Vapor-liquid-solid growth of SnO2 nanowires utilizing alternate source supply and their photoluminescence properties

    Vapor-liquid-solid growth of SnO 2 nanowires utilizing alternate source supply and their photoluminescence properties Tomoaki Terasako, Kohki Kohno, Masakazu Yagi PII: DOI: Reference:

S0040-6090(17)30650-8 doi:10.1016/j.tsf.2017.05.053 TSF 36193

To appear in:

Thin Solid Films

Received date: Revised date: Accepted date:

24 March 2017 15 May 2017 17 May 2017

Please cite this article as: Tomoaki Terasako, Kohki Kohno, Masakazu Yagi, Vaporliquid-solid growth of SnO2 nanowires utilizing alternate source supply and their photoluminescence properties, Thin Solid Films (2017), doi:10.1016/j.tsf.2017.05.053

This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

ACCEPTED MANUSCRIPT

IP

alternate source supply and

T

Vapor-liquid-solid growth of SnO2 nanowires utilizing

SC R

their photoluminescence properties

Graduate School of Science and Engineering, Ehime University, 3 Bunkyo-cho,

MA

a

NU

Tomoaki Terasakoa,*, Kohki Kohnob, Masakazu Yagic

Matsuyama, Ehime, 790-8577, Japan Faculty of Engineering, Ehime University, 3 Bunkyo-cho, Matsuyama, Ehime,

D

b

c

TE

790-8577, Japan

National Institute of Technology, Kagawa College, 551 Koda, Takuma-cho,

CE P

Mitoyo, Kagawa, 769-1192, Japan

AC

Vapor-liquid-solid (VLS) growth of tin dioxide (SnO2) nanowires (NWs) on the c-plane sapphire substrate coated with the Au film (5-30 nm in thickness) was performed by atmospheric-pressure CVD utilizing alternate source supply (ASS) of tin (Sn) and water (H2O). X-ray diffraction measurements and scanning electron microscope (SEM) observations revealed the successful growth of SnO2 NWs by the ASS technique. The ASS technique was found to be effective for suppressing the enhancement of the NWs average diameter caused by the increase in growth temperature (Tg). In the cycle number range from 300 to 800, the NWs average diameter was almost independent of cycle number, indicating that the radial growth due to vapor-solid (VS) growth was suppressed. PL spectra of the NWs were dominated by a 1

ACCEPTED MANUSCRIPT

broad orange band (OB) emission (~1.9 eV) associated with oxygen vacancy (VO)

T

and/or interstitial tin atom (Sni). Mirror reflection symmetry between the PL and PLE

IP

spectra with a very large Stokes shift indicated a strong coupling between the deep-level

SC R

defect and phonons in the optical transition process. The e nhancement of the OB emission with increasing Tg is probably due to the increase in the density of the

SnO2; nanowire; X-ray diffraction; scanning electron microscope;

MA

Keywords:

NU

structural defects composed of VO and/or Sni.

D

photoluminescence

TE

*Corresponding author. Tel: +81 89 927 9789; fax: +81 89 927 9790.

AC

CE P

E-mail address: [email protected]

2

ACCEPTED MANUSCRIPT

1. Introduction

T

Tin dioxide (SnO2) with a rutile structure has a band gap energy of -3.6 eV [1]

IP

and high transmittance in the visible region, and exhibits n-type conduction with

SC R

resistivity values ranging from 10-3 to 10-4 cm [2]. These features enable SnO2 to apply transparent conducting films [3,4], solar cells [5,6], gas-sensing devices [7,8], phosphors [9-11] and field effect transistors [12]. Especially, for the gas-sensing

NU

applications, the use of the nanostructured materials has been expected to be effective

MA

for enhancing the sensing performance [13-15]. Various quasi-one dimensional (1D) SnO2 nanostructures, such as nanobelts (NBs), nanofibers (NFs), nanowires (NWs) and

thermal oxidation [20], carbothermal reduction [21],

TE

thermal evaporation [16-19],

D

nanorods (NRs), have been synthesized by a variety of growth techniques, such as

chemical vapor deposition (CVD) [22-24], solution process [25,26], and so on. Among

CE P

these growth techniques, a vapor-liquid-solid (VLS) growth mechanism using a metal catalyst is widely used. The advantages of the VLS growth mechanism are as follows:

AC

(1) the NWs can be selectively grown on the area coated with the catalyst particles (or the catalyst film), and (2) the diameter of the NW can be controlled by the diameter of the catalyst particle (or the thickness of the catalyst film). However, the controllability of the diameters of the NWs are declined by the following reasons: ① the enhancement of the diameters of the catalyst particles caused by the coalescence among the neighboring catalyst particles during the rise in substrate temperature and ② the contribution of the film growth on the NW’s side surface due to a vapor-solid (VS) growth mechanism [27-29]. We have reported the successful VLS growth of the quasi 1D structures of the binary oxides, ZnO, CdO, and β-Ga2O3, by atmospheric-pressure CVD (AP-CVD) 3

ACCEPTED MANUSCRIPT

using metal (Zn, Cd and Ga) and water (H2O) as source materials so far [30-34]. The

T

main reasons for the use of H2O as the oxygen precursor are its safety and good

IP

controllability. Moreover, there is a possibility that the H atom involving H2O acts as a

SC R

donor impurity [35,36]. It was found that the ZnO and CdO NRs grown at higher substrate temperatures exhibited tapered shapes, resulting from the competition between the axial growth due to the VLS growth mechanism and the radial growth due to the VS

NU

growth mechanism [31,32].

MA

When both the Zn and O atoms are supplied simultaneously to the growth surfaces of the NWs, the film growth process due to the VS growth mechanism

D

proceeds. Therefore, the considerable way to avoid the VS growth mechanism is to

TE

shorten the time when the adsorbed Zn and O atoms coexist on the same growth surface as much as possible. From another point of view, the catalyst particles can be regarded

CE P

as the reservoir of Zn atoms by forming the Au-Zn alloy droplets. Even if the Zn supply from the outside is stopped, the Zn atoms reserved in the catalyst particles are

AC

continuously provided only to the growth front of the VLS growth, i.e. the interface between the catalyst particle and the tip of the NW. Under the alternate supply of Zn and O sources, therefore, the VS growth mechanism is suppressed, while the VLS growth mechanism proceeds with lower growth rate. In our previous paper, the successful growth of ZnO NRs by the alternate source supply of Zn and H2O and its effectiveness in suppressing the tapering behavior have been confirmed [31]. In this paper, we will explore the adaptability of the VLS growth utilizing the alternate source supply of Sn and H2O to the diameter controlled growth of SnO2 NWs and their photoluminescence properties.

4

ACCEPTED MANUSCRIPT

2. Experimental

T

2.1. Sample preparation

IP

Tin (Sn) powder and water (H2O) were used as Sn and oxygen (O) precursors,

SC R

respectively. A piece of c-plane sapphire (-Al2O3) wafer (Techno Chemics) with the gold (Au) film deposited by DC sputtering method was used as the substrate material. The thickness of the Au film was changed in the range from 5 to 30 nm. The AP-CVD

NU

system used in this study is schematically shown in Fig. 1. The AP-CVD system had

MA

two temperature zones. One of the two temperature zones, which was upstream one, was used for heating both the substrate and the aluminum boat containing Sn powder

D

(denoted hereafter by “Sn boat”). The growth reactor was a quartz tube (1,000 mm in

TE

length and 35 mm in inner-diameter). An inner-quartz tube with an inlet tube for a nitrogen (N2) carrier gas on its upstream edge was placed in the growth reactor. Both the

CE P

substrate and the Sn boat were enclosed in the inner tube. The substrate was set up at the downstream of the Sn boat. In this paper, the temperature for heating both the substrate and the Sn boat is named “growth temperature” and denoted by “Tg”. The growth

AC

temperature Tg was changed in the range from 800 to 925 C. Water was vaporized in its own vaporizer kept at 54 C. The vapor of H2O was transported into the growth reactor by the N2 carrier gas. Both the N2 carrier gaseous for Sn and H2O were supplied through completely separated paths. The N2 carrier gas for H2O and the N2 purge gas flew in the different paths, but were introduced into the growth reactor through the same inlet port. To switch the gas flow to either the supply side or the exhaust side, each path has a set of a normally open-type and a normally closed-type air operated valves. Opening and closing operations of these air operated valves were done using the compressed air supplied via the solenoid valves. The solenoid valves were controlled by the electrical 5

ACCEPTED MANUSCRIPT

signals from the rely board in conjunction with a computer. The order of source supply

T

was as follows:

IP

Sn (3 s) → Purge N2 (10 s) → H2O (3 s) → Purge N2 (10 s),

SC R

where the periods of the electrical signals sent to the corresponding solenoid valves are given in the parentheses. The periods mentioned above were determined empirically by referring the experimental results for the VLS growth of the ZnO NRs utilizing the

NU

alternate source supply technique [31]. The number of repeating times of this sequence,

MA

denoted by “cycle number” was changed in the range from 300 to 800. The flow rates of the carrier gaseous for Sn and H2O and the N2 purge gas were 40 SCCM, 30 SCCM and

D

112 SCCM, respectively. Detailed study on the periods and flow rates of the N2 carrier

TE

gas and the N2 purge gas will be required for achieving higher shape controllability of the NWs. This will probably be a future problem. In this paper, the usual “simultaneous

AC

respectively.

CE P

source supply” and the “alternate source supply” are abbreviated to “SSS” and “ASS”,

2.2. Characterization The crystal structure of the samples was examined by X-ray diffraction (XRD) method using the Cu K radiation with a conventional θ-2θ goniometer (Rigaku RINT2100). Surface morphologies were observed by scanning electron microscope (SEM: Hitachi S-3100H). Photoluminescence (PL) and photoluminescence excitation (PLE) spectra were taken by a spectrofluorometer equipped with a 450 W Xe lamp and two double-grating monochromators (HORIBA Fluorolog-3 Model FL-3-22) at room temperature. Some of the PL measurements were carried out under the excitation of the 325 nm line from a He-Cd laser (Kinmon, IK3052R-BR, Output: 5mW). In this case, 6

ACCEPTED MANUSCRIPT

PL from the sample was detected by a high resolution fiber multichannel spectrometer

T

(Ocean Optics HR2000, Grating: 600 lines blazed at 400 nm, Detector: 2048 pixel

IP

linear CCD array). For photoacoustic (PA) measurements, a Xe lamp (Wacom,

SC R

KXL500F) in conjunction with a spectrometer (Ritsu Oyo Kogaku MC-20L) was used as the excitation light source. The PA cell was composed of a closed vessel with a small channel in which a microphone was attached. The output signal was amplified by a

NU

lock-in amplifier (NF Circuit Design Block 5610B). The modulation frequency of the

MA

light was 14 Hz.

D

3. Results and discussion

TE

3.1. Growth temperature dependence

Fig. 2 shows the XRD patterns of the samples grown by the AP-CVD utilizing the

CE P

ASS under the different conditions. All the XRD peaks can be indexed to the planes of the rutile structure of SnO2 except for -Al2O3 (006), Au (111) and Au (200) peaks. No

AC

characteristic peaks peculiar to the secondary phases, such as Sn and SnO, were observed. This result indicates the successful growth of SnO2 crystallites. The diffraction angles of SnO2(200) and Au(111) peaks are close each other, so that these peaks cannot be resolved completely. Both the XRD pattern of the sample grown using the 30 nm thick Au film at Tg=925 C and that grown using the 20 nm thick Au film at Tg=900 C are dominated by SnO2(200) peak. On the XRD patterns of the samples grown using the 5 and 20 nm thick Au films at the same Tg=825 C, the intensity of SnO2(111) peak is stronger than that of SnO2(200) peak. The epitaxial relationship between the c-plane sapphire and the film is SnO2(100)//Al2O3(0001) with SnO2[100]//Al2O3<1120> [37-39]. Therefore, the appearance of the dominant 7

ACCEPTED MANUSCRIPT

SnO2(200) peak implies the possibility of the epitaxial growth of the NWs. Several

T

researchers obtained epitaxial VLS grown SnO2 NWs at the growth temperatures of 900

IP

C [40,41]. The critical temperature for the epitaxial growth of the NWs must be in the

possibility of the epitaxial growth of the NWs.

SC R

range from 825 to 900 C. However, further study is needed for confirming the

Fig. 3 shows the surface SEM images of the samples whose XRD patterns are

NU

shown in Fig. 2. All the NWs were grown only on the area covered with the Au film.

MA

Some of the NWs had rounded particles on their tips, suggesting that they were grown via the VLS growth mechanism. On the SEM image of the sample grown using the 30 nm thick Au film at Tg=925 C and that grown using the 20 nm thick Au film at Tg=900

TE

D

C, many straight NWs with several μm in length are aligned relatively in particular directions. The samples grown using the 20 and 5 nm thick Au films at the same Tg=825

CE P

C are composed of the bent and tangled NWs. Fig. 4 shows the histograms of the NWs diameters obtained from the surface SEM images shown in Fig. 3. The NWs grown

AC

using the Au film with the largest thickness of 30 nm at the highest Tg of 925 C are distributed in the range from 200 to 550 nm. With decreasing Au film thickness and Tg, the distribution of the NW diameter shifts towards the smaller side overall. The diameters of the NWs grown using the 20 nm thick Au film at Tg=825 and 900 C are distributed in the range from 100 to 400 nm, and the considerable amount of NWs has the diameter smaller than 300 nm. The diameter distribution of the NWs grown using the 5 nm thick Au at Tg=825 C exhibits a relatively symmetric distribution centering 250-300 nm. Fig. 5 shows the variations of the NWs average diameters Davs plotted as a function of 1000/Tg for the Au film with the thicknesses of 5, 10, 20 and 30 nm. In our 8

ACCEPTED MANUSCRIPT

previous paper, the growth temperature (Tg) dependence of the NWs average diameter

 ,  

SC R

 E Dav  exp   A  kT  B g

IP

T

Dav was approximated by the following equation:

(1)

NU

where EA is the activation energy and kB is the Boltzmann constant [34]. The

MA

experimental data points are distributed in the range from 127 to 320 nm which are approximately 10 to 25 times larger than the Au film thicknesses. The broken line was

D

obtained by substituting the activation energy EA of 0.3eV into Eq (1) and fitted to the

TE

experimental data points. Generally, the thicker the thickness of the catalyst film is, the larger the NW diameter is [42]. However, no clear correlation is found between the Au

CE P

film thickness and the diameter of the NW. For the β-Ga2O3 NWs grown on the Au/c-plane sapphire substrates by the SSS,

AC

the activation energy EA increased from 2.03 to 2.95 eV as the growth time increased from 120 to 300 min. For the SnO2 NWs grown on the Au (30-50 nmthick)/SiO2/Si(100) substrates by the SSS, the activation energy EA was estimated to be 1.25 eV. The activation energy EA of 0.3 eV obtained for the SnO2 NWs grown by the ASS is approximately one-tenth to a quarter of those for the SnO2 and β-Ga2O3 NWs grown by the SSS, indicating that the ASS is effective for suppressing the increase in NW diameter induced by the increase in Tg. In our previous paper, it was found that for the β-Ga2O3 NWs grown at Tg=875 C using the 20 nm thick Au catalyst film, the diameters of the Au catalyst particles observed on the tips of the NWs were larger in comparison with the thickness of the 9

ACCEPTED MANUSCRIPT

catalyst film and widely distributed from 70 to 230 nm [34]. Therefore, it seems likely

T

that the dispersion of the distribution of the experimental data points around the broken

IP

line reflects those of the diameters of the Au catalyst particles formed from the Au

SC R

film during the temperature rising.

3.2. Cycle number dependence

NU

For the SSS, the NWs average diameter Dav increases exponentially with the

MA

increase in growth time t and can be expressed approximately by the following relation:

(2)

TE

D

Dav  exp rt 

where r is defined as a reaction rate constant. The reaction rate constants for the

CE P

β-Ga2O3 NWs grown on the Au (30 nm thick)/c-plane sapphire substrates at Tg=850, 875 and 925 C by the SSS were estimated to be 4.1×10-3, 5.5×10-3 and 7.5×10-3 min-1,

AC

respectively [34]. In the same manner, the reaction constant of 4.0×10-3 min-1 was obtained for the SnO2 NWs grown on the Au (30-50 nm thick)/SiO2/Si(100) substrates at Tg= 900 C by the SSS, which is close to those for the β-Ga2O3 NWs grown by the SSS mentioned above. Therefore, it is speculated that the reaction rate constant for the metal oxide NWs grown by our AP-CVD under the SSS condition is typically in the order of 10-3 min-1. In Fig. 6, the NWs average diameter Dav is independent of cycle number for the SnO2 NWs grown by the ASS. Regardless of the difference in the thickness of the Au film, the Davs are scattered around 250 nm. Once the VLS growth mechanism begins to work, the Au catalyst droplets are lifted up from the substrate surface. As a consequence, 10

ACCEPTED MANUSCRIPT

the Au catalyst droplets can no longer grow bigger by the migration followed by the

T

coalescence. Therefore, it seems reasonable that the reaction rate constants obtained for

IP

the NWs grown by the SSS are governed by the film growth on the side surfaces due to

SC R

the VS growth mechanism. Because the cycle number for the ASS can be interpreted as the growth time for the SSS, the cycle number-independent Dav observed for the SnO2 NWs grown by the ASS provides an experimental evidence for the suppression of the

3.3. Photoluminescence properties

MA

NU

VS growth mechanism.

D

Fig. 7a shows the PL and PLE spectra for the NWs grown under the different

TE

growth conditions by the ASS. All the PL spectra are dominated by an orange band (OB) emission with a peak at 1.9 eV (652 nm). The NWs grown using the 5 and 30 nm

CE P

thick Au films at Tg=925 C exhibit a near-band-edge (NBE) emission at ~3.6 eV other than the OB emission. It was confirmed that the NBE emission tended to be observed

AC

on the NWs grown at higher Tgs, indicating that the increase in Tg contributes to the improvement of the crystalline quality of the NWs. For all the NWs, the PL and PLE spectra are roughly mirrored. The PLE peak appears at ~4.2 eV, which is extremely higher than the band gap energy of SnO2. It was reported that the band gap energy of the SnO2 nanoparticles (NPs) increased drastically from ~3.6 to ~4.4 eV when the NP size decreased from 25 to 4 nm, indicating the contribution of the quantum size effect [43]. As mentioned above, the average diameters of the NWs are larger than 100 nm. Therefore, the quantum size effect is unlikely to be the origin of the appearance of the PLE peak of 4.2 eV. The difference between the PL and PLE peaks, so-called Stokes shift, is estimated to be about 2.3 eV, suggesting a strong phonon coupling in the optical 11

ACCEPTED MANUSCRIPT

transition. As shown in Fig. 7b, the PA spectra for the NWs grown using the 5 nm thick

T

Au film at Tg=875 and 925 C are composed of a wide absorption band spread over a

IP

whole visible region. On both the PA spectra, the main peak appears at ~4.2 eV, which is

SC R

in accordance well with PLE peak energy mentioned above. This result strongly suggests that the excitation and/or recombination process associated with the OB emission is accompanied by the non-radiative process.

NU

The recombination of carriers at deep-level defects usually results in a broad PL

MA

band due to strong electron-phonon coupling. The concept of configuration coordinate (CC) model is often employed to describe the excitation and relaxation processes of the

D

carrier strongly coupled with phonons. In this model, the total potential energy of the

TE

carrier and surrounding atoms (electron-phonon system) is given by a quadratic function (parabola) of the configuration coordinate [44-46]. Fig. 7c shows a schematic CC

CE P

diagram for the OB emission, involving three parabolas corresponding to the ground state and two excited states with the different equilibrium coordinates. The electron

AC

transition is much faster than the motion of the lattice atoms. When the optical excitation is done by the use of the light of 4.2 eV, therefore, the parabola of the total potential energy moves vertically upward from the ground state to the excited state without changing the equilibrium coordinate (the excited state immediately after the electron transition is denoted by “Excited state 1”), corresponding to the absorption process (transition A→B). After that, the lattice atoms surrounding the deep-level defect are rearranged so that the total potential energy is minimized, the so-called lattice relaxation. The excess energy generated during the lattice relaxation is emitted in the form of phonons together with the shift of the equilibrium coordinate from Q0 to Q1 (transition B→C). The excited state after the lattice relaxation is denoted by “Excited 12

ACCEPTED MANUSCRIPT

state 2” as distinguished from “Excited state 1”. When the electron-phonon system at

T

the point C on the excited state 2 returns to the point D on the ground state, a photon

IP

with a photon energy of 1.9 eV is emitted (emission). The difference in photon energy

SC R

between the optical absorption (4.2 eV) and emission (1.9 eV) corresponds to the Stokes shift of 2.3 eV. On the ground state, the electron-phonon system relaxes from the point D to the point A by emitting several phonons (transition D→A). Besides this, there is a

NU

non-radiative route from the point C to the point A via the intersection E of the parabola

MA

for the excited state 2 and that for the ground state (transition C→E→A) proceeds simultaneously with the radiative transition C→D. The main peak at 4.2 eV on the PA spectra is probably due to the phonons generated by the non-radiative transitions B→C,

TE

D

D→A and C→E→A. It is worth noting that the excitation energy of 4.2 eV exceeds the band gap energy of SnO2 (3.5-3.6 eV). Similar behavior was also observed for the

CE P

Cu-related 1.01 eV band in n-type InP [47] and the defect CAs-Oi in AlxGa1-xAs [48]. This result suggests the contribution of the transitions of an electron from the valence

AC

band to the conduction band and following non-radiative capture of a hole by a deep acceptor to the OB emission [49]. Fig. 8a shows the PL spectra of the SnO2 NWs grown using the 30-nm thick Au film at Tg=875, 900 and 925 C. In contrast to the PL spectra taken under the excitation by the 292 nm line, the PL spectra taken under the excitation by the 325 nm line from the He-Cd laser exhibit two peaks at 600 and 660 nm. Moreover, the NBE emission became extremely weak or disappeared. The shape of the PLE spectrum for the NBE emission differed remarkably from that for the OB emission. The intensity of the PLE spectrum for the NBE emission decreased rapidly with the wavelength on the longer wavelength side than 292 nm. Therefore, it is hard to observe the NBE emission on the 13

ACCEPTED MANUSCRIPT

PL spectrum taken under the excitation by the 325 nm line. Details on the excitation and

T

radiative recombination processes of both the NBE and OB emissions will be presented

IP

elsewhere. When Tg increases from 875 to 925 C, the PL intensity increases

SC R

approximately by a factor of six. By referring the reported absorption spectrum [50], the penetration depths of the excitation lights with the wavelengths of 292 and 325 nm are estimated approximately to be 33 and 200 nm, respectively. Taking into account that the

NU

NWs average diameter was typically about 250 nm, the PL taken under the excitation by

MA

the 292 nm comes mainly from the surface region of the NW. Therefore, possible reasons for the change in the spectral shape of the OB emission depending on the

D

wavelength of the excitation light source are the contribution of the band bending due to

TE

the depletion layer caused by the dangling bonds and/or the adsorbed atoms and molecules on the side surface [51] and the difference in the stress applied to the

CE P

deep-level defects [52].

Fig. 8b shows the variations of the PL intensity at 640 nm, which is the middle

AC

point of the two peaks, as a function of Tg for the SnO2 NWs grown using the Au films with the different thicknesses of 5, 20 and 30 nm. Regardless of the difference in Au film thickness, all the measuring points are distributed around the broken line. Moreover, it was also confirmed that the PL intensity at 640 nm was independent of the cycle number. Several researchers reported that the radiative recombination center for the OB emission was related to an oxygen vacancy (VO) and/or an interstitial Sn atom (Sni) [23, 53-60]. Taking into account these reports, it can be inferred that increase in Tg promotes the formation of the high density of the deep-level defects involving VO and/or Sni. To discuss the origin, and excitation and recombination process of the OB emission

14

ACCEPTED MANUSCRIPT

quantitatively, PL and PLE measurements involving their temperature dependences are

IP

T

also planned for the NWs with the different diameters.

SC R

4. Conclusion

In summary, vapor-liquid-solid (VLS) growth of SnO2 nanowires (NWs) on the c-plane sapphire substrate with the Au catalyst film (5-30 nm in thickness) was

NU

performed by the atmospheric-pressure CVD (AP-CVD) utilizing alternate source

MA

supply (ASS) of Sn and H2O. Main conclusions obtained in the present study are as follows: (1) X-ray diffraction measurements and SEM observations revealed the

D

successful growth of SnO2 NWs by the ASS technique, (2) the NWs average diameter

TE

increased slightly with increasing Tg (decreasing 1000/Tg) and its activation energy was estimated to be 0.3 eV, which was small in comparison with those for the SnO2 and

CE P

β-Ga2O3 NWs grown by the usual simultaneous source supply, (3) the NWs average diameter was independent of cycle number, (4) the experimental results of (2) and (3)

AC

suggest that the ASS is effective for reducing the influence of vapor-solid (VS) growth, (5) photoluminescence (PL) spectra from the NWs were dominated by the orange band (OB) emission (~1.9 eV) related to the deep-level defect strongly coupled with phonons, and (6) the increase in Tg resulted in the increase in PL intensity of the OB emission, indicating the increase in the density of the structural defect composed of an oxygen vacancy and/or an interstitial tin atom.

Acknowledgments The authors would like to thank Mr. Akira Miyata for his help with the experiments. This work was supported by JSPS KAKENHI Number JP26390029. 15

ACCEPTED MANUSCRIPT

References

T

[1] R. Summitt, J. A. Marley, N. F. Borrelli, The ultraviolet absorption edge of stannic

IP

oxide (SnO2), J. Phys. Chem. Solids 25 (1964) 1465-1469.

SC R

[2] S. Das, V. Jayaraman, SnO2: A comprehensive review on structures and gas sensors, Prog. Mater. Sci. 66 (2014) 112-255.

[3] C. Ferekides, J. Britt, CdTe solar cells with efficiencies over 15%, Sol. Energy Mat.

NU

Sol. Cells 35 (1994) 255-262.

MA

[4] N. Amin, T. Isaka, A. Yamada, M. Konagai, Highly efficient 1 μm thick CdTe solar cells with textured TCOs, Sol. Energy Mat. Sol. Cells 67 (2001) 195-201. [5] S. Chappel, S.-G. Chen, A. Zabon, TiO2-Coated Nanoporous SnO2 Electrodes for

TE

D

Dye-Sensitized Solar Cells, Langmuir 18 (2002) 3336-3342.

CE P

[6] S. Gabbala, V. Chakrapani, V. Kumar, M. K. Sunkara, Bad-Edge Engineered Hybrid Structures for Dye-Sensitized Solar Cells Based on SnO2 Nanowires, Adv. Funct. Mater. 18 (2008) 2411-2418.

[7] W. Göpel, K. D. Schierbaum, SnO2 sensors: current status and future prospects, Sens.

AC

Actuators B 26 (1995) 1-12.

[8] C. O. Park, S. A. Akbar, Ceramics for chemical sensing, J. Mat. Sci. 38 (2003) 4611-4637. [9] T. Moon, S.-T. Hwang, D.-R. Jung, D. Son, C. Kim, J. Kim, M. Kang, B. Park, Hydroxyl-Quenching Effects on the Photoluminescence Properties of SnO2:Eu3+, J. Phys. Chem. C 111 (2007) 464-4167. [10] F. Gu, S. F. Wang, M. K. Lü, Y. X. Qi, G. J. Zhou, D. Xu, D. R. Yuan, Luminescent properties of Mn2+-doped SnO2 nanoparticles, Inorg. Chem. Commun. 6 (2003) 882-885.

16

ACCEPTED MANUSCRIPT

[11] F. Gu, S. F. Wang, M. K. Lü, Y. X. Qi, G. J. Zhou, D. Xu, D. R. Yuan, Luminescent

T

characteristics of Eu3+ in SnO2 nanoparticles, Opt. Mater. 25 (2004) 59-64.

IP

[12] E. N. Dottol, Q. Wan, W. Guo, Y. Chen, X. Pan, W. Lu, Fully Transparent

SC R

Thin-Film Transistor Devices Based on SnO2 Nanowires, Nano Lett. 7 (2007) 2463-2469. [13] N. Yamazoe, K. Shimanoe, New perspectives of gas sensor technology, Sens.

NU

Actuators B 138 (2009) 100-107.

MA

[14] E. Comini, Metal oxide nanocrystals for gas sensing, Analytica Chimica Acta 568 (2006) 228-40. [15] E. Comini, G. Faglia, G. Sberveglieri, Z. Pan, Z. L. Wang, Stable and highly

TE

D

sensitive gas sensors based on semiconducting oxide nanobelts, Appl. Phys. Latt. 81 (2002) 1869-1871.

CE P

[16] Y. Chen, X. Cui, K. Zhang, D. Pan, S. Zhang, B. Wang, J. G. Hou, Bulk-quantity synthesis and self-catalytic VLS growth of SnO2 nanowires by lower-temperature evaporation, Chem. Phys. Lett. 369 (2000) 16-20.

AC

[17] Y. Shen, T. Ymazaki, Z. Liu, D. Meng, T. Kikuta, N. Nakatani, M. Saito, M. Mori, Microstructure and H2 gas sensing properties of undoped and Pd doped SnO2 nanowires, Sens. Actuators B 135 (2009) 524-529. [18] Y. Shen, T. Yamazaki, Z. Liu, D. Meng, T. Kikuta, Hydrogen sensors made of undoped and Pt-doped SnO2 nanowires, J. Alloy Cmpd. 488 (2009) L21-L25. [19] H. Li, H. Ma, Y. Zeng, A. Pan, Q. Zhang, H. Yu, T. Wang, Y. Wang, B. Zou, Hierarchical SnO2 Nanostructures: Linear Assembly of Nanorods on the Nanowire Backbones, J. Phys. Chem. C 114 (2010) 1844-1848. [20] J. Q. Hu, Y. Bando, D. Golberg, Self-catalyst growth and optical properties of novel SnO2 fishbone-like nanoribbons, Chem. Phys. Lett. 372 (2003) 758-762.

17

ACCEPTED MANUSCRIPT

[21] S. Phandunghitidhada, S. Thanasanvorakun, P. Mangkorntong, S. Choopun, N. Mangkorntong, D. Wangratanaphisan, SnO2 nanowires mixed nanodendrites for high

IP

T

ethanol sensor response, Curr. Appl. Phys. 11 (2011) 1368-1373. [22] P. G. Li, M. Lei, W. H. Tang, X. Guo, X. Wang, Facile route to straight SnO 2

SC R

nanowires and their optical properties, J. Alloy Comp. 477 (2009) 515-518. [23] D. Calestani, M. Zha, A. Zappettini, G. Salviati, L. Zanotti, G. Sberveglieri,

NU

Structural and optical study of SnO2 nanobelts and nanowires, Mater. Sci. Eng. C 25 (2005) 625-630.

MA

[24] W. Yin, B. Wei, C. Hu, In situ growth of SnO2 nanowires on the surface of Au-coated Sn grains using water-assisted chemical vapor deposition, Chem. Phys. Lett. 471 (2009) 11-16.

TE

D

[25] Y. Wang, X. Jiang, Y. Xia, A Solution-Phase, Precursor Route Polycrystalline SnO2 Nanowires That Can Be Used for Gas Sensing under Ambient Conditions, J. Am. Chem.

CE P

Soc. 125 (2003) 16176-16177.

[26] B. Cheng, J. M. Russel, W. Shi, L. Zhang, E. T. Samulski, Large-Scale,

AC

Solution-Phase, Growth of Single Crystalline SnO2 Nanorods, J. Am. Chem. Soc. 126 (2004) 5972-5973. [27] Y. Wang, V. Schmidt, S. Senz, U. Gӧsele, Epitaxial growth of silicon nanowires using an aluminum catalyst, Nat. Nanotechnol. 1 (2006) 186-189. [28] J. Johansson, Regime change for nanowire growth, Nat. Nanotechnol. 2 (2007) 534-535. [29] A. Klamchuen, M. Suzuki, K. Nagashima, H. Yoshida, M. Kanai, F. Zhuge, Y. He, G. Meng, S. Kai, S. Takeda, T. Kawai, T. Yanagida, Rational concept for designing vapor-liquid-solid growth of single crystalline metal oxide nanowires, Nano Lett. 15 (2015) 6406-6412. [30] T. Terasako, M. Yagi, M. Ishizaki, Y. Senda, H. Matsuura, S. Shirakata, Growth of zinc oxide films and nanowires by atmospheric-pressure chemical vapor deposition 18

ACCEPTED MANUSCRIPT

T

using zinc powder and water as source materials, Surf. Coat. Technol. 201 (2007) 8924-8930.

IP

[31] T. Terasako, T. Fujiwara, M. Yagi, S. Shirakata, Various Shapes of ZnO and CdO Nanostructures Grown by Atmospheric-Pressure Chemical Vapor Deposition, Jpn. J.

SC R

Appl. Phys. 50 (2011) art. no. 01BJ15.

[32] T. Terasako, T. Fujiwara, Y. Nakata, M. Yagi, S. Shirakata, Structural and optical properties of CdO nanostructures prepared by atmospheric-pressure CVD, Thin Solid

NU

Films 528 (2013) 237-241.

MA

[33] T. Terasako, H. Ichinotani, M. Yagi, Growth of β-gallium oxide films and nanostructures by atmospheric-pressure CVD using gallium and water as source materials, Phys. Stat. Sol. C 12 (2015) 985-988.

D

[34] T. Terasako, Y. Kawasaki, M. Yagi, Growth and morphology control of β-Ga2O3

TE

nanostructures by atmospheric-pressure CVD, Thin Solid Films 620 (2016) 23-29.

CE P

[35] C. G. Van de Walle, Hydrogen as a Cause of Doping in Zinc Oxide, Phys. Rev. Lett. 85 (2000) 1012-1015.

AC

[36] Ҫ. Kiliç, A. Zunger, n-type doping of oxides by hydrogen, Appl. Phys. Lett. 81 (2002) 73-75. [37] J. E. Dominguez, L. Fu, X. Q. Pan, Epitaxial nanocrystalline tin dioxide thin films grown on (0001) sapphire by femtosecond pulsed laser deposition, Appl. Phys. Lett. 79 (2001) 614-616. [38]Y.-W. Kim, S. W. Lee, H. Chen, Microstructural evolution and electrical property of n-doped SnO2 films grown on Al2O3 (0001) by metalorganic chemical vapor deposition, Thin Solid Films 405 (2002) 256-262. [39] L. C. Tien, D. P. Norton, J. D. Budai, Epitaxial growth of transparent tin oxide films on (0001) sapphire by pulsed laser deposition, Mater. Res. Bull. 44 (2009) 6-10.

19

ACCEPTED MANUSCRIPT

[40] L. Mazenia, Y. N. Picard, D. Caldwell, E. R. Glaser, S. M. Prokes, Growth and photoluminescence properties of vertically aligned SnO2 nanowires, J. Cryst. Growth

IP

T

311 (2009) 3158-3162.

Applications, Nano Lett. 14 (2014) 3014-3022.

SC R

[41] X. Wang, N. Aroonyadet, Y. Zhang, M. Mecklenburg, X. Fang, H. Chen, E. Goo, C. Zhou, Aligned Epitaxial SnO2 Nanowires on Sapphire: Growth and Device

[42] P. Yang, H. Yan, S. Mao, R. Russo, J. Johnson, R. Saykally, N. Morris, J. Pham, R.

NU

He, H.-J. Choi, Controlled Growth of ZnO Nanowires and Their Optical Properties, Adv.

MA

Funct. Mater. 12 (2002) 323-331.

[43] V. Bonu, A. Das, S. Amirthapandian, S. Dhara, A. K. Tyag, Photoluminescence of oxygen vacancies and hydroxyl group surface functionalized SnO2 nanoparticles, Phys.

D

Chem. Chem. Phys. 17 (2015) 9794-9801.

TE

[44] O. Madelung, Introduction to Solid-State Theory, Springer, Berlin, 1978, p.413.

CE P

[45] G. F. Neumark, K. Kosai in: R. K. Willardson, A. C. Beer (Eds), Semiconductor and Semimetals Vol. 19, Academic Press, New York, 1983, p. 1.

AC

[46] J. I. Pankove, OPTICAL PROCESSES IN SEMICONDUCTORS, Dover Publications, Inc., New York, 1971, p. 113. [47] D. Pal, D. N. Bose, Study of photoluminescence and computation of configuration coordinate diagram of Cu related deep levels in InP, Bull. Mater. Sci. 20 (1997) 401-407. [48] S. Sinha, A. K. Srivastava, S. Banerjee, B. M. Arora, 0.8-eV photoluminescence band in AlxGa1-xAs, Phys. Rev. B 44 (1991) 10941-10944. [49] M. A. Reshchikov, H. Morkoç, Luminescence properties of defects in GaN, J. appl. Phys. 97 (2005) art. no. 061301. [50] T. Arai, S. Adachi, Simple wet chemical synthesis and photoluminescence characterization of SnO2:Eu3+ reddish –orange phosphor, J. Lumin. 153 (2014) 46-53. 20

ACCEPTED MANUSCRIPT

T

[51] S. Shi, J. Xu, X. Zhang, L. Li, Effect of annealing on the structural and photoluminescence properties of ZnO nanorod arrays grown at low temperature, J. Appl.

IP

Phys. 109 (2011) art. no. 103508.

SC R

[52] J. Hu, Y. Bando, Q. Liu, D. Golberg, Laser-ablation growth and optical properties of wide and large single-crystal SnO2 ribbons, Adv. Funct. Mater. 13 (2003) 493-496. [53] G. Salviati, L. Lazzarini, M. Z. Zha, V. Grillo, E. Carlino, Cathodoluminescence

MA

NU

spectroscopy of single SnO2 naowires and nanobelts, Phys. Stat. Sol. a 202 (2005) 2963-2970. [54] J. D. Prades, J. Arbiol, A. Cirera, J. R. Morante, M. Avella, L. Zanotti, E. Comini, G. Faglia, G. Sberveglieri, Defect study of SnO2 nanostructures by cathodoluminescence

D

analysis: Application to nanowires, Sens. Actuators B 126 (2007) 6-12.

TE

[55] L.-W. Chang, M.-W. Huang, C.-T. Li, H. C. Shih, Controllable two-step growth and

CE P

photoluminescence of waterweed-like SnO2 nanowires, Appl. Surf. Sci. 279 (2013) 167-170. [56] S. Luo, J. Fang, W. Liu, M. Zhang, Z. Song, C. Liu, X. Wu, P. K. Chu, Synthesis and low-temperature photoluminescence properties of SnO2 nanowires and nanobelts,

AC

Nanotechnology 17 (2006) 1695-1699. [57] H. T. Chen, X. L. Wu, S. J. Xiong, W. C. Zhang, J. Zhu, Red photoluminescence mechanism in SnO2 nanostructures, Appl. Phys. A 97 (2009) 365-368. [58] H. T. Chen, S. J. Xiong, X. L. Wu. J. Zhu, J. C. Shen, Tin Oxide Nanoribbons with Vacancy Structures in Luminescence-Sensitive Oxygen Sensing, Nano Lett. 9 (2009) 1926-1931. [59] L. Li, Growth and photoluminescence properties of SnO2 nanobelts, Mater. Lett. 98 (2013) 146-148. [60] T. Gao, T. Wang, Vapor phase growth and optical properties of single-crystalline SnO2 nanobelts, Mater. Res. Bull. 43 (2008) 836-842. 21

ACCEPTED MANUSCRIPT

Figure captions

IP

T

Fig. 1. Schematic diagram of our AP-CVD system.

Fig. 2. XRD patterns of the samples grown by the alternate source supply (ASS) under

NU

SC R

the different growth conditions: (a) 30 nm thick Au film, Tg=925 C; (b) 20 nm thick Au film, Tg=900 C; (c) 20 nm thick Au film, Tg=825 C; and (d) 5 nm thick Au film, Tg=800 C. Expanded XRD patterns (×300) are also indicated for the two samples from the top. The peak marked by “?” cannot be identified. Fig. 3. Surface SEM images of the samples grown by the alternate source supply

MA

(ASS) under the different growth conditions: (a) 30 nm thick Au film, Tg=925 C; (b) 20

D

nm thick Au film, Tg=900 C; (c) 20 nm thick Au film, Tg=825 C; and (d) 5 nm-thick Au film, Tg=800 C.

TE

Fig. 4. Histograms for the diameters of the ASS SnO2 NWs determined from the SEM images shown in Fig. 3.

CE P

Fig. 5. Variations of NWs average diameters plotted as a function of 1000/Tg for the ASS SnO2 NWs grown using the Au films with the thicknesses of 5, 10, 20 and 30 nm.

AC

Fig. 6. Variations of NWs average diameters plotted as a function of cycle number for the ASS SnO2 NWs grown using the Au films with the thicknesses of 5, 10 and 30 nm. Fig. 7. (a) Photoluminescence (PL) and photoluminescence excitation (PLE) spectra of the ASS SnO2 NWs grown under the different growth conditions. The growth temperature (Tg) and the thickness of the Au film are shown in the figure. (c) photoacoustic (PA) spectra of the ASS SnO2 NWs grown using the 5 nm thick Au film at Tg=875 and 925 C. (c) Schematic configuration coordinate (CC) diagram for the orange band (OB) emission. Fig. 8. (a) Photoluminescence (PL) spectra of the ASS SnO2 NWs grown using the 30 nm-thick Au film at Tg=875, 900 and 925 C. (b) Variations of PL intensity at 640 nm plotted as a function of Tg for the ASS SnO2 NWs grown using the Au films with the thicknesses of 5, 20 and 30 nm.

22

D

MA

NU

SC R

IP

T

ACCEPTED MANUSCRIPT

AC

CE P

TE

Figure 1.

23

MA

NU

SC R

IP

T

ACCEPTED MANUSCRIPT

AC

CE P

TE

D

Figure 2.

24

TE

D

MA

NU

SC R

IP

T

ACCEPTED MANUSCRIPT

AC

CE P

Figure 3.

25

MA

NU

SC R

IP

T

ACCEPTED MANUSCRIPT

AC

CE P

TE

D

Figure 4.

26

MA

NU

SC R

IP

T

ACCEPTED MANUSCRIPT

AC

CE P

TE

D

Figure 5.

27

MA

NU

SC R

IP

T

ACCEPTED MANUSCRIPT

AC

CE P

TE

D

Figure 6.

28

TE

D

MA

NU

SC R

IP

T

ACCEPTED MANUSCRIPT

AC

CE P

Figure 7.

29

MA

NU

SC R

IP

T

ACCEPTED MANUSCRIPT

AC

CE P

TE

D

Figure 8.

30

ACCEPTED MANUSCRIPT

Highlights

IP

T

・The VLS growth mechanism was used for preparing the SnO2 nanowires. ・SnO2 nanowires were successfully grown by the alternate source supply of Sn and H2O.

AC

CE P

TE

D

MA

NU

SC R

・The alternate source supply was effective for suppressing the VS growth mechanism. ・The grown SnO2 nanowires exhibited a dominant orange band emission. ・The PL and PLE spectra of the orange band emission exhibited very large Stokes shift.

31