Cao Z. (Ed.) Thin Film Growth: Physics, materials science and applications

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208 Thin ﬁlm growth

can be an exponential-like function of 1/T. With decreasing temperature,

the number of carriers decreases while the mobility generally increases,

but in different manners; consequently, the resulting electrical resistivity is

exclusively temperature dependent. For narrow band-gap semiconductors

and semimetals, the situation is more complicated than in the conventional

semiconductors like Si or Ge. The carrier distribution function obeys the

Fermi–Dirac statistics which in this case cannot be approximated by the

Maxwell–Boltzmann formula. Moreover, the conduction band is Kane-type

that the effective mass of electron, thus the mobility of electron, changes

with the carrier density (Lovett 1977). The vanishing TCR occurs when

this compound has narrowed its band-gap to be a semimetal for which the

temperature dependence of the number of the carriers becomes rather mild-

mannered in contrast to the situation in a semiconductor of a denite band-

gap. It results from a delicate balance between the opposite changes of the

number of carriers and of the carrier mobility with temperature, and reasonably

it appears only at temperatures below 240 K – at higher temperatures the

increasing number of carriers can easily compensate for the loss in carrier

mobility to give rise to a rapidly decreasing electrical resistivity. The current

discovery, though lacking a clear scenario of detailed microscopic processes

due to the difculty in calculating the band structure for the off-stoichiometric

compounds such as Cu

NPd

0.238

and in performing Hall measurement for some

supplementary information, does demonstrate the possibility that a balanced

change of carrier density and carrier mobility over a wide temperature range

is in principle possible in a class of semimetallic materials. Considering the

unlikelihood of a temperature-independent electrical resistivity for solids,

the current result also bears some signicance for solid state physics.

Resistivity (10

–6

Wm)

1.4

1.3

1.2

1.1

1.0

0 50 100 150 200 250 300

Temperature (K)

(f)

8.15 Continued.

ThinFilm-Zexian-08.indd 208 7/1/11 9:42:10 AM

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209Thin ﬁlm growth for thermally unstable noble-metal nitrides

8.5 Conclusions

By carefully adjusting the working parameters including the applied power, gas

pressure and the ratio of nitrogen in the working gas, and substrate temperature,

thin lms of noble-metal nitrides such as Cu

N and Cu

NPd

can be successfully

prepared by using the reactive magnetron sputtering method. For stoichiometric

N, the lms are found to comprise distinct crystallites in a typical size

of 40–60 nm. The peculiarity of growing noble-metal nitride lms lies in the

fact that the reaction of sputtered metal atoms or clusters with nitrogen occurs

most probably only at the target surface, hence the matching of the density and

energy of nitrogen ions with the sputtering yield for a chosen target material is

the key factor to realizing stoichiometry in the resulting copper nitride lms. A

large ion energy is unfavorable since the higher sputtering yield of the metal

atoms and the weakened combination with energetic nitrogen ions will give

rise to a metal-rich deposit. A low substrate temperature is critical since at

slightly elevated temperatures nitrogen re-emission occurs, which may even

result in microsized protruding features. Semiconducting to (semi)metallic

transition can be observed in doped Cu

N structures wherein the excessive

atoms are expected to sit at the cell centers of the Cu

N lattice. A vanishing

temperature coefcient of resistivity over a temperature of 200 K has been

realized in Cu

NPd

0.238

. Many other surprising properties are anticipated in

the noble-metal nitrides, which await discovery in the future.

8.6 References

Asano M, Umeda K, and Tasaki A (1990), Cu

N thin lm for a new light recording

media, Jpn. J. Appl. Phys. 29, 1985–1986.

Baranov O, Romanov M, Wolter M, Kumar S, Zhong X X, and Ostrikov K (2010),

Low-pressure planar magnetron discharge for surface deposition and nanofabrication,

Physics of Plasma 17, 053509.

Berg S, and Nyberg T (2005), Fundamental understanding and modeling of reactive

sputtering processes, Thin Solid Films 476, 215–230.

Borsa D M, Grachev S, Presura C, and Boerma D O (2002), Growth and properties of

N lms and Cu

N/g-Fe

N bilayers, Appl. Phys. Lett. 80, 1823–1825.

Cao Z X (2001), Electron cyclotron waveresonance plasma assisted deposition of cubic

boron nitride thin lms, J. Vac. Sci. Technol. A 19, 485–489.

Cao Z X (2002), Plasma enhanced deposition of silicon carbonitride lms and property

characterization, Diamond & Relat. Mater. 11, 16–21.

Cao Z X, and Oechsner H (2004), Effect of concurrent N

-ion bombardment on the physical

vapor deposition of nitrides thin lms, J. Vac. Sci. Technol. A 22(2), 321–323.

Chambers S A (2010), Epitaxial growth and properties of doped transition metal and

complex oxide lms, Adv. Mater. 22(2), 219–248.

Cremer R, Witthaut M, Neuschutz D, Trappe C, Laurenzis M, Winkle O, and Kurz H

(2000), Deposition and characterization of metastable Cu

N layers for applications

in optical data storage, Mikro. Acta. 133, 299–302.

Dorranian D, Dejam L, Sari A H, and Hojabri A (2010), Structural and optical properties

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

210 Thin ﬁlm growth

of copper nitride thin lms in a reactive Ar/N

magnetron sputtering system, Euro.

Phys. J.-Appl. Phys. 50, 20503.

Du Y, Ji A L, Ma L B, Wang Y Q, and Cao Z X (2005), Electrical conductivity and

photoreectance of nanocrystalline copper nitride thin lms deposited at low temperature,

J. Crystal Growth 280, 490–494.

Hahn U, and Weber W (1996), Electronic structure and chemical-bonding mechanism of

N, Cui

NPd, and related Cu(I) compounds, Phys. Rev. B 53, 12684–12693.

Hojabri A, Haghighian N, Yasserian K, and Ghoranneviss M (2010), The effect of nitrogen

plasma on copper thin lm deposited by DC magnetron sputtering, IOP Conf. Ser.:

Mater. Sci. Eng. 12, 012004.

Jacobs H, and Zachwieja U (1991), Kupferpalladiumnitride, Cu

N mit x = 0.020

und 0.989, Perowskite mit ‘bindender 3d

-4d

-Wechselwirkung’, J. Less-common

Met. 170, 185–190.

Ji A L, Li C R, Du Y, Ma L B, Song R, and Cao Z X (2005), Formation of rosette pattern

in copper nitride thin lms via nanocrystals gliding, Nanotech. 16, 2092–2095.

Ji A L, Huang R, Du Y, Li C R, and Cao Z X (2006a), Growth of stoichiometric Cu

thin lms by reactive magnetron sputtering, J. Crystal. Growth 295, 79–83.

Ji A L, Li C R, and Cao Z X (2006b), Ternary Cu

NPd

exhibiting invariant electrical

resisitivity over 200 K, Appl. Phys. Lett. 89, 252120.

Ji A L, Du Y, Li C R, and Cao Z X (2007), Formation of symmetrical relief features in

nanocrystalline copper nitride thin lms, J. Vac. Sci. Technol. B 25 (1), 208–211.

Juza R, and Hahn H (1939), Copper nitride (in German), Z. Anorg. Allg. Chem. 241,

172–178.

Kim K J, Kim J H, and Kang J H (2001), Structural and optical characterization of

N lms prepared by reactive RF magnetron sputtering, J. Crystal Growth 222,

767–772.

Lovett D R (1977), Semimetals & Narrow-bandgap Semiconductors, Pion Limited,

London.

Ma G M, Alejandro M, and Noboru T (2004), Ab initio total energy calculations of copper

nitride: the effect of lattice parameters and Cu content in the electronic properties,

Solid Stat. Sci. 6, 9–14.

Maruyama T, and Morishita T (1996), Copper nitride and tin nitride thin lms for write-

once optical recording media, Appl. Phys. Lett. 69, 890–891.

Moreno-Armenta M G, Martínez-Ruiz A, and Takeuchi N (2004), Ab initio total energy

calculations of copper nitride: the effect of lattice parameters and Cu content in the

electronic properties, Solid State Sciences 6, 9–14.

Musil J, Baroch P, Vlcek J, Nam K H, and Han J G (2005), Reactive magnetron sputtering

of thin lms: present status and trends, Thin Solid Films 475, 208–218.

Norskov J K, and Hammer B (1995), Why gold is the noblest of all the metals, Nature

376, 238–240.

Nosaka T, Yoshitake M, Okamoto A, Ogawa S, and Nakayama Y (2001), Thermal

decomposition of copper nitride thin lms and dots formation by electron beam

writing, Appl. Surf. Sci. 169–170, 358–361.

Reddy K V S, Reddy A S, Reddy P S, and Uthanna S (2007), Copper nitride lms

deposited by dc reactive magnetron sputtering, J. Mater. Sci. – Materials in Electronics

18, 1003–1008.

Terada S, Tanaka H, and Kubota K (1989), Heteroepitaxial growth of Cu

N thin lms,

J. Crystal Growth 94, 567–568.

Wang D Y, Nakamine N, and Hayashi Y (1998), Properties of various sputter-deposited

Cu–N thin lms, J. Vac. Sci. Technol. A 16, 2084–2092.

ThinFilm-Zexian-08.indd 210 7/1/11 9:42:10 AM

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211

Growth of graphene layers for thin films

B. H. Hong and H. R. Jeon, Sungkyunkwan University,

South Korea

Abstract: graphene has received considerable attention due to its fascinating

physical properties such as quantum electronic transport, a tunable band gap,

extremely high mobility, high elasticity and electromechanical modulation.

This chapter describes the direct synthesis of large-scale graphene lms

using chemical vapour deposition on thin nickel layers, and presents two

different methods of patterning the lms and transferring them to arbitrary

substrates. Furthermore, we detail the roll-to-roll production and wet

chemical doping of predominantly monolayer 30-inch graphene lms grown

by chemical vapour deposition onto exible copper substrates.

Key words: graphene growth, chemical vapour deposition, half-integer

quantum Hall effect, roll-to-roll production.

9.1 Introduction

In order to be used as nano-electronic devices, it is essential to grow graphene

on a large scale. Among methods for graphene synthesis, graphene growth

by chemical vapour deposition (CVD) on ni and Cu has the advantage of

growing graphene at large scale with uniform thickness. This chapter deals

with graphene growth on ni and Cu by CVD. First we describe large-scale

graphene lm synthesis using CVD on nickel layers, their pattern and their

transfer to arbitrary substrates.

In next section we detail the roll-to-roll

production and wet chemical doping of predominantly monolayer 30-inch

graphene lms grown by CVD onto exible copper substrates.

Further, we

detail how graphene electrodes were incorporated into a fully functional

touchscreen panel device capable of withstanding high strain.

9.2 Large-scale pattern growth of graphene ﬁlms

for stretchable transparent electrodes

9.2.1 Direct synthesis of large-scale graphene ﬁlms

Ni lms were e-beam evaporated on SiO

/Si substrates, and thermally

annealed. After exposing the ni surface to a CH

, H

and Ar mixture in

atmospheric conditions, the system was rapidly cooled to room temperature.

This fast cooling rate is critical in suppressing formation of multiple layers

ThinFilm-Zexian-09.indd 211 7/1/11 9:42:28 AM

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212 Thin ﬁlm growth

and for separating graphene layers efciently from the substrate in the later

process.

A scanning electron microscope (SEM) image of graphene lms on a thin

nickel substrate shows clear contrast between areas with different numbers

of graphene layers (Fig. 9.1a). Transmission electron microscope (TeM)

images (Fig. 9.1b) show that the lm consists mostly of less than a few layers

of graphene. After transfer of the lm to a silicon substrate with a 300 nm

thick Sio

layer, optical and confocal scanning Raman microscope images

were made of the same area (Fig. 9.1c,d).

The brightest area in Fig. 9.2(d)

corresponds to monolayers, and the darkest area is composed of more than

ten layers of graphene. Bilayer structures appear to predominate in both

TeM and Raman images for this particular sample, which was prepared

from 7 minutes of growth on a 300 nm thick nickel layer. It is found that

the average number of graphene layers, the domain size and the substrate

coverage can be controlled by changing the nickel thickness and growth

process, thus providing a way of controlling the growth of graphene for

different applications.

Atomic force microscope (AFM) images often show the ripple structures

caused by the difference between the thermal expansion coefcients of nickel

and graphene (Fig. 9.1c).

It is believed that these ripples make the graphene

lms more stable against mechanical stretching,

making the lms more

expandable. Multilayer graphene lms are preferable in terms of mechanical

strength for supporting large-area lm structures, whereas thinner graphene

lms have higher optical transparency. It is found that a ~300 nm thick

nickel layer on a silicon wafer is the optimal substrate for the large-scale

CVD growth that yields mechanically stable, transparent graphene lms to

be transferred and stretched after they are formed, and that thinner nickel

layers with a shorter growth time yield predominantly mono- and bilayer

graphene lm for microelectronic device application.

9.2.2 Transfer processes for large-scale graphene ﬁlms

The transfer of CVD-derived graphene lms to non-specic substrates is

enabled by FeCl

etching of the underlying Ni lm. Use of buffered oxide

etchant (BOE) or hydrogen uoride solution removes silicon dioxide layers,

so the patterned graphene and the nickel layer oat together on the solution

surface. After transfer to a substrate, further reaction with Boe or hydrogen

uoride solution completely removes the remaining nickel layers.

A dry-transfer process has been developed for the graphene lm using

a soft substrate such as polydimethylsiloxane (PDMS) stamp.

Here the

PDMS stamp is rst attached to the CVD-grown graphene lm on the

nickel substrate (Fig. 9.2d). The nickel substrate can be etched away using

FeCl

as described above, leaving the adhered graphene lm on the PDMS

ThinFilm-Zexian-09.indd 212 7/1/11 9:42:28 AM

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213Growth of graphene layers for thin ﬁlms

(a)

(c)

(b)

(d)

5 µm

2 µm

5 µm

4–5 layers Bilayer

0.34 nm

>10 layers

3 layers

Intensity (a.u.)

l = 532 nm

>4 layers

3 layers

Bilayer

Monolayer

1500 2000 2500

Raman shift (cm

–1

)

(e)

9.1 Various spectroscopic analyses of large-scale graphene ﬁlms

grown by CVD. (a) SEM images of as-grown graphene ﬁlms on

thin (300 nm) nickel layers and thick (1 mm) Ni foils (inset). (b)

TEM images of graphene ﬁlms of different thickness. (c) An optical

microscope image of the graphene ﬁlms transferred to a 300 nm

thick silicon dioxide layer. The inset AFM image shows typical rippled

structures. (d) A confocal scanning Raman image corresponding to

(c). The number of layers is estimated from the intensities, shapes

and positions of the G-band and 2D-band peaks. (e) Raman spectra

(532 nm laser wavelength) obtained from the corresponding coloured

spots in (c) and (d). a.u. = arbitrary units. (From Ref. [1].)

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214 Thin ﬁlm growth

substrate (Fig. 9.2e). By using the pre-patterned nickel substrate (Fig. 9.2c),

we can transfer various sizes and shapes of graphene lm to an arbitrary

substrate. This dry-transfer process turns out to be very useful in making

large-scale graphene electrodes and devices without additional lithography

(a) (b) (c)

(e)

(d)

(f) (g)

(h)

2 cm

Stamping Patterned graphene

9.2 Transfer processes for large-scale graphene ﬁlms. (a) A

centimeter-scale graphene ﬁlm grown on Ni(300 nm)/SiO

(300 nm)/Si

substrate. (b) A ﬂoating graphene ﬁlm after etching the nickel layers

in 1M FeCl

aqueous solution. After the removal of the nickel layers,

the ﬂoating graphene ﬁlm can be transferred by direct contact with

substrates. (c) Various shapes of graphene ﬁlms can be synthesized

on top of patterned nickel layers. (d, e) The dry-transfer method

based on a PDMS stamp is useful in transferring the patterned

graphene ﬁlms. After attaching the PDMS substrate to the graphene

(d), the underlying nickel layer is etched and removed using

FeCl

solution (e). (f) Graphene ﬁlms on the PDMS substrates are

transparent and ﬂexible. (g, h) The PDMS stamp makes conformal

contact with a silicon dioxide substrate. Peeling back the stamp (g)

leaves the ﬁlm on a SiO

substrate (h). (From Ref. [1].).

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215Growth of graphene layers for thin ﬁlms

processes (Fig. 9.2f–h). Microscopically, these graphene lms often show

linear crack patterns with an angle of 60°C or 120°C, indicating a particular

crystallographic edge with large crystalline domains. In addition, the Raman

spectra measured for graphene lms on nickel substrates show a strongly

suppressed defect-related D-band peak. This D peak grows only slightly

after the transfer process (Fig. 9.1e), indicating overall good quality of the

resulting graphene lm. Further optimization of the transfer process with

substrate control makes possible transfer yields approaching 99%.

9.2.3 Optical and electrical properties of the graphene

ﬁlms

For the macroscopic transport electrode application, the optical and electrical

properties of 1 ¥ 1 cm

graphene lms were measured by ultraviolet-visible

spectrometer and four-probe Van der Pauw methods, respectively (Fig. 9.3a,b).

The transmittance is measured using an ultraviolet-visible spectrometer after

transferring the oating graphene lm to a quartz plate (Fig. 9.3a). In the

visible range, the transmittance of the lm grown on a 300 nm thick nickel

layer for 7 minutes is ~80%, a value similar to those found for previously

studied assembled lms.

8,9

Because the transmittance of the graphene layer

is ~2.3%,

this transmittance can be increased to ~93% by further reducing

the growth time and nickel thickness, resulting in a thinner graphene lm.

Ultraviolet/ozone etching is also useful in controlling the transmittance under

ambient conditions (Fig. 9.4a, upper inset). Indium electrodes were deposited

on each corner of the square (Fig. 9.3a, lower inset) to minimize contact

resistance. The minimum sheet resistance is ~280W per square, which is

~30 times smaller than the lowest sheet resistance measured on assembled

lms.

8,9

The values of sheet resistance increase with the ultraviolet/ozone

treatment time, in accordance with the decreasing number of graphene layers

(Fig. 9.3a)

For microelectronic application, the mobility of the graphene lm is

critical. To measure the intrinsic mobility of a single-domain graphene

sample, the graphene samples are transferred from a PDMS stamp to a

degenerate doped silicon wafer with a 300 nm deep thermally grown oxide

layer. Monolayer graphene samples were readily located on the substrate

from optical contrast

and their identication was subsequently conrmed

by Raman spectroscopy.

electron beam lithography was used to make

multi-terminal devices (Fig. 9.3b, lower inset). notably, the multi-terminal

electrical measurements showed that the electron mobility is ~3750 cm

–1

at a carrier density of ~5 ¥ 10

–2

(Fig. 9.3b). For a high magnetic

eld of 8.8T, the half-integer quantum Hall effect (Fig. 9.3b) is observed,

indicating that the quality of CVD-grown graphene is comparable to that of

mechanically cleaved graphene.

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216 Thin ﬁlm growth

Transmittance (%)

83.7% UV for 6 h

80.7% UV for 4 h

79.1% UV for 2 h

76.3% Initial

at 550 nm

LIV

(kW per

square)

1.2

1.0

0.8

0.6

0.4

0.2

0 1 2 3 4 5 6

Time (h)

Tr (%)

400 600 800 1000 1200

Wavelength (nm)

(a)

–60 –40 –20 0 20 40 60

(V)

(b)

5 µm

Magnetoresistance (kW)

Resistance (kW)

–5

–10

–15

–60 0 60

(V)

9.3 Optical and electrical properties of graphene ﬁlms. (a)

Transmittance of the graphene ﬁlms on a quartz plate. The

discontinuities in the absorption curves arise from the different

sensitivities of the switching detectors. The upper inset shows the

ultraviolet (UV)-induced thinning and the consequent enhancement

of transparency. The lower inset shows the changes in transmittance,

, and sheet resistance, R

, as functions of ultraviolet illumination

time. (b) Electrical properties of monolayer graphene devices

showing half-integer quantum Hall effect and high electron

mobility. The upper inset shows a four-probe electrical resistance

measurement on a monolayer graphene Hall bar device (lower inset)

at 1.6 K. A gate voltage, V

, is applied to the silicon substrate to

control the charge density in the graphene sample. The main panel

shows longitudinal (R

) and transverse (R

) magnetoresistances

measured in this device for a magnetic ﬁeld B = 8.8T. The monolayer

graphene quantum Hall effect is clearly observed, showing the

plateaux with ﬁlling factor u = 2 at R

= (2e

/h)

–1

and zeros in

. (Here e is the elementary charge and h is Plank’s constant.)

Quantum Hall plateaux (horizontal dashed lines) are developing for

higher ﬁlling factors. (c) Variation in resistance of a graphene ﬁlm

transferred to a ~0.3 mm thick PDMS/PET substrate for different

distances between holding stages (that is, for different radii). The

left inset shows the anisotropy in four-probe resistance, measured

as the ratio R

, of the resistances parallel and perpendicular to

the binding direction, y. The right inset shows the bending process.

(d) Resistance of a graphene ﬁlm transferred to a PDMS substrate

isotropically stretched by ~12%. The left inset shows the case in

which the graphene ﬁlm is transferred to an unstretched PDMS

substrate. The right inset shows the movement of holding stages

and the consequent change in shape of the graphene ﬁlm. (From Ref.

[1].)

In addition to the good optical and electrical properties, the graphene lm

has excellent mechanical properties when used to make exible and stretchable

electrodes (Fig. 9.3c,d).

The foldability of the graphene lms, transferred

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217Growth of graphene layers for thin ﬁlms

to a polyethylene terephthalate (PET) substrate (thickness, ~100 mm) coated

with a thin PDMS layer (thickness, ~200 mm; Fig. 9.3c), is evaluated by

measuring resistances with respect to bending radii. The resistances show

little variation up to the bending radius 2.3 mm (approximate tensile strain

of 6.5%), and are perfectly recovered after unbending. notably, the original

resistance can be restored even for the bending radius of 0.8 mm (approximate

tensile strain of 18.7%), exhibiting extreme mechanical stability in comparison

with conventional materials used in exible electronics.

The resistance of graphene lms transferred to pre-strained and unstrained

PDMS substrates was measured with respect to uniaxial tensile strain ranging

from 0 to 30% (Fig. 9.3d). Similar to the results in the folding experiment,

the transferred lm on the unstrained substrate recovers its original resistance

after stretching by ~6% (Fig. 9.3d, left inset). However, further stretching

often results in mechanical failure. Thus, it was tried transferring the lm

to pre-strained substrates

to enhance the electromechanical stabilities by

creating ripples similar to those observed in the growth process (Fig 9.1c,

inset). The graphene transferred to a longitudinally strained PDMS substrate

does not show much enhancement, owing to the transverse strain induced

by Poisson’s effect.

To prevent this problem, the PDMS substrate was

isotopically stretched by ~12% before transferring the lm to it (Fig. 9.3d).

Surprisingly, both longitudinal and transverse resistance (R

and R

) appear

stable up to ~11% stretching and show only one order of magnitude change

at ~25% stretching. It is supposed that further uniaxial stretching can change

the electronic band structure of graphene, leading to the modulation of sheet

resistance. These electrochemical properties thus show graphene lms to be

not only the strongest

but also the most exible and stretchable conducting

transparent materials so far measured.

Resistance (W)

1st 2nd 3rd

Stretching cycles

0 3 6 0 3 6 0 3 6

Stretching (%)

Stable

Resistance (W)

Resistance (kW)

0 5 10 15 20 25 30

Stretching (%)

(d)

Flat 3.5 2.7 2.3 1.0 0.8 Flat

Bending radius (mm)

(c)

Anisotropy, R

0.0 0.4 0.8 1.2

Curvature, k

(mm

–1

)

Bending Recovery

9.3 Continued

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