Inverse StranskiâKrastanov Growth in Single-crystalline Sputtered Cu

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Inverse Stranski–Krastanov Growth in Single-crystalline Sputtered Cu Thin Films for Wafer-Scale Device Applications Seunghun Lee, Ho-Yeol Park, Su Jae Kim, Hyangsook Lee, Ik-Jae Lee, Chae Ryong Cho, Eunha Lee, Se-Young Jeong, and Young Hee Lee ACS Appl. Nano Mater., Just Accepted Manuscript • DOI: 10.1021/acsanm.9b00673 • Publication Date (Web): 30 Apr 2019 Downloaded from http://pubs.acs.org on May 1, 2019

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ACS Applied Nano Materials

Inverse Stranski–Krastanov Growth in Singlecrystalline Sputtered Cu Thin Films for Wafer-Scale Device Applications Seunghun Lee†,‡, Ho-Yeol Park†,#, Su Jae Kim&, Hyangsook Lee§, IkJae Lee¶, Chae Ryong Cho‖, Eunha Lee§,*, Se-Young Jeong#,&,ǂ,*, and Young Hee Lee,* ‡Center

for Nanophysics and Advanced Materials, and Department of

Materials Science and Engineering, University of Maryland, College Park, Maryland 20742, United States #Department

of Cogno-mechatronics Engineering, Pusan National

University, Busan 46241, Republic of Korea &Crystal

Bank Research Institute, Pusan National University, Miryang, 50463, Republic of Korea

§Analytical

Engineering Group, Samsung Advanced Institute of

Technology (SAIT), Samsung Electronics Co., Suwon 16678, Republic of Korea

ACS Paragon Plus Environment

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ACS Applied Nano Materials 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

¶Pohang

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Accelerator Laboratory, Pohang University of Science and Technology, Pohang 37673, Republic of Korea

‖Department

of Nano Fusion technology, Pusan National University, Busan 46241, Republic of Korea

ǂDepartment

of Optics and Mechatronics Engineering, Pusan

National University, Busan 46241, Republic of Korea Center

for Integrated Nanostructure Physics, Institute for Basic

Science, Sungkyunkwan University Suwon 16419, Republic of Korea [email protected] (Y. H. Lee) [email protected] (S.-Y. Jeong) [email protected] (E. H. Lee) KEYWORDS: sputtering system, copper thin film, single-crystalline thin film growth, grain-free, incoherent twin boundary

ABSTRACT

The

single-crystalline

copper

(Cu)

thin

film

is

a

platform

substrate for the growth of numerous materials and has been a significant

issue

for

the

scientific

community.

The

primary

concern is the inevitable presence of stacking faults and twin


2

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boundary

formation

structures.

Here,

crystalline

Cu(111)

conventional

in we

inherent

report

thin

a

films

sputtering

face-centered-cubic method

on

an

deposition.

for

growing

(FCC) single-

Al2O3

substrate

The

desired

using growth

configuration is induced by hidden incoherent twin boundaries (HITBs) embedded during the early growth stages. Two possible FCC stacking orders of Cu atoms, A-B-C-A and A-C-B-A rotated by 60˚, give

rise

to

diffraction

HITBs

phi-scan

between

islands,

mapping.

Such

as

confirmed

islands

merge

by

X-ray

every

three

layers, triggering layer-by-layer growth that subsequently leads to an inverse Stranski–Krastanov growth mode. Single-crystalline Cu(111)

thin

transmission

film

growth

electron

is

confirmed

microscopy

and

by

high-resolution

electron

backscatter

diffraction mapping. Our approach paves the way for mass production of single-crystalline metal thin film and thus leads to substantial advanced research and electronic device application.

INTRODUCTION Metal thin films are key components of microelectronic devices and

have

been

used

in

a

variety

of

applications

such

as

metallization, optical coatings, catalyst, and plamonics.1,2 Copper (Cu),

in

particular,

has

numerous

advantages,

such

as

high

electrical and thermal conductivities, as well as a low cost.3 However, Cu has crucial weaknesses with regard to its practical


3


application,

mainly

its

vulnerability

Page 4 of 31

to

oxidation.

Oxygen

penetrates the Cu surface through grain boundaries, resulting in a degradation of the physical properties of Cu and its performance reliability.4-6

Cu

has

recently

been

revisited

due

to

its

availability as a practical substrate for graphene synthesis via low carbon solubility7-10; however, the inevitable grain boundary formations and rough surfaces in Cu foil remain as obstacles to achieving high-quality graphene.11 Therefore, the growth of singlecrystalline Cu thin films has been a significant issue, given that this film is used as a platform substrate for the growth of numerous materials. There have been various attempts to understand the epitaxial growth mode of Cu thin films, and to improve the crystallinity of the film structure.12-28 The face-centered cubic (FCC) structure of Cu is inherently susceptible to stacking-fault and

twin-boundary

formations,

in

addition

to

point

and

line

defects, during thin film growth.29-31 In this regard, the formation of twin boundaries and their effects on the physical properties of Cu thin films have been intensively investigated.32-44 Wafer-scale single-crystal-like Cu thin films have been grown by precisely controlling the plasma in a conventional sputtering process employing a single-crystal Cu target45; this approach allows precise control over the size and homogeneity of sputtered specimens. The extremely flat surface and high crystallinity of Cu have allowed successful constructions of flexible transparent


4

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conducting crystalline realize

electrode monolayer

ultra-flat,

via

patterning46

micro-size

graphene.47 grain-free

Here, Cu

our

films

and

objective

with

low

singlewas

to

synthesis

temperature, which meets the practical demand for mass production. We observed inverse Stranski–Krastanov (ISK) growth at the atomic scale by resolving the formation of unique hidden incoherent twin boundaries (HITBs) in wafer-scale and single-crystalline sputtered Cu thin films.

EXPERIMENTAL SECTION Cu thin films were fabricated using an rf-sputtering method. The single-crystal

Cu

target

used

in

the

sputtering

process

was

prepared from a single-crystal Cu ingot grown by the Czochralski method,

which

was

supported

by

Crystal

Bank,

Pusan

National

University, Korea. A 2-inch disk-shaped target was prepared by electron discharge machining. The optimal growth parameters for Cu thin films grown on a 2-inch Al2O3 (0001) substrate were secured by adjusting various parameters, such as the temperature, rfpower, and working pressure. We first adjusted the rf-power and working pressure, which provide stable plasma during deposition, and the distance between the Al2O3 substrate and the single crystal Cu target, and the rf-power were further adjusted to set the growth rate ~ 1 Å/s. The film thickness was fixed to be 200 nm for


5


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measurements. The optimal growth temperature was investigated in a systematic manner (from room temperature to 200 ˚ C). The base pressure was fixed at ~10-6 Torr, and the working pressure was 10-3, adjusted by high-purity argon gas (99.999 %). The growth rate was 1 Å/s. The fabricated thin films were kept in an air atmosphere; over a 2–3 year time period, no changes in surface quality were visible to the naked eye. The θ-2θ x-ray diffraction (XRD) measurements were performed using a PANalytical Empyrean Series 2 instrument equipped with a Cu-Kα source (40 kV, 30 mA). Data were collected within the range 20˚ < 2θ < 90˚, with a step size of 0.0167˚ and a dwell time of 0.5 s per point in all cases. Synchrotron ϕ-scan XRD measurements were performed at the BL1D of the Pohang Light Source (PLS), Korea, operating at 3.0 GeV with a maximum storage current of 300 mA. The X-ray beam was focused using a toroidal mirror and monochromatized to 8.0495 keV (0.154027 nm) using a double-bounce silicon (Si) (111) monochromator. Atomic force microscopy (AFM) measurements were carried out using an XE-100 system (Park Systems, Inc.). Scanning

electron

microscopy

(SEM),

electron

backscattering

diffraction (EBSD), and pole figure measurements were performed with a Zeiss SUPRA40 VP with scanning electron microprobe. Highresolution transmission electron microscopy (HR-TEM) analyses were performed using an FEI Titan3 G2 60-300 equipped with double


6

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aberration

correctors

(image

and

probe)

and

a

monochromator

operating at an acceleration voltage of 200 kV. Transmission electron microscopy (TEM) samples were prepared with a focused ion beam (Helios450F1; FEI).

RESULTS AND DISCUSSION ISK growth mode is observed in the single-crystalline Cu thin films fabricated using a modified sputtering system equipped with a single-crystal Cu target (see the supplementary information for experimental details, including photographs of the Cu thin films (Figure S1)). Stranski–Krastanov (SK) growth can be described as initially two-dimensional layer-by-layer growth and subsequent three-dimensional

island

growth

(Frank–van

der

Merwe

mode

+

Volmer–Weber mode, respectively), as determined by the competition of cohesive forces between adatoms and the surface.48,49 Figure 1 depicts the ISK growth mode of HITBs experimentally observed. As the surface plane of the Al2O3(0001) substrate is analogous to the cubic close packing structure of the Cu(111) plane, with a small lattice mismatch of ~7 %,13 we consider layers A–C of the Cu FCC structure above the basal plane (Figure 1a). Initially, due to the lower interfacial tension energy of the Al2O3 substrate compared to that of the Cu vapors (i.e., the sputtered specimen), Cu vapors are led to form clusters (i.e., islands) to minimize the total


7


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free energy. Therefore, the thin film growth initially follows the three-dimensional

Volmer-Weber

growth

mode.49

Under

epitaxial

growth, the island should have one of two possible FCC stacking orders, ABCABC⋯ and ACBACB⋯, as depicted in Figure 1a. When two islands (ABC⋯/ACB⋯) with different stacking orders form next to each other, Cu atoms begin to fill the lateral spaces between them (Figure 1b). As Cu atoms continue to build up on the basal plane, a small vacant gap less than the size of an atom is inevitably left over. To insert one Cu atom, the gap requires twice the dspacing of the (110) plane (i.e.,

2a= 5.08 Å), which is wider

than the spatial distance between layers B and C above the basal layer (3.69 Å), denoted by “v”. The adatoms in the second layer denoted by “i" are placed on top of the vacant gap in the first layer with sufficient space (5.16 Å) at the boundary between the two stacking configurations, i.e., Σ3{112} ITB (Figure 1c). The atoms in the third layer relax fully to form atoms identical to the basal plane by merging two islands (Figure 1d). ITBs in the kink region are repeated several times to resolve strain due to the lattice misfit and recover to an ideal boundary-free lattice (Figure 1e) which triggers subsequent layer-by-layer growth with the formation of a Σ3{111} coherent twin boundary (CTB) (the blue thick line in Figure 1f). We first investigated an optimal growth temperature in systematic manner, and show that the growth temperature of 190 ˚C is one of


8

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the key parameters to achieve the single-crystalline Cu thin film. Figure 2a shows the θ−2θ XRD patterns of the sputtered Cu thin films deposited at room temperature (RT), 100, 180, and 190 ˚ C; the patterns exclusively exhibited a Cu(111) peak, regardless of the deposition temperature. Notably, the deposition temperature had a significant influence on the peak intensity (Figure 2a). We emphasize that the 100-nm-thick Cu thin film deposited at 190 ˚C had a full-width half-maximum (FWHM) of 0.025˚, much narrower than that of Al2O3 (0.048˚) (Figure S2). From AFM measurements, the line profile of the film deposited at 190 ˚ C in the bottom panel of Figure 2e revealed a root mean square (RMS) surface roughness of less than 0.2 nm (Figure 2e), which was markedly reduced (by ~15 times) compared with that (3.2 nm) of the film deposited at RT. To further confirm the smooth surface observed in the sputtered Cu thin film deposited at 190 ˚C, we measured the RMS surface roughness of Cu thin films of varying thicknesses (20–600 nm); our results revealed a narrow thickness dependence of 0.4 ± 0.2 nm (Figure S3). Normal direction (ND) electron backscattering diffraction (EBSD) measurements were conducted to validate the single crystal quality of the sputtered Cu (111) thin films at various temperatures (Figure 3a–d); the radio-frequency (rf) dependence is shown in Figure S4. At RT, the EBSD mapping in the (111) direction was


9


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dominant; however, inverse pole figures (IPFs) revealed a blurred (111) and additional (001) or (101). Similar {100} pole figures were visible with six primary spots accompanied by various, less obvious speckles. The spots other than those in the (111) direction indicated the presence of polycrystalline regions in the film (Figure

S5

shows

additional

IPF

data).

As

the

temperature

increased, the EBSD spots other than the (111) direction became less prominent. The sole spots associated with the (111) direction were distinct at 190 ˚C. The {100} pole figure holds an ideal cubic three-fold symmetry; however, the six-fold symmetry of the {100} pole figure in Figure 3d implies the presence of twin boundaries in the sputtered Cu thin film. To elucidate the rotation angle of the twin boundary, we examined the distribution of the misorientation angle in the EBSD (rolling direction (RD) mode) mappings, as shown in Figure 4a–d, in which the rotation angle between grains is indicated by color. The occurrence of grain boundary angles in the sputtered Cu thin film deposited at RT was widely distributed (lower panel). The grain boundary occurrence at 60˚ became more prominent at high deposition temperatures,

reaching

nearly

100

%

at

190

˚ C,

whereas

the

occurrence of grain boundaries at angles other than 60 ˚

was

negligible (Figure 4d). Two grains rotated by 60˚ were identified, each with an ABCABC⋯ (white) and ACBACB⋯(gray) stacking order in


10

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three-fold symmetry; the boundary lines of the misorientation angle (red) formed at the interface between the two stacking orders (Figure 4d and Figure S6). Two regions with ABC and ACB stacking orders are schematically depicted in Figure 4e. The six-fold symmetry observed in the ϕscan synchrotron XRD data on Cu (111) (Figure 4f) can be explained in terms of the overlap of the two three-fold symmetries of ABC and ACB stackings, which is in good agreement with the six-fold symmetry from EBSD pole measurements (the bottom figure of Figure 3d). Figure S7 shows scanning electron microscopy (SEM) images of the Cu(111) film deposited at various temperatures. High-resolution

high-angle

annular

dark-field

scanning

transmission electron microscopy (HAADF-STEM) observations were conducted near the Al2O3 interface in an attempt to identify, at the atomic scale, the thin film growth mode in the Cu thin film prepared at 190 ˚C. A kink in the rectangular box near the Al2O3 substrate was observed by two adjacent stacking orders in different [111] orientations, denoted by A and B (Figure 5a). Fast Fourier transform (FFT) electron diffraction patterns for regions A and B (Figure 5b and c) clearly show the twin boundary with respect to the [110] zone axis. The folded diffraction pattern for the full regions of A and B confirmed the twin boundary, with precise twofold symmetry of the spots (Figure 5d).


11


Page 12 of 31

To clarify the nature of the twin boundary, an atomic-resolution image was extracted from inverse fast Fourier transform (IFFT) data,

as shown in Figure

5d.

Three

layers

were continuously

repeated from the bottom substrate, with two stacking orders of ABCABC⋯ to the left of the kink and ACBACB⋯ to the right (Figure 5e); the two stacking orders assumed a disordered arrangement in the kink region. Remarkably, the two emergent stacking orders recovered from the kink disorder and merged to form epitaxial single-crystalline Cu layers by the 22nd layer (from the basal substrate), or equivalently, at ≈ 6 nm from the substrate. The magnified region in Figure 5f shows the disorders at the kink site as vacancies and strained atoms, resulting in an ITB with two stacking orders. Given that disorders at a kink site are embedded by subsequent epitaxial layers, this feature is defined as a HITB which is not the ITBs observed in Fig. 4d since such thin layers (≈ 6 nm) are unlikely to produce strong diffraction. To investigate the film growth mode of sputtered Cu thin film (because real-time observation is not available), we measured cross-sectional TEM images of Cu thin films prepared with different deposition times (Figure 6), which directly evinces the suggested growth mechanism depicted in Figure 1. From the cross-sectional TEM image of a sample prepared with a short deposition time (note that the growth rate that we estimated from a thick and complete thin film form was ≈ 1 Å/s), one can identify the initial growth


12

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mode: if the thin film growth indeed follows two-dimensional layerby-layer growth even in the early growth stage (i.e., Frank-van der Merwe mode), the cross-sectional image is expected to show a complete thin-film form with a uniform thickness. However, as shown in Figure 6a, the sputtered Cu specimens form Cu islands (6-nm height with 5-nm spacing, i.e., not a complete thin film form) on the Al2O3 substrate after 20-s deposition, which is indicative of three-dimensional island growth, namely Volmer-Weber growth mode in the early stage. The TEM image of the sputtered Cu specimens after 30-s deposition still shows a similar island form on the substrate, but the spacing between islands becomes narrower and more obscure (Figure 6b), and the complete thin-film form (6 nm thick) is observed for the sputtered Cu thin film after 50-s deposition

(Figure

6c).

Undoubtedly,

the

series

of

TEM

measurements reveal that the spaces between the initially formed Cu islands are filled by additional sputtering, and the islands become merged in this stage. It is worthy of note that the thickness of the primary thin-film form is consistent with the thickness of 22 layers in [111] FCC stacking (see Figure 5e), which accounts for the presence of the incoherent twin boundaries near the substrate. Once sputtered Cu forms a complete thin film form, the interfacial tension initially present between Al2O3 and Cu vapor is significantly reduced, and the two-dimensional layer-bylayer growth mode can be triggered, which is supported by the


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extremely smooth surface of the sputtered Cu thin films (rms roughness ≈ 4 Å; note that the thickness of a layer in [111] FCC stacking is expected to be ≈ 2 Å) regardless of the film thickness (see Figure S3, Supplementary information). A homogenous lattice pattern is indeed observed in Figure 6d, which is a cross-sectional TEM image of a 100 nm thick sputtered Cu thin film above the interface region between the substrate and thin film. More detailed information on the sputtered Cu thin film and a comparison with polycrystalline

Cu

thin

film

can

be

found

in

Supplementary

information (Figure S10–S12).

CONCLUSIONS The primary achievement of our work is the wafer-scale singlecrystalline Cu thin film growth by using a conventional sputtering system with the single-crystal Cu target; the very low deposition temperature (=190 ˚C) is highly adequate for practical industrial application, which will also boost many advanced researches. ISK growth was observed at the atomic scale by resolving the unique HITB

formation,

which

is

the

key

mechanism

in

the

single-

crystalline Cu thin film growth. While gold films have been the film of choice in the fields of metallic nanostructure and nanopatterning of meta-materials, surface plasmonics, and nano-optical device

applications,2,50-53

the

ability

to

produce

high-quality


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single-crystalline Cu thin films, as presented in this research, provides a low-cost, reliable alternative to Au films for longterm use.

ASSOCIATED CONTENT Supporting Information. The following files are available free of charge. Additional experimental details and figures (PDF)

AUTHOR INFORMATION Corresponding Author [email protected] (Y. H. Lee) [email protected] (S.-Y. Jeong) [email protected] (E. H. Lee) Author Contributions ‡These authors contributed equally. Notes The authors declare no competing financial interest. Acknowledgement This research was supported by Basic Science Research Program


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Page 16 of 31

through the National Research Foundation of Korea (NRF) funded by the Ministry of Science, ICT & Future Planning (No.NRF2017R1A2B3011822) and by the Institute for Basic Science (IBSR011-D1).

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Atomistic

Structure

of

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Cu(111)/Α-


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Figure 1.

Schematic diagram of the growth mechanism of epitaxial

copper (Cu) thin films: several stages of a) early formation of Cu islands with designated stacking layers A–C on the basal plane, adding b) the first layer with a vacant gap between layers B and C. c) Second layer with an Σ3{112} incoherent twin boundary (ITB) along the vacancy line (thick orange line). d) Third layer relaxed to the basal plane, and e) the kink denoted by long-dashed lines and repeated stacking, consequently resulting in inverse Stranski– Krastanov (ISK) growth and f) layer-by-layer (LbL) growth (from red dashed line, indicated by arrow), at which the Σ3{111} coherent twin boundary forms (blue thick line).


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Figure 2. X-ray diffraction (XRD) data and surface morphologies of Cu thin films grown at various temperatures: a) θ−2θ XRD patterns. b)-e) Surface morphologies from atomic force microscopy (AFM) images. Each lower panel shows the line profile across the surface.


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Figure

3.

a)-d),

Electron

backscattering

diffraction

(EBSD)

orientation maps (top), inverse pole figures (IPFs) (middle), and {100}

pole

figures

(bottom)

of

sputtered

Cu(111)

thin

films

deposited at different temperatures. The EBSD measurements were carried out along the normal direction (ND).


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Log intensity (a.u.)

1 2 3 a b c d 4 RT 100 °C 180 °C 190 °C 5 6 7 8 9 10 11 3 μm 3 μm 3 μm 3 μm 12 13 14 Min Max Min Max Min Max Min Max % % % % 15 2 25 8.7 2 25 8.6 2 25 0.0 2 25 0.0 16 28.8 25 59 25 59 16.1 25 59 3.3 25 59 0.4 17 18 59 60 61.6 59 60 74.5 59 60 96.2 59 60 99.4 19 20 e f 21 22 23 24 25 26 27 28 29 -180 -120 -60 0 60 120 180 30 (deg.) 31 32 33 4 34 | PresenceFigure of ITB line mapping in Cu thin in film. a-d, EBSD maps Cu showing 4. Presence ofsingle-crystal ITB line mapping single-crystal thin the 35 ntation angle distribution of sputtered Cu thin films deposited at various temperatures. The distribution film: a)–d) rolling direction (RD) mode EBSD maps showing the 36 undary rotation angles between min and max (degrees) in the lower panels from each boundary line ma 37 misorientation of sputtered films possible orientations of the firstangle layer indistribution the hexagonal basal plane (001). Cu Left thin hexagon represents the 38 [111] direction; the red rhombuses denote a 60° of rotation. on structure along 39 of a cubic deposited at the various temperatures. The distribution the f, ϕ-scan 40 otron XRD pattern of therotation Cu (111) angles plane. Diffraction peaks marked as alternating arrows correspond to t boundary between min and max (degrees) in the lower 41 ographic orientations of squares in d. 42 panels from each boundary line map. e) Two possible orientations 43 44 of the first layer in the hexagonal basal plane (001). Left hexagon 45 represents the projection of a cubic structure along the [111] 46 47 direction; the red rhombuses denote a 60 ˚ rotation. f) ϕ-scan 48 49 synchrotron XRD pattern of the Cu (111) plane. The film deposited 50 51 at 190 ˚C were used for the ϕ-scan XRD. Diffraction peaks marked 52 53 as alternating arrows correspond to the crystallographic 54 orientations of squares in (d). 55 56 57 58 59 ACS Paragon Plus Environment 60

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Figure 5. Hidden incoherent twin boundary (HITB) formation near the bottom substrate. a) A high-resolution high-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) image of the kink in the A and B regions consisting of different stacking orders deposited at 190 ˚C. b)-d), The corresponding fast Fourier transform (FFT) images of regions A and B and the full region in (a). e) Inverse FFT (IFFT) image of the square region marked by the white dashed square in (a). The white lines denote the atomic layer shared by regions A and B. From the red line, a single stacking ordering is observed. f) Magnified ITB of the square region in (e). Vacancies and strained atoms are represented as white dotted circles and red solid circles, respectively.


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Figure 6. ISK growth of an epitaxial single-crystalline Cu thin film: Cross-sectional HAADF-STEM images of the sputtered Cu thin film deposited at 190 ˚ C with varying deposition times; a) Cu islands at 20 s. b) Coalescence stage of two islands at 30 s, and c) ISK growth followed by layer-by-layer growth at 50 s, creating a ≈6-nm-thick Cu thin film. d) Epitaxial single-crystalline growth was observed at 30 min in the ≈100-nm-thick film. The insets show the corresponding magnified images.


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Inverse StranskiâKrastanov Growth in Single-crystalline Sputtered Cu

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Inverse StranskiâKrastanov Growth in Single-crystalline Sputtered Cu