justbelow those of sapphire substrate in 2h. The peak indicatedby open-triangle was probably due to oxygen deﬁcient of
substrate. Any other peaks did not appear in XRD patterns.
The lattice spacing was 0.363 nm similar to that of substrate. As shown in Fig. 2b, peak intensity was the highest
at 500 C with shoulder at right hand side of the peak indicated by arrow. Peaks became decreasing in intensity andbroader in ﬁlms grown 300 C, 400 C and 600 C. Fig. 3
shows the full width at half maximum (FWHM) of rocking
curve for Cr2O3(1 1 0 2) as a function of substrate temper-

10

600ºC

8
6
4

500ºC

400ºC

2

300ºC
24.0

24.2

24.4

24.6

24.8

25.0

2θ (degrees)Fig. 2. XRD patterns of Cr2O3 ﬁlm grown at (a) 500 C on R-cut sapphireand (b) magniﬁed images around Cr2O3(1 1 0 2) at various substratetemperature. Open circles and triangle are come from the substrate. Bragg
reﬂections of Cr2O3 thin ﬁlm were clearly observed.

1.1

FWHM of rocking curve for Cr2O3(1102)

3. Results and discussion

Intensity (arbitrary units)

12

Cr2O3 ﬁlms were grown by oﬀ-axis DC–RF hybrid magnetron sputtering method with two-in. Cr metal target on
the substrates of R-cut sapphire and STO with CeO2 buﬀer
layer. Substrates surface of sapphire was etched by HF acidfor 10 s (HF:H2O = 1:7) and then annealed at 1000 C for
12 h in air in advance. For Cr2O3 growth substrate temperatures, DC current, RF power and sputtering gas were300–630 C, 0.04 A, 100 W and 0.26 Pa (Ar:O2 = 4:1),
respectively. CeO2 buﬀer layer was deposited by on-axis
RF magnetron sputtering method with pressed CeO2 powder at 380 C. The substrate was 0.5 wt.% Nb-doped
STO(1 0 0). A post-annealing was carried out at 1000 C
for 4 h in air prior to growth of Cr2O3.
The X-ray diﬀraction (XRD) 2h h scans from 5 to 90,using Rigaku RAD-C system with Cu Ka1 and Ka2 radiation generated with a tube power of 40 kV and 30 mA,
examined the grown ﬁlms. The diﬀraction area was always
same with divergence slit of 1/2 and sample size of
5 · 10 mm2. The surface morphology was observed in a
dynamic force mode by an atomic force microscopy(DFM, Seiko Instruments Inc. SII: NanoNavi Station).

1.0
0.9
0.8
0.7
0.6
0.5

300

350

400
450
500
550
Substrate Temperature ( C)

600

Fig. 3. The FWHM of rocking curve for Cr2O3(1 1 0 2) as a function ofsubstrate temperature. At 500 C the sharpest peak was obtained.

ature. At 500 C the rocking curve was 0.58 and the
sharpest.
Generally the FWHM of Bragg reﬂections in 2h h
examination depends on a ﬁlm thickness in particular lessthan 500 nm and a ﬂuctuation of a lattice spacing. Since

N. Iwata et al. / Physica C 463–465 (2007) 1005–1008

the thickness of the grown ﬁlm was constant approximately
190 nm in these experiments, the FWHM is a function of
only the ﬂuctuation of the Cr2O3(1 1 0 2) lattice spacing.
Considering the appearance of the shoulder at the righthand side of the Cr2O3(1 1 0 2) Bragg reﬂection for theﬁlm grown at 500 C, which was derived from the separation of the peak by Ka1 and Ka2 of diﬀerent X-ray wave
lengths, the ﬂuctuation of the lattice spacing was reduced
and then highly oriented ﬁlm was obtained. Although a lattice mismatch is +4.01% along the [1 1 2 0] direction and is+4.34% along the [ 1 1 0 1] direction, an expected compressive stress in-plane is seemed to be released form the resultof the same lattice spacing to the bulk value.
Figs. 4 and 5 show the surface images of the Cr2O3 ﬁlms
with the size of 2 lm · 2 lm and line proﬁles of thoseimages. In Table 1, grain sizes as a function of substrate(a) 400ºC, Ra=1.5 nm(b) 500ºC, Ra=2.3 nm

1007

Fig. 5. Line proﬁles related to the index of the surface images of Fig. 4.
Solid arrows indicate the diﬀerence in height at the grain boundary where
coalescence did not take place. Dotted arrows indicate the steps composedof multiples of 1/3 unit cell along the [1 1 0 2] direction inside one grain.

(b-L1)

Table 1Substrate temperature vs. grain size(a-L1)(b-L2)

0 nm

13.6 nm(c) 550ºC, Ra=1.9 nm

0 nm(d) 600ºC, Ra=4.2 nm(c-L2)(d-L1)

16.9 nm

0 nm33.2 nm(e) 630ºC, Ra=2.8 nm(e-L1)

0 nm

Grain size (nm · nm)

400
500
550
600

100 · 100–200
100 · 100–500
200 · 200–600
400 · 400

21.0 nm(c-L1)

0 nm

Sub. temp. (C)

24.3 nmFig. 4. Surface images with 2 · 2 lm2 of Cr2O3 thin ﬁlms grown at (a)400 C, (b) 500 C, (c) 550 C, (d) 600 C and (e) 630 C. Scale bars wereattached below each image. The ﬁlm grown at (c) 550 C showed thesmoothest surface, the roughness Ra of which was 1.9 nm. Rectangulargrains with the size of 200 · 200–600 nm2 coalesced throughout thesurface expect for the grains separated by deep grain boundaries.

temperature are summarized. Rectangular grains weremainly observed with the size of 100 · 100–200 nm2 at
400 C in Fig. 4a. The diﬀerence in height between grains
was approximately 3.5 nm indicated by solid arrows in
Fig. 5a-L1. One grain contained 2 or 3 steps at the surface,indicated by dotted arrows, with the height of 0.3–0.9 nm,
which was integer multiple of 1/3 unit cell along the[1 1 0 2] direction, 0.363 nm. At 500 C larger square
grains with the size of 100 · 100–500 nm2 were observed
and those grains were coalesced in Fig. 4b. In one grain
the steps of 1–3 integer multiples of 1/3 unit cell was also
observed as shown in Fig. 5b-L1 indicated by arrows.
The grains growing from diﬀerent nucleus sites, where coalescence did not occur, had a diﬀerence of 2–10 nm inheight as shown in Fig. 5b-L2 and sometimes deep grainboundaries were formed. The deepest depth of the boundaries was 34 nm. At 550 C as shown in Fig. 4c rectangulargrains became larger and longer along a particular direction and the smoothest surface was obtained with the surface roughness Ra of 1.9 nm. Almost all of the grains
coalesced except for the grains separated by deep grainboundaries as shown in Fig. 5c-L1 indicated by arrows.
The density of deep boundaries decreased compared withthat of Fig. 4b. As indicated by dotted arrows in Fig. 5cL1 and c-L2, steps of 1/3 unit cell were clearly observedin one grain. Film growth looked like 2-dimensional island.
At 600 C although the largest grains were obtained,bunching steps, the value of which was approximately
5 nm, were observed as shown in Figs. 4d and 5d-L1.

1008

N. Iwata et al. / Physica C 463–465 (2007) 1005–1008

4

Current Density (A/cm2)

10

2

10

0

10

-2

10

-4

10

-6

10
0.00

0.05

0.10

0.15

0.20

0.25

0.30

0.35

0.40

Electric Field (MV/cm)
Fig. 6. I–V characteristic of Cr2O3/CeO2 bilayered ﬁlm on STO substrate.
Below 0.24 MV/cm the ﬁlm was insulating.

The surface smoothness was degraded down to the Ra of
4.2 nm due to the bunching steps. The surface of the terrace
was so ﬂat that even steps of 1/3 unit cell did not appear. In
addition, spherical grains were also observed with the sizeof B40 lm. At 630 C as shown in Figs. 4e and 5e-L1 ﬁlm
3-dimensionally grew without a regularity in grains.
As the substrate temperature is increased, the migration
length increases and the density of nucleus sites as well asof grains decrease. At 400 C where coalescence did nottake place, the surface was covered by many grains andboundaries with the height of 3.5 nm due to the shortmigration length. However at 500 C where start to coalesce but not to obtain suﬃcient migration length to belarge grain, the deep grain boundary was expected to be
introduced. As increasing substrate temperature up to550 C, increased migration length led to larger grain andthen the density of deep grain boundary was reduced. At
500–550 C, the terrace in one grain was composed of several steps of 1/3 unit cell at most, indicating the 2-dimensional island growth. The unit of –O–Cr–O–Cr–O–,
corresponding to the height of 1/3 unit cell, stacks along
the [1 1 0 2] direction in Cr2O3 single crystal with rhombohedral structure. In the highly oriented Cr2O3 ﬁlms, it isreasonable that the integer multiples of 1/3 unit cell appear
because of the repulsive force between O2 ions at theinterface of stacked 1/3 unit cell. Above 600 C despite sufﬁcient migration length to be large grain, Cr2O3 is expected
to be decomposed due to deﬁcient of oxygen, demonstrated
by the results of XRD and the present of spherical grains.
Fig. 6 is a leakage current density as a function of electric ﬁeld (I–V) of a bilayered ﬁlm of Cr2O3/CeO2//

STO(1 0 0) observed at room temperature. The thickness
of Cr2O3 and CeO2 was 220 nm and approximately
10 nm, respectively. The diameter of measurement area
was B0.45 mm. From the XRD result Cr2O3 ﬁlm was oriented along the c-direction normal to the substrate surface.
Leakage current rapidly increased above 0.24 MV/cm indicating that the Cr2O3 ﬁlm works as a ME material below
the electric ﬁeld.
We estimated the induced magnetic ﬁeld tentatively withthe obtained data. Assuming that the ME coeﬃcient of
Cr2O3 along the c-direction around 40 K azz is 10 4 (inCGS – gauss unit) and the demagnetizing ﬁeld is neglected,
induced magnetic ﬁeld is estimated at 1.0 G [1–4]. For
example, to obtain 0.5 mA by a shielding current, whichis representative for base current in SFQ device, with100 lm2 and 100 pH superconducting loop, 5 G is
required. We expect to achieve such the value by optimizing the growth condition of CeO2 for epitaxial growth of
Cr2O3 on it.
4. Conclusion
We propose to apply the ME materials of single phasecompounds to the SFQ device as the replacement of a basecurrent and stabilization of IC spread. The ME material
generates induced magnetic ﬁeld with electric ﬁeld applied.
As a representative ME material, Cr2O3 thin ﬁlms weregrown on R-cut sapphire and STO substrates in order tointegrate with high-Tc superconductors. At 550 C,
0.26 Pa and a ﬂow rate of Ar:O2 = 4:1, highly oriented
Cr2O3 thin ﬁlms grew on sapphire substrate, the surfaceof which was composed of coalesced rectangular grains
with the size of 200 · 200–600 nm2. The surface roughness
Ra was 1.9 nm. Estimated induced magnetic ﬁeld from theresult of Cr2O3/CeO2//STO was 1.0 G around 40 K. To
eliminate the induced magnetic ﬂux, shielding current is
expect to ﬂow superconducting circuit, which is substituted
for a base current of SFQ device by a current source.
References[1] K.K. Likharev, K. Semenov, IEEE Trans. Appl. Supercond. 1 (1991) 3.
[2] M. Fiebig, J. Phys. D: Appl. Phys. 38 (2005) R123.
[3] H. Wigelmann, A.G.M. Jansen, P. Wyder, J.P. Rivera, H. Schmid,
Ferroelectrics 162 (1994) 141.
[4] N. Iwata, K. Matsuo, N. Ootsuka, H. Yamamoto, Mat. Res. Soc.
Symp. Proc. 811 (2004) 425.

Physica C 463–465 (2007) 965–968
www.elsevier.com/locate/physc

Improvement of ramp fabrication for high-Tc ramp-edge
Josephson junction
K. Igarashi a,b,*, K. Higuchi a,b, H. Wakana a, S. Adachi a, N. Iwata b,
H. Yamamoto b, K. Tanabe aaSuperconductivity Research Laboratory, ISTEC, 1-10-13 Shinonome, Koto-ku, Tokyo 135-0062, Japanb
Nihon University, 7-24-1, Narashinodai, Funabashi-shi, Chiba 274-8501, Japan
Accepted 6 March 2007
Available online 29 May 2007

Abstract
The fabrication process of ramps for HTS ramp-edge Josephson junctions was studied. We reﬁned a photolithography method for
producing ramps having smooth surfaces without a tail structure and similar angles in a chip. Optimization of resist reﬂowing parameters
was crucially important. It was found that excessive reﬂowing gave rise to an increased spread in the angles. We demonstrated fabricationof smooth ramps with a small spread in their angles using a mask pattern for a toggle ﬂip-ﬂop circuit pattern by using the optimized
reﬂowing parameters.
 2007 Published by Elsevier B.V.
PACS: 81.65.Cf; 74.72.Bk; 82.37.Gk

1. Introduction
Ramp-edge Josephson junctions with interface-modiﬁed
barriers using YBCO-type high-Tc superconductors (HTS)
have been widely studied because of small spread of junction properties [1]. Before deposition of the counter electrode, the ramp surface of the base-electrode waschanged to an amorphous structure by Ar ion bombardment. After deposition of the counter electrode on themodiﬁed ramp, the barrier layer was formed at the interface of the two electrodes. In fabrication of the junctions,several parameters that inﬂuence junction properties havebeen reported; for example, substrate temperature during
the counter-electrode deposition [2], irradiation angle during Ar ion bombardment to form the modiﬁed ramp surfaces [3] and so on. The ramp angle is also one of the*Corresponding author. Address: Nihon University, 7-24-1, Narashinodai, Funabashi-shi, Chiba 274-8501, Japan. Tel.: +81 3 3536 5709; fax:
+81 3 3536 5717.
E-mail address: kyu_igarashi_1214@istec.or.jp (K. Igarashi).

0921-4534/$ - see front matter  2007 Published by Elsevier B.V.
doi:10.1016/j.physc.2007.03.474

parameters inﬂuencing the spread of junction properties[4]. It is necessary to develop a fabrication technique forthe ramps having smooth surfaces and well-controlled
angles [5,6]. To reduce the spread of junction properties,
many ramps with the identical shape must be fabricated.
To fabricate the ramp-edge structures, the standard
photolithography with reﬂowed photoresist and Ar ionbeam etching is generally applied. The cross-sectional
images of an insulator/base-electrode bilayer ﬁlm covered
with reﬂowed photoresist before and after ion-beam etching are sketched in Fig. 1a. The ion-beam is irradiatedalong the slope of the reﬂowed photoresist. A smooth ramp
can be fabricated. Horstmann et al. [6] pointed out thenecessity of the reﬂowing process to make smooth ramps.
They tried to fabricate the ramps without reﬂowing a photoresist. Two parts of the slope could be distinguished in
the ramp. A steeper part of approximately 20 was followed by a long and shallow tail of only a few degrees. Thissituation is sketched in Fig. 1b. The resist stencil shades anarea of signiﬁcant width for incident angle b < 90. During
the ion irradiation, not only thin ﬁlms but also photoresist

966

K. Igarashi et al. / Physica C 463–465 (2007) 965–968

Fig. 1. Schematic illustrations for formation of the ramp proﬁles using (a) a reﬂowed photoresist and (b) a not-reﬂowed one.

are etched. As the photoresist is reduced, the shaded area
becomes narrower. This narrowing produces a tail structure in the etched ﬁlm. On the other hand, Blank et al.
[7] and Schoop et al. [8] developed a fabrication method
not using the reﬂowing process. They used TiN as the hard
and thin mask, instead of photoresist. However, their
method needs an additional vacuum chamber for TiNdeposition.
Recently, we have fabricated HTS electronic devices
containing many interface-modiﬁed ramp-edge junctions[9]. A 1r-spread in critical current of approximately 8%
was achieved for a 1000-junction series-array on a superconducting ground plane [10]. Proper operations of several
elementary circuits for SFQ electronic devices were demonstrated [11]. For further reduction of the spread of thesejunctions, we initiated reinvestigation of our fabrication
process. In this work, we focus on the ramp structure
before the counter-electrode deposition. In our process,the resist reﬂowing process was chosen for convenience tofabricate the ramps. We investigated the structures oframps located at diﬀerent positions in a chip using an
atomic force microscope (AFM) and observed the spreadin ramp angle. After discussion of the correlation between
the ramp structure and the shape of the reﬂowed photoresist, we propose a more suitable reﬂowing process to fabricate smooth ramps with the same angle and demonstrate it.
2. Experimental
Multilayer structures including an Y0.9La0.2Ba1.9Cu3Oy(La-YBCO) superconductor and a SrSnO3 (SSO) insulator
were prepared on 15 · 15 mm2 MgO (1 0 0) substrate by oﬀaxis magnetron sputtering. Their sequence was SSO(250 nm)/La-YBCO (200 nm)/SSO (250 nm)/La-YBCO(200 nm)//MgO. The upper and the lower La-YBCO layers
are a base-electrode and a ground plane, respectively.
Ramps were formed at the edges of the upper SSO(250 nm)/La-YBCO (200 nm) bilayer patterns by standard

photolithography with resist reﬂowing and Ar ion-beametching.
Our standard process for ramp fabrication is as follows.
At ﬁrst, a mask of photoresist was made by spin-coating
Shipley S1813 and baked in an electric furnace at 85 C
for 20 min. The thickness of the mask was 1.4 lm. After
patterning by ultraviolet exposure and photoresist development, the sample was baked at 145 C for 5 min to reﬂow
the photoresist. Ramps were formed by Ar ion-beam etching with an acceleration voltage of 280 V and an incident
angle of 30 to the substrate with sample rotation. Finally,
the mask was removed by O2-plasma ashing. We tried
using another mask of Shipley S1808 (0.8-lm thick). The
thinner mask was baked under diﬀerent conditions.
For morphology studies on the obtained ramps, an‘‘open air’’ Nanopics 2000 was used in AFM dumping
mode. The ramp angle for the base-electrode was
evaluated.
3. Results and discussion
We fabricated ramps by our standard process. Ramps
with no tail were successfully obtained. Their angles were
investigated by measuring a line proﬁle of the ramps fabricated at edges of patterned bilayers with a variety of sizes,
as illustrated in the inset of Fig. 2. Three samples were prepared by the same procedure. The ramp angles are plottedagainst the length of the pattern, l, in Fig. 2. It is found thatthe ramp angle depends on the size of the patterned bilayer.
Sato et al. [10] investigated the shapes of photoresists(reﬂowed and not reﬂowed) and the formed ramps usinga scanning electron microscope (SEM). They describedthat the shape of photoresist was copied to the ramp. In
our experiments, the ramps with signiﬁcantly steeper angles
were obtained for l = 10 and 20 lm, as shown in Fig. 2. It
is deduced that the shapes of photoresist at the edges of thebilayer patterns with large and small size must be diﬀerent
after reﬂowing.

K. Igarashi et al. / Physica C 463–465 (2007) 965–968

967

Fig. 2. Ramp angles for these samples including patterned insulator/baseelectrode bilayers with various sizes, illustrated in the inset.

Before the reﬂowing, the photoresist at the edges has thesame shape regardless the size of the bilayer pattern. As thebaking for reﬂowing proceeds, the shape of the resist
deforms gradually. In the above experiment, the 5-minbaking made a diﬀerence in the shapes of the photoresistat the edges. It seems that the diﬀerence resulted in theobserved spread in ramp angle, as shown in Fig. 2. It is
expected that baking for a shorter period would makethe diﬀerence less signiﬁcant. Optimization of the reﬂowing
condition, or the baking period and/or temperature, is crucially important. However, insuﬃcient baking cannot eliminate the shaded portion which produces a tail structure(see Fig. 1). We changed the thick resist (1.4 lm) to a thinner one (0.8 lm) and optimized the reﬂowing conditions,that is the baking period at 145 C.

Fig. 3. AFM line proﬁles for the ramps. The ramps were fabricated using
a thinner mask of S1808 (thickness, 0.7 nm) and reﬂowed at 145 C for (a)1, (b) 2 and (c) 5 min, respectively.

Fig. 3 shows the line proﬁles for the ramps fabricatedusing the photoresists reﬂowed for 1, 2 and 5 min. Reﬂowing for 1 min was not enough to eliminate the tail. Smooth
ramps without tails were fabricated using resist reﬂowed
for 2 min and more. The ramp angles were 32.5, 30 and25 for the reﬂowing periods of 1, 2 and 5 min, respectively.

Fig. 4. (a) The optical microscope image of a toggle ﬂip-ﬂop circuit patterns and (b) the measured angles of the ramps. The numbered arrows in the image
indicate the measured positions. The square and closed circles are ramp-angles fabricated using a thinner photoresist mask with 2-min-reﬂowing and a
thicker photoresist mask with 5-min-reﬂowing, respectively.

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K. Igarashi et al. / Physica C 463–465 (2007) 965–968

The spreads in the ramp-edge angles made using 2-minand 5-min-reﬂowed resists were compared. The angles foreight edges of bilayer patterns with diﬀerent size
(15 · 17 lm2–65 · 80 lm2) in a chip were measured. The
2-min- and 5-min-reﬂowed resists yielded ramps with30.0 ± 0.5 and 25.0 ± 2.0, respectively. The shorter
baking period for reﬂowing resist resulted in the smaller
spread, in accordance with our expectation. It is thus demonstrated that the use of the thin mask reﬂowed for 2 min
can produce smooth ramps without a tail and improves the
spread of the ramp angles in a chip.
We tried applying this improved method to prepare
ramps using a pattern of an actual elementary circuit.
The obtained ramp angles were compared with those fabricated using the standard method. Fig. 4a shows an optical
microscope image of patterned bilayer for a toggle ﬂip-ﬂop.
It contains base-electrodes with various size. The arrows in
the ﬁgure indicate the positions for junction fabrication.
The obtained results are plotted in Fig. 4b. The improved
and standard methods yielded ramps with 29.5 ± 0.7and 25.5 ± 0.3, respectively. A smaller spread was actually achieved by using the improved method. Additionally,
we tried to draw the break lines along the obtained data by
eye guide in Fig. 4b. It is found that two curves wind similarly. We considered that this tendency is due to the designof the circuit. For the further improvement of the ramp
fabrication, we need to focus on the circuit design.
4. Conclusion
We investigated the ramp-edge angles of base-electrodes
with a variety of size in a chip. An appreciable spread wasobserved and the deviation in the angle depended on thesize. It was speculated that the spread in the angle originated from a diﬀerence in the edge-shape of resist mask.
We reﬁned a photolithography method containing a resist
reﬂowing process. A method to shape the resist mask hav-

ing identical structures at edges was proposed. A thin photoresist mask was utilized in order to avoid formation of atail structure in the ramps. It was found that excessive
reﬂowing gave rise to an increased spread in the angle.
By using the optimized resist reﬂowing parameters, wedemonstrated fabrication of smooth ramps with a small
spread in their angles for a toggle ﬂip-ﬂop circuit pattern.
Acknowledgment
This work was supported by the New Energy and Industrial Technology Development Organization (NEDO) as
Superconductors Network Device Project.
References[1] T. Satoh, M. Hidaka, S. Tahara, IEEE Trans. Appl. Supercond. 9(1999) 3141.
[2] H. Wakana, S. Adachi, M. Horibe, Y. Ishimaru, O. Horibe, Y.
Tarutani, K. Tanabe, Physica C 378–381 (2002) 1353.
[3] H. Wakana, S. Adachi, M. Horibe, Y. Ishimaru, O. Horibe, Y.
Tarutani, K. Tanabe, Jpn. J. Appl. Phys. 41 (2002) 239.
[4] M. Horibe, Y. Inagaki, G. Matsuda, N. Hayashi, K. Kawai, M.
Maruyama, A. Fujimaki, H. Hayakawa, Supercond. Sci. Technol. 12(1999) 726.
[5] T. Makita, K. Toma, K. Ishikawa, H. Zama, T. Utagawa, U.
Kawabe, K. Tanabe, IEEE Trans. Appl. Supercond. 11 (2001) 155.
[6] C. Horstmann, P. Leinenbach, A. Engelhardt, R. Gerber, J.L. Jia, R.
Dittmann, U. Memmert, U. Hartmann, A.I. Braginski, Physica C 302(1998) 176.
[7] D.H.A. Blank, G.J.H.M. Rijnders, R.M.H. Bergs, M.A.J. Verhoeven,
H. Rogalla, IEEE Trans. Appl. Supercond. 7 (1997) 3323.
[8] U. Schoop, M. Schonecke, S. Thienhaus, S. Schymon, L. Alﬀ, R.
Gross, Physica C 351 (2001) 200.
[9] Y. Tarutani, Y. Ishimaru, H. Wakana, S. Adachi, K. Nakayama, Y.
Oshikubo, O. Horibe, Y. Morimoto, K. Tanabe, Physica C 412–414(2004) 1528.
[10] H. Sato, F.J.G. Roesthuis, A.H. Sonnenberg, A.J.H.M. Rijnders, H.
Rogalla, D.H.A. Blank, Supercond. Sci. Technol. 13 (2000) 522.

Physica C 445–448 (2006) 917–920
www.elsevier.com/locate/physc

Fabrication of ramp-edge Josephson junction by PLD
method using double plume
K. Higuchi a,b,*, K. Igarashi a,b, H. Wakana a, K. Nakayama a, S. Adachi a,
N. Iwata b, H. Yamamoto b, K. Tanabe aaSuperconductivity Research Laboratory, ISTEC, 1-10-13 Shinonome, Koto-ku, Tokyo 135-0062, Japanb
Nihon University, 7-24-1, Narashinodai, Funabashi-shi, Chiba 274-8501, Japan
Available online 21 June 2006

Abstract
Fabrication of high-Tc ramp-edge Josephson junctions with interface-modiﬁed barriers in a wide area have been studied. A novel
pulsed laser deposition apparatus with a double laser plume was newly designed. We succeeded in observing RSJ I–V characteristics withan excess current ratio as low as approximately 15% for all junctions dispersively located on a chip with a size of 15 · 15 mm2. It was
found that the characteristics of the junctions signiﬁcantly depend on the geometric conﬁguration of the laser plume and the ramps during counter-layer deposition.
 2006 Elsevier B.V. All rights reserved.
PACS: 74.72.Bk; 74.50.+r; 81.15.Fg
Keywords: Y0.9La0.2Ba0.9Cu3Oy; Ramp-edge Josephson junction; Double laser plume

1. Introduction
Multilayer structures of various oxide thin ﬁlms including Josephson junctions are indispensable in high-Tc electronic devices such as single ﬂux quantum (SFQ) digital
circuits [1–4]. SFQ circuits require reproducible fabricationof Josephson junctions with high IcRn products and small
1 rIc spreads. Interface-modiﬁed ramp-edge Josephson
junctions have attracted much attention as reliable junctions, because of their higher IcRn products and smaller
Ic spreads in a chip than those reported in other types ofhigh-Tc Josephson junctions [5,6]. The barriers of interface-modiﬁed junctions are fabricated from an amorphouslayer, that is, made by an ion bombardment process. This
amorphous layer is changed to the barrier layer through*Corresponding author. Address: Superconductivity Research Laboratory, ISTEC, 1-10-13 Shinonome, Koto-ku, Tokyo 135-0062, Japan. Tel.:+81 3 3536 5709; fax: +81 3 3536 5717.
E-mail addresses: Higuchi@istec.or.jp (K. Higuchi), adachi@istec.or.jp(S. Adachi).
0921-4534/$ - see front matter  2006 Elsevier B.V. All rights reserved.
doi:10.1016/j.physc.2006.05.071

appropriate annealing and counter-layer deposition processes [7]. Wakana et al. [8] have recently fabricatedinterface-modiﬁed ramp-edge Josephson junctions with
La-doped YbBa2Cu3Oy (Yb0.9La0.2Ba1.9Cu3Oy:La–YbBCO)
counter-layer deposited by a pulsed laser deposition (PLD)
method. It is well known that the properties of these junctions have a strong dependence on the recrystallization
process of the interface-modiﬁed barrier. In the case ofcounter-layer deposition by PLD, in particular, the characteristics of the interface-modiﬁed ramp-edge Josephson
junctions signiﬁcantly depend on the geometric conﬁguration of the laser plume and the ramps during counter-layer
deposition [8].
In order to execute counter-layer deposition in a morehomogeneous condition for the ramp surfaces, we designed
a novel PLD apparatus with a double laser plume. A photographic view of the double plume during deposition isshown in Fig. 1. We expected that use of the double laser
plume would enable us to fabricate ramp-edge Josephson
junctions with more homogeneous characteristics in a
wider area on a chip than the case for deposition with a

918

K. Higuchi et al. / Physica C 445–448 (2006) 917–920

Fig. 1. Double laser plume during deposition.

single plume. In this work, we fabricated the junctions on a
15 · 15 mm2 substrate using the novel apparatus and investigated their properties. Samples with the counter-layers
deposited employing diﬀerent conﬁguration of the double
laser plume and the substrate were prepared. Propertiesof the junctions that were put all over a chip were
examined.
2. Experimental
Firstly, multilayer ﬁlms including a base-electrode layer(BE) on a ground plane (GP) were prepared on a
15 · 15 mm2 MgO (1 0 0) substrate by oﬀ-axis magnetron
sputtering. Y0.9La0.2Ba1.9Cu3Oy (La–YBCO) and SrSnO3(SSO) were used as a superconductor and an insulator,
respectively. The layer sequence was SSO (250 nm)/La–
YBCO (BE; 200 nm)/SSO (250 nm)/La–YBCO (GP;
200 nm)/MgO substrate. Ramp-edge structures for theupper two layers were fabricated by the standard photolithography using a resist reﬂow process and Ar ion millingat 280 V with an incident angle of 30. After removing thephotoresist mask, an amorphous layer was formed on theramp surfaces by Ar ion bombardment for 3 min at anacceleration voltage of 500 V and the incident angle of90 to the substrate.
A La–YbBCO counter-layer was deposited by the twostep deposition process [9]. A 20-nm-thick primary-layer,
which considerably inﬂuences the junction properties, wasﬁrstly deposited. Then a 200-nm-thick secondary-layer covered it to form a counter layer with low sheet inductance.
Fig. 2 shows schematic illustrations of the substrate, target
and the laser plumes for the deposition of the primary andsecondary layers. During deposition of the former layer,
the substrate was not rotated and the double plume stayedat two positions separated by a distance of d, as shown in
Fig. 2(a). Points A, B, C and D in the ﬁgure are given formarking the projection of the plume centers in a plane
including the substrate surface. We tried two distances ofd = 15 and 25 mm. Deposition was carried out at a laser
frequency of 5 Hz in 20-Pa oxygen atmosphere. The substrate-target distance was 67 mm. The energy density on
the target was 1.7 J/cm2 for each beam. For the secondary-layer deposition, the substrate was rotated and the
double plume with a constant distance of 25 mm wasscanned on the target, as shown in Fig. 2(b). The laser frequency, energy density and the substrate-target distancewere changed to 2 Hz, 1.2 J/cm2 and 70 mm, respectively.
In this work, we did not use Ag-paste for ﬁxing a substrate on a sample holder to improve reproducibility. Possibility of contamination due to Ag-paste can be avoided.
The substrate was directly heated by thermal radiationfrom a heater located at the backside of the substrate. Tem-

Fig. 2. Schematic illustration of the substrate, the target and the laser plumes for depositions of (a) the primary and (b) the secondary layers. Characters of
A–D are markers of the plume centers.

K. Higuchi et al. / Physica C 445–448 (2006) 917–920

perature of the heater was set at 840 C during the counterlayer deposition. The actual substrate temperature was
unknown. Finally, a 450-nm-thick Au ﬁlm was depositedon the La–YbBCO counter-layer, and the counter-electrode was patterned to width of 5 lm by photolithography
and Ar ion milling. Current–voltage (I–V) characteristics
for the fabricated junctions were evaluated at 4.2 K by afour-probe method.
3. Results and discussion
Fig. 3 shows typical I–V characteristics for the junction
with the primary-layer prepared at d = 15 mm. The
obtained junction exhibits resistively-shunted-junction(RSJ) I–V characteristics with an excess current ratio as
low as approximately 15%. All the junctions on a chipshowed RSJ I–V characteristics with similar excess current.
The distribution of Ic values on a chip is shown in Fig. 4.

Fig. 3. Typical I–V characteristics for the junction with a primary layer
prepared at d = 15 mm.

919

Directions of all the examined ramps were identical, as
drawn in the inset of Fig. 4. The height of each column
indicates the average Ic for 10 junctions within an area of
200 · 200 lm2. The 1 rIc spreads in each column was
13.7–29.7%. The columns in the ﬁgure indicates that thejunctions on the chip with a size of 15 · 15 mm2 exhibited
Ic = 0.21–1.1 mA. The 1 r spread in Ic was 47%. It wasreported that the spreads of Ic for ramp-edge junctions
without ground plane in a 10 · 10 mm2 substrate was
20% [10], being smaller than the obtained spread of 47%
in the present study. We thought that the larger chip size
and the junction fabrication on a ground plane made thereduction of the spreads diﬃcult.
A contour map of Ic is also drawn in Fig. 4. The 15 · 15
area of a chip is divided to three regions in which the junctions exhibited Ic over 1 mA, 1.0–0.5 mA and less than
0.5 mA. It can be seen that Ic values are not distributed
randomly on the chip. Assuming that the observed variation in Ic was caused by inhomogeneous substrate temperature, the Ic values should be concentrically distributed.
The observed distribution is not so, as seen in Fig. 4. It suggests that the spatial conﬁguration of the plumes and theramps rather inﬂuenced the junction characteristics.
In the case of d = 25 mm, only a few junctions exhibited
RSJ I–V characteristics with an excess current of 50%, asshown in Fig. 5(a). The other junctions show ﬂux-ﬂow-like
I–V characteristics as in Fig. 5(b). The Ic mapping is illustrated in Fig. 6. This result implies that the formed barrier
layers were inhomogeneous. The obtained results indicatethat the junction properties signiﬁcantly depend on geometric conﬁguration of the laser plume and the ramps during counter-layer deposition even for the case of double
plume PLD. Further optimization of the conﬁguration isneeded to improve the Ic spreads.

Fig. 4. Mapping of the average Ic values on a chip with a size of 15 · 15 mm2. The primary layer prepared at d = 15 mm. Characters of A and B aremarkers of the plume centers. Illustration of a typical junction is given in the inset. Directions of all junctions are identical. A contour map of Ic is in theinset.

920

K. Higuchi et al. / Physica C 445–448 (2006) 917–920

Fig. 5. Typical I–V characteristics for the junctions with a primary layer prepared at d = 25 mm.

layer deposition considerably inﬂuenced the junction
characteristics.
Acknowledgements
This work was supported by the New Energy and Industrial Technology Development Organization (NEDO) as
Superconductors Network Device Project.
References
Fig. 6. Mapping of the average Ic values on a chip with a size of
15 · 15 mm2. The primary layer prepared at d = 25 mm. Characters of C
and D are markers of the plume centers. Directions of all junctions are
identical, (see Fig. 4).

4. Conclusion
We fabricated interface-modiﬁed ramp-edge Josephson
junctions using a newly designed PLD apparatus with adouble laser plume. The junctions exhibiting RSJ I–V characteristics with an excess current 15% on a chip with a sizeof 15 · 15 mm2. It proved that the PLD with a double laser
plume could be utilized to realize high-Tc superconducting
electronic devices with a large scale. Distribution of Ic on achip was investigated. It was found that the spatial conﬁguration of the plumes and the ramps during the counter-[1] A. Fujimaki, K. Kawai, N. Hayashi, M. Horibe, M. Maruyama,
H. Hayakawa, IEEE Trans. Appl. Supercond. 9 (1999) 3436.
[2] T. Satoh, J.G. Wen, M. Hidaka, S. Tahara, N. Koshizuka, S. Tanaka,
IEEE Trans. Appl. Supercond. 11 (2001) 770.
[3] J. Yoshida, S. Inoue, H. Sugiyama, T. Nagano, IEEE Trans. Appl.
Supercond. 11 (2001) 784.
[4] Y. Soutome, T. Fukazawa, K. Saitoh, A. Tsukamoto, K. Takagi,
IEICE Trans. Electron E-85C (2002) 759.
[5] B.H. Moeckly, K. Char, Appl. Phys. Lett. 71 (1997) 2526.
[6] T. Satoh, M. Hidaka, S. Tahara, IEEE Trans. Appl. Supercond. 9(2001) 770.
[7] J.G. Wen, N. Koshizuka, S. Tanaka, T. Satoh, M. Hidaka, S. Tahara,Appl. Phys. Lett. 75 (1999) 2470.
[8] H. Wakana, S. Adachi, M. Horibe, Y. Ishimaru, O. Horibe,
Y. Tarutani, K. Tanabe, Jpn. J. Appl. Phys. 41 (2002) L 239.
[9] H. Wakana, S. Adachi, M. Horibe, Y. Ishimaru, Y. Oshikubo,
Y. Tarutani, K. Tanabe, Jpn. J. Appl. Phys. 41 (2002) L 1366.
[10] H. Wakana, S. Adachi, M. Horibe, Y. Ishimaru, O. Horibe,
Y. Tarutani, K. Tanabe, Physica C 378–381 (2002) 1353.

Physica C 445–448 (2006) 904–907
www.elsevier.com/locate/physc

IcRn product of YBa2Cu3O7 x ramp-edge junctionsat temperatures higher than 30 K
Y. Morimoto a,b, Y. Tarutani a,*, H. Wakana a, S. Adachi a, K. Tsubone a, Y. Ishimaru a,
K. Nakayama a, Y. Oshikubo a, O. Horibe a, N. Iwata b, H. Yamamoto b, K. Tanabe aaSuperconductivity Research Laboratory, ISTEC, 1-10-13, Shinonome, Koto-ku, Tokyo 135-0062, Japanb
Nihon University, 7-24-1, Narashinodai, Funabashi-shi, Chiba 274-8501, Japan

Abstract
We examined temperature dependence of the IcRn (Ic being the critical current and Rn being the normal-state resistance) product forinterface-engineered YBa2Cu3O7 x ramp-edge junctions formed on a groundplane. There were two types for the temperature dependenceof the IcRn product. One was the linear temperature dependence, and the other was a dependence that shows the steeper drop with temperature. There were two factors that were responsible for the temperature dependence of the IcRn value. These were critical temperature
Tc of the junctions and temperature dependence of the Rn value. The IcRn value more than 1 mV at 40 K was obtained for the junctionthat had the Tc of 84 K and the positive coeﬃcient of Rn value on temperature.
Ó 2006 Published by Elsevier B.V.
PACS: 74.78.Bz; 85.25.Dq
Keywords: IcRn product; Oxide SFQ circuit; Hysteresis; Excess current

1. Introduction
In an oxide single-ﬂux-quantum (SFQ) circuit, IcRn
product of Josephson junctions is very important for
improving an operating frequency [1]. The IcRn productis also important for a driver circuit because an output
voltage is almost proportional to the IcRn product. The
IcRn product much more than 2 mV at 4.2 K has easily
been obtained for oxide junctions. This value is more thanﬁve times as much as the IcRn product of resister-shunted
Nb junctions that have a critical current density less than2000 A/cm2.
Oxide SFQ circuits are supposed to be operated at temperatures from 30 K to 40 K when oxide SFQ chips arecooled down by a single-stage refrigerator. The IcRn product required for the oxide junctions may be more than*Corresponding author. Tel.: +81 3 3536 5709; fax: +81 3 3536 5705.
E-mail address: y-morimoto4038@istec.or.jp (Y. Tarutani).

0921-4534/$ - see front matter Ó 2006 Published by Elsevier B.V.
doi:10.1016/j.physc.2006.05.068

1 mV at these operating temperatures. This is because the
IcRn value of 1 mV is the least minimum value for operating
SFQ circuits at frequencies more than 100 GHz [2]. However, the IcRn product rapidly decreases as a temperature
is increased. The temperature dependence of the IcRn value
for oxide junctions has not systematically been studied.
In this report, YBa2Cu3O7 x ramp-edge junctions arefabricated and their IcRn products are investigated. Temperature dependence of the IcRn products is measured andthe IcRn products from 30 K to 40 K are compared among
junctions that show diﬀerent characteristics. Requirement
to the characteristics for the junctions whose IcRn products
are more than 1 mV from 30 K to 40 K are found out for
applying the ramp-edge junctions to SFQ circuits.
2. Junction fabrication and measurement
The ramp-edge junctions were made by the interfaceengineering method [3,4]. A La0.2Y0.9Ba1.9Cu3O7 x ﬁlm

Y. Morimoto et al. / Physica C 445–448 (2006) 904–907

was used for a base electrode and a La0.2Yb0.9Ba1.9Cu3O7 x ﬁlm was used for a counter electrode, respectively. The La0.2Y0.9Ba1.9Cu3O7 x was also used for agroundplane. SrSnO3 ﬁlms was used for insulating layers
between two of these electrodes. A MgO single crystal
was used for a substrate.
The base electrode and the counter electrode were
deposited by the sputtering method and the pulse-laserablation method, respectively. Substrate temperature during the counter layer deposition was changed among junction fabrication runs. Other fabrication parameters
including the base layer deposition condition were kept
the same. Junction barrier layer was formed by irradiating
Ar beam onto surface of the base electrode.
Current–voltage (I–V) characteristics of the ramp-edge
junctions were measured by a four-terminal probe method.
Chip temperatures were measured by using a carbon-glass
resistor. A package and socket on which a junction chip
was mounted were kept in close contact with a Cu block.
The Cu block was covered by double permalloy cylinders
that magnetically shielded the chip. The chip was cooled

905

down by keeping the sample holder into a liquid He cryostat. The chip temperature was controlled by adjusting thesample distance from the liquid He surface.
3. IcRn product of ramp-edge junctions
Current–voltage (I–V) curves of six ramp-edge junctions
were measured at temperatures from 4.2 K to 90 K. Thesejunctions were denoted as (A) throughout (F).
The I–V curves were measured at every four or ﬁve
degrees from 4.2 K throughout 90 K. The I–V curves ofjunctions (A), (B), (C) and (F) measured at 40 K are shownin Fig. 1. These I–V curves are representing typical characteristics of ramp-edge junctions. These I–V curves diﬀer in
Ic or IcRn values. These junctions show no hysteresis in the
I–V curves although all these junctions showed diﬀerent
degrees of hysteresis at 4.2 K.
Temperature dependences of the IcRn products for thejunctions (A), (B), (C) and (F) are shown in Fig. 2. The
IcRn values were plotted from the I–V curves. Junctions(A) and (B) show the IcRn products at 4.2 K more thanFig. 1. I–V curves of ramp-edge junctions (A), (B), (C) and (F). Measured temperature was 40 K.

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Y. Morimoto et al. / Physica C 445–448 (2006) 904–907

Dependences of the IcRn (40 K/4.2 K) values on thejunction Tc are shown in Fig. 3. Tc of the base electrodewas from 85 K to 86 K. The Tc of junction was corresponding to that of counter electrodes. As shown in
Fig. 3, ratio of the IcRn product at 40 K to that 4.2 K
did not much decrease when the junction Tc was increased
to that of the base electrode. Dependence of the junction
Tc on the substrate temperature is shown in Fig. 4 The
junction Tc increased as the deposition temperature wasincreased.
Temperature dependences of the normal-state resistanceRn for the junctions from (A) to (F) are shown in Fig. 5.
Samples (A) and (D) show positive coeﬃcients of the Rn

Fig. 2. The IcRn product – temperature characteristics.

2 mV. On the other hand, junctions (C) and (F) show the
IcRn products less than 1.5 mV and 1 mV, respectively.
Slope of the IcRn products for junction (A) and junction(B) are much diﬀerent. The IcRn product of junction (B)at 40 K is almost comparable to that of sample (C) as aresult of the steeper temperature dependence. Temperaturedependences of the IcRn products for junction (D) andjunction (E) resembled those for junctions (A) and (B),respectively. From these result, requirements for the IcRnvalue to be more than 1 mV at 40 K are not only the IcRn
product being much more than 2 mV at 4.2 K but also
dependence of the IcRn product on temperature being
weak. Therefore, the ratio of the IcRn value at 40 K to thatat 4.2 K is an important parameter. This value is denotedas IcRn (40 K/4.2 K).

Fig. 3. Dependence of the IcRn (40 K/4.2 K) value on the junction Tc.

Fig. 4. Substrate temperature dependence of the Tc.

Fig. 5. Dependence of the Rn value on temperature.

Y. Morimoto et al. / Physica C 445–448 (2006) 904–907

value on temperature. On the other hand, samples (B) and(F) show negative coeﬃcients of the Rn value on temperature. The IcRn values of samples (A) and (D) almost linearly depended on temperature. Dependences of the IcRnvalues on temperature for samples (B) and (F) was muchsteeper than that for samples (A) and (D). Decrease of
the Rn values with temperature increase for samples (B)and (F) was one of the factor for the much greater decreaseof the IcRn products at 40 K.
The temperature dependence of the IcRn value is determined by two factors. One is the typical temperaturedependence that is determined by Josephson formulae [5].
The other is the conduction mechanism through a barrier.
The normal resistance may change from 4.2 K to 90 K
depending on the conduction mechanism. The negative
coeﬃcient of the Rn value for junction (F) indicates thatthe conduction through barrier is due to carrier hopping[6]. In this case, normal current is enhanced at higher temperatures in addition to tunnel current that owes the supercurrent. This normal current component may havesuppressed the IcRn value for junction (F) at higher temperatures. In this manner, there are two factors that are
required to increase the IcRn value more than 1 mV at
40 K. One is the IcRn value at 4.2 K being more than
2 mV. The other is the temperature dependence of the IcRnvalue being linear. Both factors can be satisﬁed by increasing the Tc of the counter electrode to that of the base electrode. Such a junction has a higher critical current density.
It is necessary to decrease the junction size depending on
the Ic value.

907

4. Conclusions
In conclusion, the temperature dependence of the IcRn
products for the ramp-edge junctions was examined. There
were two types for the temperature dependence of the IcRn
product. One was the linear temperature dependence, andthe other was the nonlinear dependence that showed rapid
drop of the IcRn product with temperature. The junctions
having the IcRn value more than 1 mV at 40 K had the linear temperature dependence of the IcRn value in addition to
the IcRn value more than 2 mV at 4.2 K.
Acknowledgement
This work was supported by the New Energy and Industrial Technology Development Organization (NEDO) as
Superconductors Network Device Project.
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Tarutani, K. Tanabe, Physica C 392–396 (2003) 1362.
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Kawabe, K. Tanabe, Jpn. J. Appl. Phys. 41 (2002) 1998.
[5] T. Van Duzer, C.W. Turner, Principles of superconductive devices andcircuits, Elsevier, Amsterdam, 1980, p. 142.
[6] M. Kupriyanov, in: Y. Bando, H. Yamauchi (Eds.), Adv. Supercond.
5 (1992) 1049.

ARTICLE IN PRESS

Physica E 37 (2007) 49–53
www.elsevier.com/locate/physe

Position-selective growth of vertically aligned carbon nanotubes forapplication of electronic-measuring nanoprobes
Hiroki Okuyama, Nobuyuki Iwata, Hiroshi Yamamoto
College of Science and Technology, Nihon University, 7-24-1-401, Narashinodai, Funabashi, Chiba 274-8501, Japan
Available online 29 November 2006

Abstract
Position-selective growth of carbon nanotubes (CNTs) and vertically aligned CNTs (VACNTs) on patterned metal electrodes havebeen prepared by thermal chemical vapor deposition (TCVD) and DC plasma enhanced chemical vapor deposition (PECVD). We
propose newly a position-controlling method of CNTs by controlling not only a position of Ni as catalysts but also the morphology ofMo as underlayers for the catalysts. The position-selective growth of CNTs was achieved at the edges of the patterned metal by TCVD.
The morphologies of the Mo underlayer at the selected area were rough and porous. No CNTs grew on smooth Mo surfaces. The
minimum width of selectively grown CNTs, ca. 2.6 mm, was approximately one-eightieth of the patterned metal, 200 mm. VACNTs weresynthesized by a PECVD method, however, the VACNTs grew up all over the patterned metal. The Ni catalysts formed into ﬁneparticles on rough surfaces of the Mo underlayer. Then the selective growth was achieved by Ni ﬁne particles formed only at the edges ofthe metal pattern. The results of PECVD suggest that the plasma promoted the Ni catalysts to become ﬁne particles on smooth surfacesof Mo. Conclusively a position-controlling method of CNTs was demonstrated in the optimum conditions of the TCVD.
r 2006 Elsevier B.V. All rights reserved.
PACS: 81.05.Uw; 81.07.De; 81.15.Gh
Keywords: Cabon nanotube; Selective growth; Chemical vapor deposition; Vertical alignment

1. Introduction
Carbon nanotubes (CNTs) have been intensively studied
because of many featured characteristics such as highelectric conductivity [1], high permissible current density [2]and high mechanical strength in spite of high ﬂexibility [3].
These characteristics are available to prepare various kindsof nanostructured devices such as scanning probe microscopy tip [4–6], ﬁeld emitter [7–9] and so on. We are aiming
at realizing a novel nanoprobe with CNTs that enables toapproach it toward contact points on specimen easily
without an accurate control for measuring of electrictransport properties. In order to prepare such nanoscaled
CNT devices, the controlling position and direction of
CNTs is necessary. CNTs grow selectively from catalyst
Corresponding author. Tel./fax: +81 47 469 5457.

E-mail address: okuyama@yamanoya.ecs.cst.nihon-u.ac.jp(H. Okuyama).
1386-9477/$ - see front matter r 2006 Elsevier B.V. All rights reserved.
doi:10.1016/j.physe.2006.10.016

particles by using chemical vapor deposition (CVD) [10]. A
growth direction of CNTs is controlled by applying electricﬁeld during CVD growth [11].
The position-selective growth of CNTs is achieved byusing the two characteristics as follows. One is that the
CNTs grow selectively from ﬁne particle catalysts. The
other is that we can control the synthesis of the ﬁneparticle catalysts by noticing wettability on metal interfaces. Since wettability of metal interface is high, it ishard to form ﬁne particles as overlayers on smooth
underlayers. When the surfaces of underlayers are rough,overlayered metals form ﬁne particles. In this paper, we
propose newly a position-controlling method of CNTs
by controlling not only the position of catalysts but
also the surface morphology of metal underlayers. And
we tried to prepare position-selective growth of CNTs andvertically aligned CNTs (VACNTs) on patterned metalelectrodes for realization of the novel electronic measuring
nanoprobe.

ARTICLE IN PRESS50

H. Okuyama et al. / Physica E 37 (2007) 49–53

2. Experimental
Patterned metals of Ni catalysts and Mo underlayers
were deposited on quartz substrates through a metal mask
(thickness: 30 mm, pattern width: 200 mm, gap: 100 mm) by
RF magnetron sputtering. A background pressure of thechamber was 3.0  10 3 Pa, and the pressure during
deposition was ﬁxed at 2.0 Pa by feeding Ar gas. The
growth conditions were 120 1C of substrate temperature,
50 W of sputtering power, 1 min of sputtering time for Mo
underlayers, and 70 1C, 20 W, 1 min for Ni catalysts,
respectively. The deposition rates of the ﬁlms were 10 nm/
min for Mo, and 1 nm/min for Ni. In order to control
incident angles of sputtered atoms, aligned ﬁns were placed
above substrates during deposition. A schematic conﬁguration of the ﬁns and the metal mask is shown in Fig. 1.
Here, ‘‘shady side’’ represents the edges of the patterns thathide behind the mask. And ‘‘sunny side’’ represents theedges that do not hide behind the mask.
CNTs were grown by thermal chemical vapor deposition(TCVD) and DC plasma enhanced chemical vapor deposition (PECVD) method. The reactor tube of the CVD
equipment was heated up to 600 1C. The backgroundpressure was 0.1 Pa. The ﬂow of H2:Ar (10:200 ccm) was
introduced into the reactor tube while heating. In the caseof a TCVD method, mixture gas of ethylene (C2H4):H2:Ar(20:10:200 ccm) was introduced into the reactor tube for
30 min for the CNT growth. The total pressure wasmaintained at 100 kPa. In the case of the PECVD method,
the CNT growth was performed by inducing a DC glow
discharge and feeding mixture gas of C2H4:H2:Ar(5:50:50 ccm). The total pressure was maintained at
2 kPa. The DC plasma was produced by applying bias
voltage of
250 V between a substrate holder andgrounded anodes apart from the substrates by approximately 5 mm.
The morphology of patterned metals was investigated by
an atomic force microscopy (AFM, SII Nanotechnology

Inc. NanoNavi Station). The structure of CNTs wasobserved by a ﬁeld emission scanning electron microscope(SEM, Hitachi S-4500) operated at 15 kV.
3. Results and discussion
Fig. 2 shows AFM images (5  5 mm2) of the patterned
metal layers just prior to CNT growth. Figs. 2(a) and (b)show the pattern edges of patterned metal layers and the
mid areas of the patterned metal, respectively. Fig. 2(c)shows scan areas of Figs. 2(a) and (b). The left of Fig. 2(a)image is the substrate surfaces; the right is the patterned
metal layers. At the pattern edges, large Ni particles with aheight of ca. 30 nm were observed on small Mo particles
with a height of ca. 3 nm. At the mid areas of the pattern(Fig. 2(b)), only a few particles were observed on smoothsurfaces.
Fig. 3 shows SEM images of the pattern edges of thepatterned metal layers after TCVD varying sputtering
conditions. The patterned metal layers of Fig. 3(a) wereprepared without aligned ﬁns, Figs. 3(c) and (d) wereprepared by installing aligned ﬁns. The optical microscopy(OM) image of the patterned metal, which was sputteredthrough the ﬁns, is shown in Fig. 3(b). The array of 12
patterns was labeled points A to L as shown in Figs. 1 and
3(b). Figs. 3(c) and (d) show the results on the patternedges of point D. The images of Fig. 3(c) are the shadyside, Fig. 3(d) are the sunny side as shown in Fig. 1. All
SEM images revealed that CNTs grew selectively at theregion of just insides from the pattern edges whichappeared as the white contrast. The diameter of individual
CNT was approximately 30 nm. On the contrary CNTs did
not grow on the mid areas of the pattern. The position
selective growth width of CNTs (WCNT) were (a) 30, (c) 17
and (d) 2.6 mm, respectively. The diameter of the CNTs was
similar to that of the Ni particles, which was observed by
AFM. At the pattern edges, many small Mo particles wereobserved under the CNTs. When the ﬁns were installed

Fig. 1. A schematic conﬁguration of the ﬁns and the metal mask.

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H. Okuyama et al. / Physica E 37 (2007) 49–53

51

Fig. 2. AFM images (5  5 mm2) of patterned metal layers (a) and (b). The scanning area of the AFM images (c). (a) The pattern edges of patterned metal
layers. The left of the AFM image shows the substrate surfaces; the right shows the patterned metal layers. (b) The middle areas of the patterned metal.

(Fig. 3(b)), the contrast of the pattern was not uniform
because the thickness of metal pattern was not uniform.
Therefore, various thicknesses of metal patterns wereappeared simultaneously on this sample. The WCNT inevery pattern is summarized in Fig. 4. The gray belts
correspond to metal mask patterns, and white belts to thepattern gaps. The results show that the WCNT was
narrower in the sunny side than in the shady side. In
addition, the WCNT depended on the position of thepattern. The minimum WCNT was 2.6 mm. This width wasapproximately one-eightieth of the pattern width.

The position-selective growth of CNTs at the patternedges was achieved because CNTs grow from only ﬁneparticle catalysts. On the smooth surface of Mo underlayers, it is hard for Ni catalysts to become ﬁne particles
due to the high wettability. The rough morphology of Mo
underlayers resulted in ﬁne particles of Ni catalysts. The
rough morphology was observed in very thin Mo regions
due to ‘‘shadow effect’’ of a sputtering method. A
schematic shadow effect is illustrated in Fig. 5. Sputteredatoms arrive at substrates with various incident angles
because of scattering by atmospheric gasses. Then the

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H. Okuyama et al. / Physica E 37 (2007) 49–53

52

Fig. 3. SEM images of the pattern edges of patterned metal layers after the TCVD at varying sputtering conditions: (a) without aligned ﬁns, (c) ‘‘shadyside’’ by aligned ﬁns, and (d) ‘‘sunny side’’ by aligned ﬁns. (b) OM image of the patterned metal. CNTs grew in the narrow region which appeared as thewhite contrast. The WCNT was about (a) 30, (c) 17 and (d) 2.6 mm. The insets are each the magniﬁed image.

Metal pattern Gap

CNT grown width [m]40
Shady side30

20
Sunny side

Fig. 5. Schematic ‘‘shadow effect’’ appeared through the mask pattern.
The pattern edges thin down by deposition of shaded sputtered atoms
under the mask edges (indicated by the oval).

10

0
A

B

C

D

E

F G
Point

H

I

J

K

L

Fig. 4. The WCNT observed at each point in every pattern prepared. The
gray belts correspond to the metal mask patterns, and white belts to thepattern gaps.

pattern edges thin down by deposition of particles shaded
from the mask pattern. The WCNT depended on theposition of the pattern because the thickness of metal
pattern was not uniform as shown in Fig. 3(b). By

increasing the thickness of the metal pattern, the WCNT
decreased because thin regions of Mo decreased. Since theeffective widths of the shadow effect depended on theincident angle of sputtered atoms, the WCNT obtained in
the sunny side was quite different from that in the shadyside. An estimated mean free path of sputtered atoms wasca. 6 mm, and was slightly longer than the ﬁn–substratedistance, 5 mm. The installing of the ﬁns was effective forcontrolling the incident angles of sputtered atoms andresulted in reducing the WCNT.

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H. Okuyama et al. / Physica E 37 (2007) 49–53

53

Fig. 6. SEM images of CNTs grown by (a) TCVD only applying bias without a discharge and (b) PECVD with a glow discharge. An applied bias voltage
was (a) 300 V (discharge current of about 5 mA) and (b) 250 V (discharge current of about 300 mA). The images of (b) were taken at 451 tilt angle. The
insets show magniﬁed images.

In order to obtain VACNTs, electric ﬁeld was applied
during CVD. Fig. 6 shows SEM images of CNTs grown by(a) TCVD only applying bias without a glow discharge and(b) PECVD with a glow discharge accompanied with
discharge current. CNTs with a diameter of ca. 15 nm grewin random directions by the TCVD method. The CNTs
grew selectively near the boundary of the metal pattern.
The effects of the applied bias during TCVD were notobserved. In the case of PECVD method, CNTs with adiameter of ca. 30 nm grew up perpendicular to thesurfaces. The VACNTs were observed all over the pattern.
The VACNTs were achieved only by PECVD method.
When a glow discharge is induced, the voltage drop takes
place mostly in the cathode dark space of which the
thickness is a few hundreds micrometers. On the other
hand, the electric ﬁeld distributes uniformly in the space
between the cathode and the anode in the case of anondischarge. The gap of electrodes–substrates was about
5 mm. Therefore the electric ﬁeld applied in the case of
PECVD is approximately ten times as large as that in
TCVD. CNTs grew up all over the pattern by PECVD
whereas the selective growth was observed by TCVD. This
result suggests that the Ni catalysts were changed into ﬁneparticles by bombardments of ions accelerated in large
electric ﬁeld. As a result CNTs were able to grow even on
the smooth metal patterns.
4. Conclusions
We tried to prepare position-selective growth of CNTs
and VACNTs on patterned metal. The position-selective
growth of CNTs was observed at the edges of the patterned
metal by the TCVD. By controlling not only the position of
catalysts but also the surface morphology of underlayers,
the position selective growth width of CNTs was minimized, approximately 2.6 mm. This width was approximately one-eightieth of the patterned metal, 200 mm. In

order to prepare VACNTs, DC negative bias voltage wasapplied to the substrates. When the glow discharge took
place in the PECVD, VACNTs were successfully synthesized. The VACNTs grew up, however, all over thepatterned metal. Conclusively the selective growth of
CNTs was achieved in the optimum conditions of the
TCVD by controlling the morphology of the Mo underlayers for Ni catalysts.
Acknowledgments
The authors are grateful for the ﬁnancial support of theGrant-in-Aid for the Futaba Foundation. This work was
supported by the Ministry of Education, Culture, Sports,
Science and Technology through a Scientiﬁc Grant-in-Aid(no. 15360173).
References[1] J. Kong, E. Yenilmez, T.W. Tombler, W. Kim, H. Dai,
R.B. Laughlin, L. Liu, C.S. Jayanthi, S.Y. Wu, Phys. Rev. Lett. 87(2001) 106801.
[2] Z. Yao, C.L. Kane, C. Dekker, Phys. Rev. Lett. 84 (2000) 2941.
[3] E.W. Wong, P.E. Sheehan, C.M. Lieber, Science 277 (1997) 1971.
[4] J.H. Hafner, C.L. Cheung, C.M. Lieber, Nature 398 (1999) 761.
[5] R. Stevens, C. Nguyen, A. Cassell, L. Delzeit, M. Meyyappan,
J. Han, Appl. Phys. Lett. 77 (2000) 3453.
[6] T. Arie, H. Nishijima, S. Akita, Y. Nakayama, J. Vac, Sci. Technol.
B 18 (2000) 104.
[7] S. Fan, M.G. Chapline, N.R. Franklin, T.W. Tombler, A.M. Cassell,
H. Dai, Science 283 (1999) 512.
[8] N. Saurakhiya, Y.W. Zhu, F.C. Cheong, C.K. Ong, A.T.S. Wee,
J.Y. Lin, C.H. Sow, Carbon 43 (2005) 2128.
[9] W.I. Milne, K.B.K. Teo, G.A.J. Amaratunga, R. Lacerda,
P. Legagneux, G. Pirio, V. Semet, V.T. Binh, Curr. Appl. Phys. 4(2004) 513.
[10] A.C. Dupuis, Prog. Mater. Sci. 50 (2005) 929.
[11] V.I. Merkulov, A.V. Melechko, M.A. Guillorn, D.H. Lowndes,
M.L. Simpson, Appl. Phys. Lett. 79 (2001) 2970.

Mater. Res. Soc. Symp. Proc. Vol. 901E © 2006 Materials Research Society

0901-Ra22-25-Rb22-25.1

Selective Chemical Vapor Growth of Carbon Nanotubes
for Application of Electronic-Measuring Nano-Probes
Hiroki Okuyama, Nobuyuki Iwata and Hiroshi Yamamoto
College of Science and Technology, Nihon University, Chiba, Japan
ABSTRACT
Selectively grown carbon nanotubes (CNTs) and/or vertically aligned CNTs (VACNTs) on
patterned metal electrodes have been prepared for applications to a novel electronic-measuringnano-probe. We propose newly a position-controlling method of CNTs by controlling not only
a position of catalysts but also the thickness of metal underlayers for the catalysts.

The

selective growth of CNTs was achieved at edges of the metal pattern by thermal chemical vapor
deposition. The width of selectively grown CNTs was minimized, about 2.6 µm. The width
was about 80 times as small as that of the metal pattern. VACNTs were synthesized by aplasma enhanced CVD method in which the high electric field in a cathode dark space wasapplied to the substrates. The VACNTs grew up, however, all over the metal pattern.

INTRODUCTION
Carbon nanotubes (CNTs) have attracted much attention because of many featured
characteristics. High electric conductivity [1] and high permissible current density [2] aresuitable for electronic devices.

CNTs grow selectively from catalyst particles by usingchemical vapor deposition (CVD) [3].

CNTs have a high aspect ratio and their growth direction

is controllable by applying electric field during CVD growth [4]. These characteristics areavailable to fabricate various kinds of nanostructures. CNTs reveal high mechanical strength inspite of high flexibility [5].

Therefore we are aiming at realizing a novel nano-probe with CNTs

which enables to approach it toward contact points on specimen easily without an accurate
control for measuring of electric transport properties. The control for position and direction of
CNTs is necessary in order to prepare the nanoscaled CNTs electronic devices. Furthermore
electrically and mechanically firm contacts between CNTs and metal electrodes are needed.
In this work, we tried to prepare selectively grown CNTs and vertically aligned CNTs
(VACNTs) on patterned metal electrodes for realization of the novel electronic measuringnano-probe. The selectively grown CNTs are prepared by using the two characteristics asfollows. One is that the CNTs grow selectively from fine particle catalysts. The other is that

0901-Ra22-25-Rb22-25.2

we can control the synthesis of the fine particle catalysts by noticing wettability between metalinterfaces. Since wettability of metal interface is high, it is hard to form fine particles asoverlayers on smooth underlayers. When the surfaces of underlayers are rough, overlayered
metals form fine particles.

Therefore we propose newly a position-controlling method of CNTs

by controlling not only a position of catalysts but also a surface morphology of metal
underlayers.

EXPERIMENTAL
In the first place Ni catalysts and Mo underlayers were deposited on quartz substrates
through a metal mask (thickness: 30 µm, pattern width: 200 µm, gap: 100 µm) by RF magnetron
sputtering. A background pressure of the chamber was 1.0 × 10-2 Pa, and the pressure during
deposition was fixed at 2.0 Pa by feeding Ar gas. The growth conditions were 120 °C of
substrate temperature, 50 W of sputtering power, 1 min of sputtering time for Mo underlayers,
and 70 °C, 20 W, 1 min for Ni catalysts, respectively.
CNTs were grown by thermal chemical vapor deposition (TCVD) and dc plasma enhanced
chemical vapor deposition (PECVD) method. The reactor tube of the CVD equipment washeated up to 600 °C. The flow of H2 : Ar (10 : 200 ccm) was introduced into the reactor tube
while heating.

The background pressure was 0.1 Pa.

In the case of a TCVD method, mixture

gas of ethylene (C2H4) : H2 : Ar (20 : 10 : 200 ccm) was introduced into the reactor tube for 30
min for the CNT growth. The total pressure was maintained at 100 kPa. In the case of a
PECVD method, the CNT growth was performed by inducing a dc glow discharge and feeding
mixture gas of C2H4 : H2 : Ar (5 : 50 : 50 ccm). The total pressure was maintained at 2 kPa.
The dc plasma was produced by applying bias voltage of -250 V between a substrate holder andgrounded anodes which apart from the substrates by about 5 mm.
electrodes were connected electrically to the substrate holder.

The patterned metal

RESULTS AND DISCUSSION
Figure 1 shows SEM images of the edge boundary of the patterned metal layers after TCVD
deposition. CNTs grew selectively on just insides from the pattern edges where thickness of
Mo underlayers varied gradually. In the CNT grown region the surfaces of Mo layers wererough and porous. On the contrary CNTs didn’t grow on the mid areas of the pattern where thesurfaces of Mo layers were smooth. By increasing Mo deposition time from 1 min to 3 min, the0901-Ra22-25-Rb22-25.3

(a)

(b)

Figure 1. SEM images of the edge boundary of patterned metal layers after the TCVD
deposition. Sputtering time of the Mo layers was (a) 1 min and (b) 3 min. The upper-right of
the image is the substrate surfaces, the lower-left is the patterned metal layers. CNTs grew inthe narrow region which appeared as the white contrast. The WCNT was about (a) 30 µm and (b)6.3µm. The insets are each the magnified image.
widths of selectively grown CNTs, WCNT decreased from 30 µm to 6.3 µm.
The selective growth of CNTs at the pattern edges was achieved because CNTs grow fromonly fine particle catalysts.

On the smooth surface of Mo underlayers, Ni catalysts don’t

become fine particles due to the high wettability. The rough morphology of Mo underlayers
resulted in fine particles of Ni catalysts. The rough morphology was appeared in very thinregions of Mo due to “shadow effect” of a sputtering method. A schematic shadow effect is
illustrated in figure 2. Sputtered atoms arrive at substrates with various incident angles becauseof scattering by atmospheric gasses.

Then the pattern edges thin down by deposition ofparticles shaded from the mask pattern. Since the rough regions of thicker Mo underlayers
became narrow, the WCNT decreased by increasing Mo thickness.
Aligned fins were installed during deposition to control incident angles of sputtered atoms.
A schematic configuration of the fins and the metal mask is shown in figure 3 (a). And optical
microscopy image of the specimen, which was prepared by installing the fins, is shown in figure
3 (b). The contrast of the pattern was not uniform. Figure 4 shows SEM images of the patternboundary after TCVD by installing the fins during Mo deposition.

The WCNT in every pattern issummarized in figure 5. The gray belts correspond to metal mask patterns, and white belts tothe pattern gaps.

The results show that the WCNT was narrower in “sunny side” than in “shady

Figure 2. Schematic “shadow effect” appeared through the mask pattern. The pattern edges
thin down by deposition of shaded sputtered atoms under the mask edges (indicated by the oval).

0901-Ra22-25-Rb22-25.4

(a)

(b)

Figure 3. A schematic configuration of the aligned fins and the mask (a).
microscopy image of the specimen prepared by installing the fins (b).

(a)

(b)

(c)

(d)

An opticalFigure 4. Typical SEM images of the pattern boundary after TCVD by installing the fins: (a)and (b) show the results on the edges of point D, (c) and (d) on point K. The images of (a) and(c) are the “shady side” and the upper-right of the image is the metal pattern. The images of (b)and (d) are the “sunny side” and the lower-left of the image is the metal pattern. CNTs grew inthe narrow region which appeared as the white contrast. The WCNT was (a) 17 µm, (b) 2.6 µm,(c) 29 µm and (d) 19 µm. The inset shows each the magnified images.
side”. In addition, the WCNT depended on the position of the pattern. The minimum WCNT was
2.6 µm.

This width was about 80 times as small as the pattern width.

The WCNT depended on the position of the pattern because the thickness of metal pattern was
not uniform as shown in figure 3 (b). By increasing the thickness of the metal pattern, the WCNT
decreased as discussed above. Since the effective widths of the shadow effect depended on theincident angle of sputtered atoms, the WCNT obtained in “sunny side” was quite different from

0901-Ra22-25-Rb22-25.5

Figure 5. The WCNT observed at each the points in every patterns prepared. The gray belts
correspond to the metal mask patterns, and white belts to the pattern gaps.
that in “shade side”. An estimated mean free path of sputtered atoms was 6 mm, and wasslightly longer than the fin-substrate distance, 5 mm. The installing of the fins was effective forcontrolling the incident angles of sputtered atoms and resulted in reducing of the WCNT.
Figure 6 shows SEM images of CNTs grown by (a) TCVD only applying bias without a glow
discharge and (b) PECVD with a glow discharge accompanying with discharge current. CNTs
with a diameter of ca. 15 nm grew in random directions by the TCVD method.

The CNTs grew

selectively at near the boundary of the metal pattern. The effects of the applied bias during
TCVD were not observed. In the case of PECVD method, CNTs with a diameter of ca. 30 nm
grew up perpendicular to the surfaces. The VACNTs were observed all over the pattern.
The VACNTs were achieved only by PECVD method. When a glow discharge is induced,
the voltage drop takes place mostly in the cathode dark space of which the thickness is few
hundreds µm. On the other hand, the electric field distributes uniformly in the space between
the cathode and the anode in the case of a non-discharge. The gap of electrodes-substrates was
about 5 mm.

Therefore the electric field applied in the case of PECVD is approximately tentimes as large as that in TCVD. CNTs grew up all over the pattern by PECVD whereas the

(a)

(b)

Figure 6. SEM images of CNTs grown by (a) TCVD only applying bias without a dischargeand (b) PECVD with a glow discharge. An applied bias voltage was (a) -300 V (discharge
current of about 0.005A) and (b) -250 V (discharge current of about 0.3A). The images of (b)were taken at 45° tilt angle. The inset shows magnified images.

0901-Ra22-25-Rb22-25.6

selectively growth was observed by TCVD. This result may suggest that the Ni catalysts werechanged into fine particles by bombardment of ions accelerated in large electric field. As aresult CNTs were able to grow even on the smooth metal patterns.

CONCLUSIONSWe tried to prepare CNTs and VACNTs grown selectively on patterned metal electrodes.
The selectively grown CNTs were observed at the edges of the patterned metal by the thermal
CVD. By controlling not only a position of catalysts but also a surface morphology ofunderlayers, the width of selectively grown CNTs was minimized, about 2.6 µm. The width of
the CNTs was about 80 times as small as that of the pattern metal. In order to prepare VACNTs,
dc negative bias voltage was applied to the substrates.

When the glow discharge took place inthe plasma enhanced CVD, VACNTs were successfully synthesized due to the high electric fieldin the cathode dark space. The VACNTs grew up, however, all over the metal pattern.
Conclusively a selective growth of CNTs was achieved in the optimum conditions of the thermal
CVD by controlling the thickness of the Mo underlayers for Ni catalysits.

It is the future

subject to obtain selectively grown VACNTs on the designated region of metal patterns.

ACKNOWLEDGMENTS
The authors are grateful for the financial support of the Grant-in-Aid for the Futaba
Foundation. This work was supported by the Ministry of Education, Culture, Sports, Scienceand Technology through a Scientific Grant-In-Aid ( No.15360173 ).

REFERENCES
1. J. Kong, E. Yenilmez, T. W. Tombler, W. Kim, H. Dai, R. B. Laughlin, L. Liu, C. S. Jayanthi
and S. Y. Wu, Phys. Rev. Lett. 87, 106801 (2001)
2. Z. Yao, C. L. Kane and C. Dekker, Phys. Rev. Lett. 84, 2941 (2000)
3. A. C. Dupuis, Prog. Mater. Sci. 50, 929 (2005)
4. V. I. Merkulov, A. V. Melechko, M. A. Guillorn, D. H. Lowndes and M. L. Simpson, Appl.
Phys. Lett. 79, 2970 (2001)
5. E. W. Wong, P. E. Sheehan and C. M. Lieber, Science, 277, 1971 (1997)Physica C 463–465 (2007) 1005–1008
www.elsevier.com/locate/physc

Preparations of magnetoelectric thin ﬁlms for superconducting devices
N. Iwata *, T. Asada, K. Nagase, T. Yamada, H. Yamamoto
Department of Electronics and Computer Science, College of Science and Technology, Nihon University, 7-24-1 Narashinodai,
Funabashi-shi, Chiba 274-8501, Japan
Received 30 October 2006; accepted 9 February 2007
Available online 26 May 2007

Abstract
Representative magnetoelectric (ME) Cr2O3 thin ﬁlms were grown on R-cut sapphire and SrTiO3(1 0 0) substrate in order to integratewith high temperature superconductor oxides. The replacement of a base current and stabilization of IC spread for single ﬂux quantum
device is expected. Highly oriented Cr2O3 thin ﬁlm grew on sapphire at 550 C by oﬀ-axis DC–RF hybrid magnetron sputtering methodwith two-in. Cr metal target. The surface was composed of coalesced rectangular grains, the surface roughness Ra of which was 1.9 nm.
Estimated induced magnetic ﬁeld with electric ﬁeld applied of 0.24 MV/cm was 1.0 G around 40 K.
 2007 Elsevier B.V. All rights reserved.
PACS: 68.55. a; 81.15. z; 81.15.Cd
Keywords: Magnetoelectric eﬀect; Cr2O3; Single ﬂux quantum; Thin ﬁlm

1. Introduction
We pay attention to apply the magnetoelectric (ME)materials of single phase compounds to single ﬂux quantum (SFQ) device integrated by high critical temperature(high-Tc) superconductors [1–4]. The ME eﬀect is deﬁned
as the dielectric polarization of a material in an appliedmagnetic ﬁeld or an induced magnetization in an external
electric ﬁeld [1–4]. The performance of the SFQ device isfeatured as a high-speed (over 100 GHz) and low consumption (less than 100 nW/gate) following Si digital devices.
However, increasing integration degree, large amount ofbase current is required, which give rise to uncontrollable
operation due to accompanying magnetic ﬁeld. In addition,
to reduce the ﬂuctuation of critical current (IC) and to stabilize the device operation, working temperature is limited
less than 40 K.

*

Corresponding author. Tel./fax: +81 47 469 5457.
E-mail address: iwata@ecs.cst.nihon-u.ac.jp (N. Iwata).

0921-4534/$ - see front matter  2007 Elsevier B.V. All rights reserved.
doi:10.1016/j.physc.2007.02.036

As shown in the Fig. 1, we propose a novel element how
the ME material is embedded in a SFQ device. Since theME material is grown to generate a base current, the basecurrent is localized and then error operations could be
reduced. It is expected that a base current in a SFQ device
is replaced by the ME materials. The electric power loss
from a current source is excluded. Furthermore, considering critical current IC of the Josephson junction is modulated by the magnetic ﬁeld penetrating the junction area,
IC spread of the junctions could be reduced by the induced
magnetic ﬁeld from the ME ﬁlm. We expect that a largerscale circuit is achievable by the induced and localized magnetic ﬁeld from ME materials integrated with SFQ devices.
In this paper, we have focused on an antiferromagnetic
insulator Cr2O3 which is a representative oxide among ME
single phase compounds. Cr2O3 has the linear ME susceptibility of axx, ayy and azz and the others are zero [1–4].
Magnetic ﬁeld is induced along the direction of electric ﬁeld
applied. Since the ME eﬀect is derived from a magnetic
superexchange interaction, highly oriented ME ﬁlms areclaimed on SrTiO3 (STO), MgO, sapphire, etc., to integratewith high-Tc superconductors.

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N. Iwata et al. / Physica C 463–465 (2007) 1005–1008

Junction

5Fig. 1. Schematic image of a part of the SFQ device with ME material we
propose. Induced magnetic ﬁeld from the ME ﬁlm generates the shielding
current along the high-Tc superconductor. Shielding current could be abase current for operation of SFQ device.

3

10

Cr2O3(3306)

insulator

Cr2O3(2204)

ME

4

10

Cr2O3(1102)

YBCO

Intensity (arbitrary units)

10

2

10

1

10

0

10

20

30

40

50

60

70

80

90

2θ (degrees)

2. Experimental

Fig. 2a shows a XRD pattern of Cr2O3 ﬁlm grown on
R-cut sapphire substrate at 500 C and Fig. 2b is the magniﬁed pattern around Cr2O3(1 1 0 2) Bragg reﬂection at various substrate temperatures from 300 C to 600 C. Bragg
reﬂections from R-plane of Cr2O3 were clearly observed justbelow those of sapphire substrate in 2h. The peak indicatedby open-triangle was probably due to oxygen deﬁcient of
substrate. Any other peaks did not appear in XRD patterns.
The lattice spacing was 0.363 nm similar to that of substrate. As shown in Fig. 2b, peak intensity was the highest
at 500 C with shoulder at right hand side of the peak indicated by arrow. Peaks became decreasing in intensity andbroader in ﬁlms grown 300 C, 400 C and 600 C. Fig. 3
shows the full width at half maximum (FWHM) of rocking
curve for Cr2O3(1 1 0 2) as a function of substrate temper-

10

600ºC

8
6
4

500ºC

400ºC

2

300ºC
24.0

24.2

24.4

24.6

24.8

25.0

2θ (degrees)Fig. 2. XRD patterns of Cr2O3 ﬁlm grown at (a) 500 C on R-cut sapphireand (b) magniﬁed images around Cr2O3(1 1 0 2) at various substratetemperature. Open circles and triangle are come from the substrate. Bragg
reﬂections of Cr2O3 thin ﬁlm were clearly observed.

1.1

FWHM of rocking curve for Cr2O3(1102)

3. Results and discussion

Intensity (arbitrary units)

12

Cr2O3 ﬁlms were grown by oﬀ-axis DC–RF hybrid magnetron sputtering method with two-in. Cr metal target on
the substrates of R-cut sapphire and STO with CeO2 buﬀer
layer. Substrates surface of sapphire was etched by HF acidfor 10 s (HF:H2O = 1:7) and then annealed at 1000 C for
12 h in air in advance. For Cr2O3 growth substrate temperatures, DC current, RF power and sputtering gas were300–630 C, 0.04 A, 100 W and 0.26 Pa (Ar:O2 = 4:1),
respectively. CeO2 buﬀer layer was deposited by on-axis
RF magnetron sputtering method with pressed CeO2 powder at 380 C. The substrate was 0.5 wt.% Nb-doped
STO(1 0 0). A post-annealing was carried out at 1000 C
for 4 h in air prior to growth of Cr2O3.
The X-ray diﬀraction (XRD) 2h h scans from 5 to 90,using Rigaku RAD-C system with Cu Ka1 and Ka2 radiation generated with a tube power of 40 kV and 30 mA,
examined the grown ﬁlms. The diﬀraction area was always
same with divergence slit of 1/2 and sample size of
5 · 10 mm2. The surface morphology was observed in a
dynamic force mode by an atomic force microscopy(DFM, Seiko Instruments Inc. SII: NanoNavi Station).

1.0
0.9
0.8
0.7
0.6
0.5

300

350

400
450
500
550
Substrate Temperature ( C)

600

Fig. 3. The FWHM of rocking curve for Cr2O3(1 1 0 2) as a function ofsubstrate temperature. At 500 C the sharpest peak was obtained.

ature. At 500 C the rocking curve was 0.58 and the
sharpest.
Generally the FWHM of Bragg reﬂections in 2h h
examination depends on a ﬁlm thickness in particular lessthan 500 nm and a ﬂuctuation of a lattice spacing. Since

N. Iwata et al. / Physica C 463–465 (2007) 1005–1008

the thickness of the grown ﬁlm was constant approximately
190 nm in these experiments, the FWHM is a function of
only the ﬂuctuation of the Cr2O3(1 1 0 2) lattice spacing.
Considering the appearance of the shoulder at the righthand side of the Cr2O3(1 1 0 2) Bragg reﬂection for theﬁlm grown at 500 C, which was derived from the separation of the peak by Ka1 and Ka2 of diﬀerent X-ray wave
lengths, the ﬂuctuation of the lattice spacing was reduced
and then highly oriented ﬁlm was obtained. Although a lattice mismatch is +4.01% along the [1 1 2 0] direction and is+4.34% along the [ 1 1 0 1] direction, an expected compressive stress in-plane is seemed to be released form the resultof the same lattice spacing to the bulk value.
Figs. 4 and 5 show the surface images of the Cr2O3 ﬁlms
with the size of 2 lm · 2 lm and line proﬁles of thoseimages. In Table 1, grain sizes as a function of substrate(a) 400ºC, Ra=1.5 nm(b) 500ºC, Ra=2.3 nm

1007

Fig. 5. Line proﬁles related to the index of the surface images of Fig. 4.
Solid arrows indicate the diﬀerence in height at the grain boundary where
coalescence did not take place. Dotted arrows indicate the steps composedof multiples of 1/3 unit cell along the [1 1 0 2] direction inside one grain.

(b-L1)

Table 1Substrate temperature vs. grain size(a-L1)(b-L2)

0 nm

13.6 nm(c) 550ºC, Ra=1.9 nm

0 nm(d) 600ºC, Ra=4.2 nm(c-L2)(d-L1)

16.9 nm

0 nm33.2 nm(e) 630ºC, Ra=2.8 nm(e-L1)

0 nm

Grain size (nm · nm)

400
500
550
600

100 · 100–200
100 · 100–500
200 · 200–600
400 · 400

21.0 nm(c-L1)

0 nm

Sub. temp. (C)

24.3 nmFig. 4. Surface images with 2 · 2 lm2 of Cr2O3 thin ﬁlms grown at (a)400 C, (b) 500 C, (c) 550 C, (d) 600 C and (e) 630 C. Scale bars wereattached below each image. The ﬁlm grown at (c) 550 C showed thesmoothest surface, the roughness Ra of which was 1.9 nm. Rectangulargrains with the size of 200 · 200–600 nm2 coalesced throughout thesurface expect for the grains separated by deep grain boundaries.

temperature are summarized. Rectangular grains weremainly observed with the size of 100 · 100–200 nm2 at
400 C in Fig. 4a. The diﬀerence in height between grains
was approximately 3.5 nm indicated by solid arrows in
Fig. 5a-L1. One grain contained 2 or 3 steps at the surface,indicated by dotted arrows, with the height of 0.3–0.9 nm,
which was integer multiple of 1/3 unit cell along the[1 1 0 2] direction, 0.363 nm. At 500 C larger square
grains with the size of 100 · 100–500 nm2 were observed
and those grains were coalesced in Fig. 4b. In one grain
the steps of 1–3 integer multiples of 1/3 unit cell was also
observed as shown in Fig. 5b-L1 indicated by arrows.
The grains growing from diﬀerent nucleus sites, where coalescence did not occur, had a diﬀerence of 2–10 nm inheight as shown in Fig. 5b-L2 and sometimes deep grainboundaries were formed. The deepest depth of the boundaries was 34 nm. At 550 C as shown in Fig. 4c rectangulargrains became larger and longer along a particular direction and the smoothest surface was obtained with the surface roughness Ra of 1.9 nm. Almost all of the grains
coalesced except for the grains separated by deep grainboundaries as shown in Fig. 5c-L1 indicated by arrows.
The density of deep boundaries decreased compared withthat of Fig. 4b. As indicated by dotted arrows in Fig. 5cL1 and c-L2, steps of 1/3 unit cell were clearly observedin one grain. Film growth looked like 2-dimensional island.
At 600 C although the largest grains were obtained,bunching steps, the value of which was approximately
5 nm, were observed as shown in Figs. 4d and 5d-L1.

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N. Iwata et al. / Physica C 463–465 (2007) 1005–1008

4

Current Density (A/cm2)

10

2

10

0

10

-2

10

-4

10

-6

10
0.00

0.05

0.10

0.15

0.20

0.25

0.30

0.35

0.40

Electric Field (MV/cm)
Fig. 6. I–V characteristic of Cr2O3/CeO2 bilayered ﬁlm on STO substrate.
Below 0.24 MV/cm the ﬁlm was insulating.

The surface smoothness was degraded down to the Ra of
4.2 nm due to the bunching steps. The surface of the terrace
was so ﬂat that even steps of 1/3 unit cell did not appear. In
addition, spherical grains were also observed with the sizeof B40 lm. At 630 C as shown in Figs. 4e and 5e-L1 ﬁlm
3-dimensionally grew without a regularity in grains.
As the substrate temperature is increased, the migration
length increases and the density of nucleus sites as well asof grains decrease. At 400 C where coalescence did nottake place, the surface was covered by many grains andboundaries with the height of 3.5 nm due to the shortmigration length. However at 500 C where start to coalesce but not to obtain suﬃcient migration length to belarge grain, the deep grain boundary was expected to be
introduced. As increasing substrate temperature up to550 C, increased migration length led to larger grain andthen the density of deep grain boundary was reduced. At
500–550 C, the terrace in one grain was composed of several steps of 1/3 unit cell at most, indicating the 2-dimensional island growth. The unit of –O–Cr–O–Cr–O–,
corresponding to the height of 1/3 unit cell, stacks along
the [1 1 0 2] direction in Cr2O3 single crystal with rhombohedral structure. In the highly oriented Cr2O3 ﬁlms, it isreasonable that the integer multiples of 1/3 unit cell appear
because of the repulsive force between O2 ions at theinterface of stacked 1/3 unit cell. Above 600 C despite sufﬁcient migration length to be large grain, Cr2O3 is expected
to be decomposed due to deﬁcient of oxygen, demonstrated
by the results of XRD and the present of spherical grains.
Fig. 6 is a leakage current density as a function of electric ﬁeld (I–V) of a bilayered ﬁlm of Cr2O3/CeO2//

STO(1 0 0) observed at room temperature. The thickness
of Cr2O3 and CeO2 was 220 nm and approximately
10 nm, respectively. The diameter of measurement area
was B0.45 mm. From the XRD result Cr2O3 ﬁlm was oriented along the c-direction normal to the substrate surface.
Leakage current rapidly increased above 0.24 MV/cm indicating that the Cr2O3 ﬁlm works as a ME material below
the electric ﬁeld.
We estimated the induced magnetic ﬁeld tentatively withthe obtained data. Assuming that the ME coeﬃcient of
Cr2O3 along the c-direction around 40 K azz is 10 4 (inCGS – gauss unit) and the demagnetizing ﬁeld is neglected,
induced magnetic ﬁeld is estimated at 1.0 G [1–4]. For
example, to obtain 0.5 mA by a shielding current, whichis representative for base current in SFQ device, with100 lm2 and 100 pH superconducting loop, 5 G is
required. We expect to achieve such the value by optimizing the growth condition of CeO2 for epitaxial growth of
Cr2O3 on it.
4. Conclusion
We propose to apply the ME materials of single phasecompounds to the SFQ device as the replacement of a basecurrent and stabilization of IC spread. The ME material
generates induced magnetic ﬁeld with electric ﬁeld applied.
As a representative ME material, Cr2O3 thin ﬁlms weregrown on R-cut sapphire and STO substrates in order tointegrate with high-Tc superconductors. At 550 C,
0.26 Pa and a ﬂow rate of Ar:O2 = 4:1, highly oriented
Cr2O3 thin ﬁlms grew on sapphire substrate, the surfaceof which was composed of coalesced rectangular grains
with the size of 200 · 200–600 nm2. The surface roughness
Ra was 1.9 nm. Estimated induced magnetic ﬁeld from theresult of Cr2O3/CeO2//STO was 1.0 G around 40 K. To
eliminate the induced magnetic ﬂux, shielding current is
expect to ﬂow superconducting circuit, which is substituted
for a base current of SFQ device by a current source.
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