Oxygen Vacancy Driven Modulations in In2O3 Pyramidal Beaded

Article pubs.acs.org/crystal

Oxygen Vacancy Driven Modulations in In2O3 Pyramidal Beaded Nanowires Muhammad Hafeez,†,‡ Tianyou Zhai,*,‡,⊥ Arshad S. Bhatti,*,† Yoshio Bando,‡ and Dmitri Golberg‡ †

Center for Micro and Nano Devices (CMND), Department of Physics, COMSATS Institute of Information Technology, Islamabad, 44000, Pakistan ‡ International Center for Materials Nanoarchitectonics (WPI-MANA), National Institute for Materials Science (NIMS), Namiki 1-1, Tsukuba, Ibaraki, 305-0044, Japan ABSTRACT: We present the growth of pyramidal beaded In2 O 3 nanowires by using hydrogen assisted thermal evaporation. Reduction reaction at the source produces different growth species having varying vapor pressures, which is responsible for the growth of oxygen deficient nanostructures. The number and nature of oxygen vacancies affect the growth rates of different planes and thus the ultimate nanostructure morphology. A detailed growth mechanism of the nanowires is proposed on the basis of thus created oxygen vacancies. Morphology of the synthesized nanostructures was interpreted using electrical and structural analysis (VESTA) software. Structural, compositional, optical and field emission (FE) characteristics were studied to further confirm the oxygen deficient growth. The phonon confinement model (PCM) was used to calculate the correlation length of defects. The regarded nanowires were found to be good field emitters with low turn-on fields, from 5.8 to 14.5 V/μm, and field enhancement factors from 1775 to 362, depending on cathode−sample distances. The experimental FE data were fitted with the Philips model and two-region field emission (TRFE) model, and the screening effect, absolute amplification factor and width of field enhancement region were calculated. Our approach to fabricate beaded nanowires may open new avenues to synthesize unique nanostructures for novel optoelectronic devices.

■

The phenomenon of field emission from a material can be explained by using Fowler Nordheim (FN) theory.26 FN theory relates the field emission to the two important parameters, namely, the work function of an emitting material and the fieldenhancement factor.27 The work function is an intrinsic parameter that varies from 1 to 7 eV for most of the inorganic semiconductors, the lower being better for the higher emission efficiency.28 The field-enhancement factor is defined as the ratio of a local field to an applied field and is determined by the shape and morphology of the emitter. The work function of In2O3 (5.0 eV) is comparable to some other popular FE materials, such as C nanotubes (5.0 eV), ZnO (5.3 eV) and CdS (4.2 eV) nanowires, thus In2O3 is among good FE materials. It has a wide band gap (3.6 eV), so it has a large band bending in a strong applied field, also In2O3 has a tendency to grow in numerous morphologies, which makes it a potential material for FE applications. Many works have been reported with respect to FE studies of different morphologies in In2O3 with no or a limited success. Turn-on field (at 1 μA/cm2) from the aligned In2O3 nanowires was measured as 7 V/μm whereas for the not-aligned In2O3 nanowires this value became 10 V/ μm.29 In the case of pyramidal nanostructures, for micro-

INTRODUCTION

Research and development in the field of one-dimensional (1D) nanostructures such as nanowires, nanobelts and nanotubes presents some of the most exciting breakthroughs in the modern era of science and technology.1−3 A broad range of 1D oxide semiconductor nanostructures attained much attention due to their use in the applied areas.4−6 As an important oxide semiconductor, indium oxide with a direct band gap of 3.6 eV can be one of the most attractive conductive oxides for field emission because of its relatively low work function (5.0 eV) and low electron affinity ∼3.5 eV.7 Due to its unique optical, chemical and electronic properties, it has found many applications in gas sensors,8 field-emission displays,9 lithium-ion batteries,10 solar cells,11 biosensors,12 photocatalysis13 etc. It is well-known that morphology plays a crucial role in the physical properties of materials, e.g., optical, electrical, field emission and sensing. Until now, many standard In2O3 nanostructures, such as nanowires,14 nanorods,15 nanotubes,16 nanosheets,17 nanobelts,18 nanoflowers,19 nanocubes,20 nanocolumns,21 nanoparticles,22 nanoarrows,23 nanotowers24 and nanopyramids,25 have successfully been synthesized. But the present study, performed on the very complex pyramidal beaded In2O3 nanostructures prepared under full control, is undoubtedly quite unique. © 2012 American Chemical Society

Received: June 27, 2012 Revised: August 26, 2012 Published: August 28, 2012 4935

dx.doi.org/10.1021/cg300870y | Cryst. Growth Des. 2012, 12, 4935−4943

Crystal Growth & Design

Article

pyramids, the reported turn-on field (at 0.1 μA/cm2) was 9.8 V/μm, whereas in the case of nanopyramids, it was either 2.7 V/μm30 or 3.34 V/μm.31 Nanopyramids have shown better FE properties than the nanowires because the former have fewer surface defects while an FE current always flows on the structure surface. There have several methods to synthesize In2O3 nanostructures, namely, thermal evaporation,32 hydrothermal technique33 and chemical vapor deposition.34Among those, the thermal evaporation is the simplest way to get large density, high-quality and single crystalline In2O3 nanostructures. During thermal evaporation several researchers have tried to evaporate In metal under oxygen as an oxidation agent35 whereas different reduction agents have been studied, such as In2O3, an In2O3 powder in carbon environment,36,37 In2O3 in deionized (DI) water38 etc. We have used hydrogen assisted thermal evaporation, which is very rare because hydrogen is more reactive than any other reduction agent. Our aim was to enhance the field emission properties, and for this purpose we synthesized novel In2O3 pyramidal beaded nanowires. This aim was achieved by introducing oxygen vacancies which had controlled the growth rate of specific atomic planes. Thus in this paper we report an efficient and simple way to fabricate In2O3 pyramidal beaded nanowires by using hydrogen assisted thermal evaporation. The process led to the formation of oxygen vacancies, which played the key role in the nanostructure growth. We also demonstrate that these vacancies have been responsible for the reduction of planar densities, which results in the enhanced energy of fast growing planes. Scanning electron microscopy (SEM) and highresolution transmission electron microscopy (HRTEM) were used to determine the morphology, sizes and growth directions of the nanostructures. X-ray diffraction (XRD) and X-ray photoelectron spectroscopy (XPS) were applied for the structural and compositional analysis, and quantification of the oxygen vacancies. Raman and photoluminescence (PL) spectroscopies were employed to determine the effect of oxygen vacancies on optical properties of the nanostructures. Phonon confinement model (PCM) was utilized to calculate the correlation length between the defects. Finally, the field emission measurements were performed to determine the gain factor, turn-on voltage and other FE parameters and to compare them with other documented FE results on various In2O3 and other FE materials.

■

T6400 with 514 nm excitation wavelength) and a microphotoluminescence (PL) spectrometer (Horiba Jobin-Yvon S.A.S using He−Cd laser with excitation wavelength of 325 nm). The field emission properties were studied at room temperature in a high-vacuum chamber (4 × 10−6 Pa) using a 1 mm2 cross sectional area copper anode. A dc voltage sweeping from 100 to 1100 V was applied to the samples.

■

RESULTS AND DISCUSSION Figure 1a−d shows the SEM and TEM micrographs of pyramidal beaded In2O3 nanowires. Low- and high-magnifica-

Figure 1. (a, b) Low- and high-magnification FESEM images and (c, d) low- and high-magnification TEM images of In2O3 pyramidal beaded nanowires.

tion FESEM images reveal tens of micrometers long dense In2O3 nanowires with an average diameter in the range of 50 to 150 nm and pyramidal beads having dimensions of 250 to 400 nm, as depicted in Figure 1a,b. Low-magnification TEM image shows that the size variations of pyramidal beads are minor during the growth. This implies that the growth of beads took place in the identical environment, as illustrated in Figure 1c. High-magnification TEM image from the selected region is displayed in Figure 1d, from which the angle between the growth plane and normal to the edge planes is measured as 35°. Since the angle between the [110] and [111] orientations is close, 35.2°, thus it can be concluded that the edge planes of the pyramids are {111}. Figure 2a depicts the TEM image of another In2O3 nanowire; its corresponding energy dispersive X-ray spectrum (EDS) is presented in Figure 2b. The EDS spectrum confirmed the In rich growth of In2O3 wires, whereas the Cu signal came from the support grid. HRTEM images of the selected nanobead region and the interface between the two beads show the lattice fringes, as depicted in Figure 2c,d. It is clear from the HRTEM images that the nanowires are single crystals with [011̅] growth direction and the d spacing is 0.71 nm. The insets in the HRTEM images of Figure 2c,d display the selected area electron diffraction (SAED) patterns. These also confirm that the nanowire is a single crystal having a bcc crystal structure and the growth direction [011̅]. Many reports have analyzed the growth of different types of In2O3 structures like nanopyramids, nanowires and nanoarrows, but the presently obtained morphology is unique. The main cause for the formation of In2O3 nanopyramids in thermal

EXPERIMENTAL SECTION

Si (100) substrates with a thin (2−3 nm) native oxide layer were used for the growth of nanostructures. A 1 nm Mn (catalyst) thin film was deposited in an ultrahigh-vacuum chamber at room temperature. The catalyst coated substrates were then loaded in a central zone of a horizontal tube furnace along with a 99.99% pure In2O3 powder (1 g) in an alumina boat placed at the center of the tube. Then the furnace was heated to 1120 °C, and the substrates were kept in the downstream at 675 °C. N2 + 5% H2 gas was used as a carrier gas to transport a flux of vapors from the source boat to the Si wafers. The flow rate of the carrier gas was maintained at 20 sccm for all experiments. The synthesized nanostructures were characterized using a field emission scanning electron microscope SEM (Hitachi SU 8000), a high-resolution transmission electron microscope (HRTEM, JEM2100F) equipped with an X-ray energy dispersive spectrometer (EDS) and an X-ray diffractometer XRD (RINT 2200HF) for analyzing surface topography, imaging and phase identification, respectively. For a study of optical properties, measurements were done at room temperature by means of a Raman spectrometer (Horiba Jobin-Yvon 4936



Article

It is believed that the mechanism of the present growth is a vapor solid (VS) process. The nucleation sites became In rich and also contained oxides of In. The schematic of the growth is shown in Figure 3a. The defects in the nuclei became the

Figure 2. (a) TEM image of In2O3 pyramidal beaded nanowire, (b) energy dispersive X-ray spectrum (EDS), (c, d) HRTEM images of the selected region from the nanobead and the interface between the two beads. Figure 3. (a) Schematics of the growth of In2O3 pyramidal beaded nanowires. (b) HRTEM image of a pyramidal bead. (c) Structural visualization by using VESTA software.

evaporation is the oxygen deficient environment. We have used hydrogen gas as a reduction agent and intentionally created an oxygen deficient environment to decorate the In2O3 nanowires with pyramidal beads. A detailed description of the growth conditions and possible mechanisms is given as follows. The growth process can be divided in two stages, e.g., possible reactions taking place at the source and growth of nanowires on the Si substrate are considered. In the first step, In2O3 powder as a source was used; its melting point was 1910 °C. So it could not be evaporated at 1020 °C, so there should be a possibility of two kinds of reactions which take place at the source side at 1020 °C. In2O3(s) + 2H 2(g) ↔ In2O(g) + 2H 2O(g)

(1)

In2O3(s) + 3H 2(g) ↔ 2In(l) + 3H 2O(g)

(2)

nucleation sites for a helical dislocation. Under ongoing hydrothermal reduction reaction, In and In2O molecules collide continuously, forming helical steps, and further grow into oxygen deficient In2O3 NWs. For In2O3 with the bcc crystal structure, it is well-known that the surface energy relation between three low-indexed crystallographic planes is given as γ{111} < γ{100} < γ{110}, while the growth rate perpendicular to the growing plane has a relation of r{111} < r{100} < r{110}.40 For In2O3 pyramids, if R is the ratio of the growth rate along the ⟨100⟩ to that of the ⟨111⟩, then R = 1.73, i.e., growth perpendicular to the plane (111) is much faster than (100). This was indeed observed in the HRTEM images: the growth direction was [011̅], as documented in Figure 3b. The observed peculiar growth morphology was interpreted with the help of visualization of electrical and structural analysis (VESTA) software. This software uses periodic bond chain (PBC) theory for structural visualization.41 As revealed by the XRD results, the grown nanostructures show the formation of bcc structure, which belongs to the space group 206. The structure of the product was visualized, and it was found that, in an ideal case, the growth rate of the fast growing direction {110} is dominant and the growth rate in the {111} direction is relatively slow. Thus in the ideal case no pyramids would form for pure In2O3 because the growth of a pyramidal structure only takes place when the growth rate on the {110} plane is comparable with the growth rate on the {111} plane. The introduction of oxygen deficiencies in the crystal resulted in decrease in the planar density. In this way, the number of unpaired bonds increased, and as a consequence, the energy of the plane increased, which reduced the growth rate of the plane. As soon as the growth rate of the {110} plane became comparable with that of the {111} plane, the growth of pyramids started, as summarized in Figure 3c. Figure 4 shows the XRD pattern taken from In 2 O 3 nanowires. All the diffraction peaks may be identified to the

The melting temperature of In is as low as 152 °C at 1 atm, whereas the boiling point for In2O is 525 °C at 1 atm. So, likely, there were two kinds of evaporating species, In2O and In. Partial pressures of the respective species may give us a clue on which one would evaporate faster. The partial pressures for In2O and In were then calculated using the following relations.39 log PIn2O =

⎛ −10, 678 ⎞ 4 ⎜ ⎟ + 9.644 = 5.7 × 10 Pa ⎝ ⎠ T

(3)

−12, 860 − 0.7 log T + 7.83 = 851 Pa (4) T From the above equations, it is clear that the partial pressure of In2O is much higher than that of pure In. So, the probability of evaporation of In2O was much higher than that of In. A large amount of In rich species coming out of In2O3 was evaporated and transported to the substrate with the help of carrier gas. In the second step, vapors of In and In2O were driven by the flowing nitrogen gas and deposited on the substrate surface to form oxygen deficient indium oxide crystalline nuclei, as was also confirmed by the XPS results presented below. log PIn =

4937



Article

crystal structure, oxygen leaves as O2 thus leaving behind two uncompensated electrons. This results in a formation of a donor state with a binding energy of 1 eV. So this addition of 1 eV is observed in a higher binding energy peak. The existence of oxygen vacancies in the nonstoichiometric In2O3−X can be calculated by stoichiometric ratio (Sij) and is given by following formula:43 Si , j =

Ci I /ASFi = i Cj Ij/ASFj

(5)

Here Ci and Cj are the concentrations and Ii and Ij are the intensities of the emission spectra and ASF is the atomic sensitivity factor. In our case, the ratio of indium to oxygen is calculated as 1:1.31, and thus the synthesized compound is of In2O2.62 composition. Further investigations of the structural features and quality of the nanostructures were carried out by Raman spectroscopy at room temperature using a 514.5 nm wavelength. According to the XRD findings, the grown nanostructures have a bixbyite bcc crystal structure belonging to the space group Ia3 with the space group number 206. According to the group theory analysis, the 52 optical modes have the irreducible representation as shown:44

Figure 4. XRD pattern of In2O3 pyramidal beaded nanowires.

cubic bixbyite structure of In2O3 with the lattice parameter a = 10.118 Å, in good agreement with the literature (JCPDF card number 06-0416). The sharpness of the peaks reflects the good quality of the sample. The XRD spectrum has two strong peaks of [222] and [440]; this additionally confirms the growth direction, in accordance with the HRTEM results. The stoichiometry, quantification and information about the phases of the synthesized product were finally confirmed by the XPS analysis. Figure 5a,b shows the high-resolution In 3d and O 1s XPS spectra of indium oxide nanostructures, respectively. The binding energies were corrected taking into account the specimen charging and by referring to C (1s) at 284.60 eV. The In (3d5/2) and O (1s) peaks are located at around 444.6 and 530.7 eV, respectively, indicating that the major constituent phase of the as-synthesized nanostructures is In2O3. The oxygen peak is resolved into two symmetric peaks centered at 530.7 and 532.6 eV. The peak at a low binding energy (530.7 eV) is referred to as a low binding energy component (LBEC) and may be ascribed to the O 1s core peak bound with In. The peak at high binding energy (532.6 eV) is referred to as a high binding energy component (HBEC) and attributed to be due to oxygen vacancies.42 During creation of oxygen vacancy in a

Γopt = 5A1g + 5E1g + 5E 2g + 17Tg + 20Tu

(6)

The A1g, E1g, E2g and Tg are Raman active, and Tu modes are infrared active. Therefore 32 active modes were expected to be present in the Raman spectra from bcc In2O3. Figure 6a shows the typical Raman spectrum, where six modes from In2O3 and one from the Si substrate are observed. The modes centered at 110, 133, 310 (E1g), 368 (E2g), 498 (A1g) and 630 cm−1 are ascribed to the typical modes of bcc In2O3.45 The blue shift in peaks was proportional to the phonon confinement region as phonon could be confined by stacking faults, vacancies, boundaries or pores. In the present case, it is believed that this has happened due to the presence of oxygen vacancies in line with the XPS data. We have observed the E1g mode in the Raman spectrum which corresponds to the stretching mode of the In2O3 and is very sensitive to the presence of oxygen vacancies.45 It was observed that the E1g mode was blue-shifted by 2 cm−1 to 310 cm−1 compared to its bulk value In2O3 (308 cm−1) and the fwhm was approximately 15 cm−1. As a result of Gaussian−

Figure 5. (a, b) High-resolution In 3d and O 1s XPS spectra of In2O3 pyramidal beaded nanowires. 4938



Article

The Eg mode of the Raman spectrum was fitted with eq 7 by using the values Δ = 20 cm−1, a = 1.118 nm and ω0 = 308 cm−1, as shown in Figure 6c. From fitting, the estimated value of the correlation length was computed as 9 nm. This confirms that the regarded shift is due to defects introduced by oxygen vacancies and is not related to other structural defects. The mode observed at 630 cm−1 was due to the In−O bonds; fwhm of this peak was around 18 cm−1. The increased value of fwhm is another signature of the oxygen vacancies' presence. It was also important to determine the nature of oxygen vacancies, i.e., are they singly ionized or doubly ionized vacancies? For this purpose, room temperature PL measurements of the synthesized nanowires were carried out. Figure 7 depicts the PL spectrum of the nanowires taken at room temperature with an excitation energy of 325 nm. Pure

Figure 6. (a) Micro Raman spectrum of In2O3 pyramidal beaded nanowires. (b) Gaussian peak fittings of the Eg mode. (c) Fitting of the Eg mode with the phonon confinement model.

Figure 7. Micro PL spectrum of In2O3 pyramidal beaded nanowires.

Lorentzian peak fitting, it was resolved as a combination of the two peaks marked in Figure 6b. The one is a high-intensity peak at 308 cm−1, and the other is a low-intensity peak at a high frequency of around 313.9 cm−1. There are two possibilities for the blue shifting of the Eg mode and its splitting: the quantum confinement effect or nonstoichiometry in the sample, i.e., oxygen vacancies in the system. The quantum confinement effect was ruled out because the Bohr radius of In2O3 is 2.3 nm but the nanostructures were much larger. So it was believed that the blue shifting of the Eg mode and its splitting were due to oxygen vacancies in the system.37 An insight into the lattice disorder was obtained through utilizing the phonon confinement model (PCM). The regarded model proposes to describe the quality of a crystalline structure by introducing a parameter known as correlation length. The latter is defined as the average size of the material homogeneity region. In the present case, the correlation length is the distance between the adjacent oxygen vacancies. According to PCM, the Raman line intensity, I(ω) at the frequency ω can be written as46

In2O3 is a wide band gap material and does not emit in the visible range; quite contrarily, a broad peak in that range was observed and was attributed to defects in nanostructures. The peak was fitted with two Gaussian functions peaked at 565 and 640 nm. In nanocrystalline oxides, oxygen vacancy is known to be the most common defect that usually acts as the radiative center in the luminescence process. Thus, the peaks at 565 and 640 nm are due to oxygen vacancies. In fact there have been reports on oxygen vacancies related luminescence of yellow, 570 nm, 580 nm and red 646 nm emissions.48−50 XRD, XPS and Raman measurements have also earlier indicated the existence of oxygen vacancies. Normally such vacancies exist in × three different charge states: V×O, V•O and V•• O . VO is a very shallow donor that is unable to produce luminescence in the visible region. So it is expected that in the present case there were singly ionized vacancies responsible for the emission at 640 nm and doubly ionized oxygen vacancies responsible for the emission at 565 nm. These oxygen vacancies could act as reservoirs for charge carriers that would be helpful for the electrical current flow. On the other hand, pyramids would be good for the field emission if these contain fewer surface defects. During FE measurements current always flows on the surface of a wire. So In2O3 pyramidal beaded nanowires may be a good option for the FE applications compared to standard In2O3 nanowires. Motivated by these predictions, we performed field emission measurements at room temperature. FE measurements of In2O3 nanowires have shown that they are indeed potentially good field emitters. Figure 8a presents FE data taken at an 80 μm gap between the anode and the

·

I(ω) =

∫

|C(0, q)|2 d3q [ω − ω(q)]2 +

BZ 2 2

Γ0 2 2

( )

(7)

2

where |C(0,q)|2 = e−q L /16π is the Fourier coefficient, q is the wave vector in units 2π/a, L is the average distance between defects, Γ0 is the fwhm and ω(q) is the energy of the Raman Eg mode and is given by47 ω(q) = Δ × [1 − cos(qa)] + ω0

(8) 4939



Article

Figure 8. (a) J−E plot of In2O3 pyramidal beaded nanowires measured at a 80 μm gap between the anode and sample; the inset shows the corresponding F−N plot. (b) J−E plots and (c) corresponding F−N plots from In2O3 pyramidal beaded nanowires with different vacuum gaps. (d) Variation in turn-on field (at 1 μA/cm2) and the field-enhancement factor with a change in vacuum gap d.

sample in a vacuum chamber maintained at a pressure of 4 × 10−6 Pa. From the response of the emission current density as a function of the applied field (JE), it is concluded that the present structures have a low turn-on field (Eto, defined as the E at which J becomes 1 μA/cm2) of 11.7 V/μm. The FE current voltage characteristics were further analyzed by the Fowler Nordheim (FN) equation,51 ⎛ Aβ 2E2 ⎞ ⎛ −BΦ3/2 ⎞ ⎟ J=⎜ ⎟exp⎜ ⎝ Φ ⎠ ⎝ βE ⎠

from the F−N theory and current saturation was also observed probably due to the contact resistance between the nanowires and the substrate. Figure 8b illustrates the FE curves measured at different vacuum gaps (d). When a vacuum gap increases from 60 to 180 μm, the turn-on field decreases from 14.5 to 5.8 V/μm. Figure 8c demonstrates FN plots obtained at different distances. The calculated field enhancement factors (obtained from the slope of the linear part of FN plots) were 362, 574, 792, 905, 1279, 1527 and 1775 for d = 60, 80, 100, 120, 140, 160 and 180 μm, respectively. Figure 8d presents the variation of field enhancement factor and turn-on field as a function of the vacuum gap. As the vacuum gap increases, the field enhancement factor also increases. Dependence of field enhancement factor on the vacuum gap could be explained on the basis of Philip’s model. According to this model the field enhancement factor is not a sole characteristic of nanowires but of the entire setup and is dependent on the vacuum gap and the radius of the emitting tip. It is written as53

(9)

or ⎛ Aβ 2 ⎞ ⎛ J ⎞ BΦ3/2 ln⎜ 2 ⎟ = ln⎜ ⎟− ⎝E ⎠ βE ⎝ Φ ⎠

(10) −6

where A and B are constants with the values of 1.54 × 10 A eV V−2 and 6.83 × 103 V/μm eV−3/2, respectively, J is the current density, β is the field-enhancement factor, E is the applied field and Φ is the work function (5 eV for In2O3) of the emitting materials. It is known that the field enhancement factor is related to the emitter geometry (such as aspect ratios), crystal structure, vacuum gaps and the spatial distribution of emitting centers.52 The inset shows that the F−N plot at a distance of 80 μm and the linear variation of ln(J/E2) with (1/ E) follow the FN behavior. The calculated field enhancement factor from the slope of the F−N plot was 574. Some deviation

⎛d ⎞ d β = 1 + S⎜ − 1⎟ ≅ 1 + S ⎝r ⎠ r

if

d≫r

(11)

Here S is the screening factor and the calculated field enhancement factor by using Philips law (βcal) for a single wire emitter (i.e., S = 1) is given as 601, 801, 1001, 1201, 1401, 1601 and 1801 for d = 60, 80, 100, 120, 140, 160 and 180 μm, 4940



Article

respectively. It was interesting to plot the ratio between the βcal and βFN (calculated by the slope of FN curves), as shown in the Figure 9a. From the trend we could see that the ratio was high

where d2 is the width of the field-enhancement region near the nanowire surface and β0 is the absolute amplification factor, which was intrinsically determined by emitters independent of d and the applied field. The values of d2 (=195 nm) and β0 (1720) were determined by the linear fits, as given by the slope and the intercept, respectively. A comparison with other field emitters shows that the value of d2 for the present In2O3 nanowires is much higher than for CdS nanocones grown on a Si substrate (d2 = 110 nm),55 ZnO nanorods grown on a Si substrate (d2 = 62 nm)56 and a ZnS tapered nanowire emitter (d2 = 55 nm).57 Table 1 summarizes the comparative key FE Table 1. Comparison of the Key FE Parameters of Standard In2O3 and Other FE Materials Found in the Literature with the Present In2O3 Pyramidal Beaded Nanowire FE Performance emitter material and morphology In2O3 aligned nanowire In2O3 unaligned nanowire In2O3 micropyramid CdS nanowire

thermal evaporation rf magnetron sputtering thermal evaporation MOCVD

CNTs with catalyst particle ZnO nanowire

CVD growth

ZnS nanowires In2O3 nanowire decorated by pyramids

Figure 9. (a) Fitting of experimental data to Philip’s model. (b) Linear fit to the experimental data based on the TRFE model.

β−1 d r

−1

≅ (β − 1)

r d

if

β

ref

7

NA

25

10

NA

25

9.8

266

26

12.2 at 10 μA/cm2 4.9

555

55

1870

58

6 at 0.1 μA/ cm2 11.7 at 0.1 μA/cm2 5.8−14.5

847

59

522

60

1775−360

present work

■

CONCLUSIONS Pyramidal beaded In2O3 nanowires were successfully synthesized by using hydrogen assisted thermal evaporation. This was explained on the basis of the reduction reaction taking place at the source side. Different vapor growth species were formed and were responsible for the growth of oxygen deficient nanostructures. Oxygen vacancies had a large impact on the growth rates of different planes and on the ultimate morphology of the nanostructures. Possible growth mechanism for the present nanowires was discussed in detail and explained with the help of visualization of electrical and structural analysis (VESTA). Structural, compositional, optical and field emission characteristics were studied to confirm the oxygen deficient growth and its effect on the experimentally measured properties. Correlation length of defects was also calculated to be around 9 nm by using the phonon confinement model (PCM). Field emission properties of synthesized product showed low turn-on fields ranging from 5.8 to 14.5 V/μm and field enhancement factors ranging from 1775 to 362, respectively, at various cathode distances. The experimental data of field emission were successfully fitted with the Philips model and the two-region field emission (TRFE) model, and

d≫r (12)

Equation 12 clearly shows the dependence of the screening factor, S, on the radius of the emitting tip and the vacuum gap between the sample and anode. The value of S varied from 0 to 1, S = 1 for a single emitter which has no effect due to the environment and S → 0 for the most dense emitters. The calculated screening factors of the present system at different values of separation are plotted in Figure 9a. It is evident that, by increasing the vacuum distance, the screening was reduced, and hence the field enhancement became high. The ratio between the field enhancement factors thus reached unity due to wiping out of the screening effect. When 1/β versus 1/d is plotted, as shown in Figure 9b, it becomes clear that 1/β follows linearly with 1/d, i.e., the larger the vacuum gap, the higher the value of β. On the basis of the two-region field-emission (TRFE) model, the results could be approximated by54 d 1 1 = 2 + β d β0

thermal evaporation vapor phase deposition hydrogen assisted thermal evaporation

turn-on field at 1 μA/cm2 (V/μm)

parameters of In2O3 nanostructures with different FE materials found in the literature. Decent emission demonstrated that the present structures could find potential applications in the coldcathode-based electronics.

in the beginning and then decreased, and finally reached unity at 180 μm. This was explained by using the following relation: S=

synthesis method

(13) 4941



Article

(21) Kar, S.; Chakrabarti, S.; Chaudhuri, S. Nanotechnology 2006, 17, 3058−3062. (22) Murali, A.; Barve, A.; Leppert, V. J.; Risbud, S. H.; Kennedy, I. M.; Lee, H. W. H. Nano Lett. 2001, 1, 287−289. (23) Yan, Y.; Zhang, Y.; Zeng, H.; Zhang, J.; Cao, X.; Zhang, L. Nanotechnology 2007, 18, 175601−175606. (24) Jean, S. T.; Her, Y. C. Cryst. Growth Des. 2010, 10, 2104−2110. (25) Zhang, Y.; Jia, H.; Yu, D.; Luo, X.; Zhang, Z.; Chen, X.; Lee, C. J. Mater. Res. 2011, 18, 2793−2798. (26) Junpeng, L.; Tao, Y.; Minrui, Z.; Haur, S. C. J. Phys. Chem. C 2011, 115 (17), 8816−8824. (27) Ye, C. H.; Bando, Y.; Fang, X. S.; Shen, G. Z.; Golberg, D. J. Phys. Chem. C 2007, 111, 12673−12676. (28) Fang, X. S.; Bando, Y.; Gautam, U. K.; Ye, C. H.; Golberg, D. J. Mater. Chem. 2008, 18, 509−522. (29) Li, S. Q.; Liang, Y. X.; Wang, T. H. Appl. Phys. Lett. 2005, 87, 143104. (30) Jia, H.; Zhang, Y.; Chen, X.; Shu, J.; Luo, X.; Zhang, Z.; Yu, D. Appl. Phys. Lett. 2003, 82, 4146. (31) Kar, S.; Chakrabarti, S.; Chaudhuri, S. Nanotechnology 2006, 17, 3058−3062. (32) Pan, C. A.; Ma, T. P. Appl. Phys. Lett. 1980, 37, 163. (33) Elouali, S.; Bloor, L. G.; Binions, R.; Parkin, I. P.; Carmalt, C. J.; Darr, J. A. Langmuir 2012, 28 (3), 1879−1885. (34) Kuo, D. H.; Hsu, C. W.; Liang, C. H. J. Electrochem. Soc. 2008, 155, K156. (35) Kar, S.; Chaudhuri, S. Chem. Phys. Lett. 2006, 422, 424−428. (36) Wu, X.; Hong, J.; Han, Z.; Tao, Y. Chem. Phys. Lett. 2003, 373, 28−32. (37) Kumar, M.; Singh, V. N.; Singh, F.; Lakshmi, K. V.; Mehta, B. R.; Singh, J. P. Appl. Phys. Lett. 2008, 92, 171907. (38) Jeong, J. S.; Lee, J. Y. Nanotechnology 2010, 21, 405601. (39) Chatterji, D.; Vest, R. W. J. Am. Ceram. Soc. 1972, 55, 575−578. (40) Hao, Y.; Meng, G.; Ye, C.; Zhang, L. Cryst. Growth Des. 2005, 5, 1617−1621. (41) Momma, K.; Izumi, F. J. Appl. Crystallogr. 2011, 44, 1272−1276. (42) Rumaiz, A. K.; Ali, B.; Ceylan, A.; Boggs, M.; Beebe, T.; Ismat Shah, S. Solid State Commun. 2007, 144, 334−338. (43) Wu, Q. H.; Song, J.; Kang, J.; Dong, Q. F.; Wu, S. T.; Sun, S. G. Mater. Lett. 2007, 61, 3679−3684. (44) Berengue, O. M.; Rodrigues, A. D.; Dalmaschio, C. J.; Lanfredi, A. J. C.; Leite, E. R.; Chiquito, A. J. J. Phys. D: Appl. Phys. 2010, 43, 045401. (45) White, W. B.; Keramidas, V. G. Spectrochim. Acta, Part A 1972, 28, 501−509. (46) Begum, N.; Bhatti, A. S.; Jabeen, F.; Rubini, S.; Martelli, F. J. Appl. Phys. 2009, 106, 114317. (47) Zhang, W. F.; He, Y. L.; Zhang, M. S.; Yin, Z.; Chen, Q. J. Phys. D: Appl. Phys. 2000, 33, 912−916. (48) Lim, T.; Lee, S.; Meyyappan, M.; Ju, S. ACS Nano 2011, 5, 3917−3922. (49) Jeong, J. S.; Lee, J. Y.; Lee, C. J.; An, S. J.; Yi, G. Chem. Phys. Lett. 2004, 384, 246−250. (50) Jean, S.-T.; Her, Y.-C. Cryst. Growth Des. 2010, 10, 2104−2110. (51) Li, L.; Fang, X. S.; Chew, H. G.; Zheng, F.; Liew, T. H.; Xu, X. J.; Zhang, Y. X.; Pan, S. S.; Li, G. H.; Zhang, L. D. Adv. Funct. Mater. 2008, 18, 1080−1088. (52) Gautam, U. K.; Fang, X. S.; Bando, Y.; Zhan, J. H.; Golberg, D. ACS Nano 2008, 2, 1015−1021. (53) Pan, N.; Xue, H.; Yu, M.; Cui, X. Nanotechnology 2010, 225707. (54) Zhong, D. Y.; Zhang, G. Y.; Liu, S.; Sakurai, T.; Wang, E. G. Appl. Phys. Lett. 2002, 80, 506−508. (55) Zhai, T. Y.; Fang, X. S.; Bando, Y.; Liao, Q.; Xu, X.; Zeng, H.; Ma, Y.; Yao, J.; Golberg, D. ACS Nano 2009, 3 (4), 949−959. (56) Li, C.; Fang, G. J.; Liu, N. S.; Li, J.; Liao, L.; Su, F. H.; Li, G. H.; Wu, X. G.; Zhao, X. Z. J. Phys. Chem. C 2007, 111, 12566−12571. (57) Hafeez, M.; Zhai, T.; Bhatti, A. S.; Bando, Y.; Golberg, D. J. Phys. Chem. C 2012, 116, 8297−8304.

the screening effect, absolute amplification factor and width of the field enhancement region were determined. Our approach to grow beaded nanowires may pave the way for fabrication of other unique nanostructures to be applied in novel optoelectronic devices.

■

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected], [email protected]. Present Address ⊥

Department of Material Science and Engineering, Tsinghua University, Beijing 100084, P. R. China. Notes

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS This work was supported by the World Premier International (WPI) Research Center for Materials Nanoarchitectonics (MANA), MEXT, Japan. The work was also supported by HEC Grant No. 261, 1770 and HEC’s development grant for Micro and Nano Fabrication facilities. M.H. is thankful to IRSIP Program of Higher Education Commission, Pakistan (HEC), and the National Institute for Materials Science (NIMS) for financial support and an Internship award.

■

REFERENCES

(1) Zhai, T. Y.; Li, L.; Ma, Y.; Liao, M.; Wang, X.; Fang, X. S.; Yao, J.; Bando, Y.; Golberg, D. Chem. Soc. Rev. 2011, 40, 2986−3004. (2) Fang, X. S.; Zhai, T. Y.; Gautam, U. K.; Li, L.; Wu, L.; Bando, Y.; Golberg, D. Prog. Mater. Sci. 2011, 56, 175−287. (3) Zhai, T. Y.; Fang, X. S.; Liao, M. Y.; Xu, X. J.; Li, L.; Liu, B. D.; Koide, Y.; Ma., Y.; Yao, J. N.; Bando, Y.; Golberg, D. ACS Nano 2010, 4, 1596−1602. (4) Li, L.; Zhang, Y.; Fang, X.; Zhai, T.; Liao, M.; Sun, X.; Koide, Y.; Bando, Y.; Golberg, D. J. Mater. Chem. 2011, 21, 6525−6530. (5) Zhai, T.; Liu, H.; Li, H.; Fang, X.; Liao, M.; Li, L.; Zhou, H.; Koide, Y.; Bando, Y.; Golberg, D. Adv. Mater. 2010, 22, 2547−2552. (6) Zhai, T.; Fang, X.; Liao, M.; Xu, X.; Zeng, H.; Yoshio, B.; Golberg, D. Sensors 2009, 9, 6504−6529. (7) Klein, A. Appl. Phys. Lett. 2000, 77, 2009. (8) Li, E.; Cheng, Z.; Xu, J.; Pan, Q.; Yu, W.; Chu, Y. Cryst. Growth Des. 2009, 9, 2146−2151. (9) Jia, H.; Zhang, Y.; Chen, X.; Shu, J.; Luo, X.; Zhang, Z.; Yu, D. Appl. Phys. Lett. 2003, 82, 4146. (10) Yang, Y.; Jeong, S.; Hu, L.; Wu, H.; Lee, S. W.; Cui, Y. Proc. Natl. Acad. Sci. U.S.A. 2011, 108, 13013−13018. (11) Meiss, J.; Riede, M. K.; Leo, K. Appl. Phys. Lett. 2009, 94, 013303. (12) Curreli, M.; Li, C.; Sun, Y.; Lei, B.; Gundersen, M. A.; Thompson, M. E.; Zhou, C. J. Am. Chem. Soc. 2005, 127, 6922−6923. (13) Yang, X.; Wang, Y.; Xu, L.; Yu, X.; Guo, Y. J. Phys. Chem. C 2008, 112, 11481−11489. (14) Li, C.; Zhang, D.; Han, S.; Liu, X.; Tang, T.; Zhou, C. Adv. Mater. 2003, 15, 143−146. (15) Zhang, R.; Jiang, K. M.; Chen, D. L.; Ding, G. Q. Mater. Lett. 2009, 63, 1044−1046. (16) Du, N.; Zhang, H.; Chen, B. D.; Ma, X. Y.; Liu, Z. H.; Wu, J. B.; Yang, D. R. Adv. Mater. 2007, 19, 1641−1645. (17) Yang, H.; Zhang, R.; Dong, H.; Yu, J.; Yang, W.; Chen, D. Cryst. Growth Des. 2008, 8, 3154−3159. (18) Kong, X. Y.; Wang, Z. L. Solid State Commun. 2003, 128, 1−4. (19) Du, J.; Yang, M.; Nam Cha, S.; Rhen, D.; Kang, M.; Kang, D. J. Cryst. Growth Des. 2008, 8, 2312−2317. (20) Tang, Q.; Zhou, W.; Zhang, W.; Ou, S.; Jiang, K.; Yu, W.; Qian, Y. Cryst. Growth Des. 2005, 5, 147−150. 4942



Article

(58) Jang, H. S.; Lee, H. R.; Kim, D. H. Thin Solid Films 2006, 500, 124−128. (59) Lee, C. J.; Lee, T. J.; Lyu, S. C.; Zhang, Y.; Ruh, H. Appl. Phys. Lett. 2002, 81, 3648−3650. (60) Chang, Y. Q.; Wang, M. W.; Chen, X. H.; Ni, S. L.; Qiang, W. J. Solid State Commun. 2007, 142, 295−298.

4943


Oxygen Vacancy Driven Modulations in In2O3 Pyramidal Beaded

Recommend Documents