Tailoring of Point Defects in Polycrystalline Indium Tin Oxide Films

Article Cite This: ACS Appl. Electron. Mater. XXXX, XXX, XXX−XXX

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Tailoring of Point Defects in Polycrystalline Indium Tin Oxide Films with Postirradiation of Electronegative Oxygen Ions Yutaka Furubayashi,*,† Makoto Maehara,‡ and Tetsuya Yamamoto† †

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Materials Design Center, Research Institute, Kochi University of Technology, Tosayamadacho-Miyanokuchi 185, Kami 782-8502, Japan ‡ Industrial Equipment Division, Sumitomo Heavy Industries, Ltd., Soubiraki-cho 5-2, Niihama 792-8588, Japan ABSTRACT: We have developed a state-of-the-art technology to tailor oxygen-related point defects such as oxygen vacancies (VO) and structural defects of polycrystalline highly conductive Sn-doped In2O3 (ITO) films by a postirradiation of electronegative oxygen (O−) ions. The intentional oxygen doping that would annihilate VO decreases carrier density (ne) from 9.3 × 1020 to 7.1 × 1020 cm−3 with an increase of Hall mobility (μH) from 44 to 51 cm2·V−1·s−1, and subsequently, ne drastically decreases down to 1.2 × 1019 cm−3 together with a decrease in μH owing to a formation of Sn−O neutral complexes with a further increase in the amount of oxygen atoms that should fill the structural defects while retaining their crystal structure. Upon the filling of the structural vacancies, the successful control of intrinsic point defects is probably caused by almost no difference in the ionic radii between In3+ having six-coordination and Sn4+ having eight-coordination. The O−-ion-irradiation technology enables one to tailor the physical properties of ITO films over a wide range while retaining the crystal structure of the films. KEYWORDS: electronegative oxygen ions, indium oxide, carrier density, point defects, ionic radii

1. INTRODUCTION Tin (Sn)-doped In2O3 (ITO) is widely applied to transparent conducting oxide (TCO) films because of its high optical transmittance and low electrical resistivity.1−3 In2O3 has the cubic bixbyite structure (space group Ia-3, number 206) with 32 cations and 48 anions in a unit cell. The distorted InO6 octahedra can be regarded as virtually cubic structures, where only six of the eight vertices of the anionic cubes are occupied by oxygen (O) atoms.4,5 This shows that In2O3 originally has structural O defects (VSO). The high carrier density (ne) of ITO films can be attributed to both partial Sn4+ substitution for In3+ and positively charged oxygen vacancies. The value of ne depends on not only the concentration of activated Sn dopant atoms6 but also strongly on the oxygen partial pressure during film growth7 or post annealing.8−10 Frank and Köstlin observed the formation of bound complexes of Sn4+ ions at doping concentrations above 5 cation at.% of Sn by the examination of the ionizable and nonionizable Sn−O defects.11 Note that the word “complex” means a local structure within a crystal network and not an individual chemical species. To explain the low activity of Sn dopants, some studies have focused on the VSO sites. The O atoms that occupy the VSO sites act as interstitial O atoms (Oi) to form SnOx complexes, with Sn atoms inactivated.5,11 Note that the neutral complexes with a large scattering cross section present a significant obstacle to free carriers, reducing the carrier mobility of ITO films. As described above, many reports on the electrical properties of ITO films have treated them as a function of Sn doping.6,11 On the other hand, in this paper, for ITO films with Sn dopants fixed at a high doping level, the © XXXX American Chemical Society

concentration of related point defects was controlled by the irradiation of electronegative oxygen (O−) ions with the aim of clarifying the characteristics of the behavior of ITO films with various ne in terms of point-defect physics. The details will be described in Section 4. In this study, we investigated the physical properties of ITO films as a function of oxidation. First, we deposited ITO films with a high ne of 9.3 × 1020 cm−3. Then, we irradiated negatively charged oxygen (O−) ions instead of positively charged oxygen (O+) ions,12 where the latter cause the chargeup of the films. The process conditions for the O−-irradiation were selected to reduce the damage such as the deterioration of crystallinity and the change in [SnIn+], owing to the movement from the original site or the evaporation of Sn atoms. We have been developing a state-of-the-art technology, which is commercially produced by Sumitomo Heavy Industries, Ltd., to generate and irradiate O− ions by using reactive plasma deposition (RPD) with direct-current (DC) arc discharge,13 exhibiting a high arc plasma density of 1012 cm−3. RPD with a high growth rate has enabled us to fabricate oxide films with a uniform spatial distribution of thickness prepared on large substrates14,15 with sizes such as 1.5 × 1.5 m2. Unlike conventional sputtering deposition, a characteristic of RPD is the use of positively charged ions with low energies, such as In+ and Zn+ ions,16 supplied by DC arc discharge; for Received: May 20, 2019 Accepted: July 4, 2019

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DOI: 10.1021/acsaelm.9b00317 ACS Appl. Electron. Mater. XXXX, XXX, XXX−XXX

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ACS Applied Electronic Materials the fabrication of ITO and ZnO films, the ion energies range from 15 to 60 eV and from 20 to 60 eV, respectively. This lowion-energy technology allows us to fabricate high-quality oxide films with the near-bulk densities. In fact, we have recently achieved a dense 10 nm-thick ZnO17 and ITO18 films having near-bulk volume densities. In this paper, we report the structural, electrical, and optical properties of polycrystalline ITO films modified with O−-ion irradiation. We found that O−-ion irradiation causes a large reduction in ne for the ITO films from 9.3 × 1020 to 1.2 × 1019 cm−3 while retaining the original crystal structure of the films.

Figure 1. Schematic diagram of RPD during the O−-ion irradiation after the growth, where the ITO film is connected to a current collector and a trolley line to apply a bias voltage to attract O− ions.

2. METHODS 2.1. Principle of Process. Thin film growth of polycrystalline ITO and a subsequent postirradiation of O− ions were performed by using the RPD equipment. During ITO film growth with a DC arc discharge, the deposition chamber is filled with a weakly ionized plasma comprising a partially ionized gas consisting of electrons (e−); positively charged ions such as O+, O2+, Ar+, In+, and Sn+ and negatively charged O− ions; a large number of neutral species such as In, Sn, O, and Ar atoms; and oxygen free radicals including a singlet oxygen (1O2). The O+, O2+, In+, Sn+, In, Sn, and O species propagating toward a substrate contribute to the orientation of crystallites on the substrates and subsequently the film growth. When the DC arc plasma is switched off while O2 gas is introduced into the deposition chamber at a constant rate, on the basis of the analysis of the data obtained by plasma diagnostics, additional O− will be formatted by the following reactions19

e− + 1O2 → O− + O

(1)

e− + O → O−

(2)

e− + O2 → O2− −

−

−

e + O2 → O + O + e

results. The electrical properties at room temperature were determined by Hall effect measurements in a van der Pauw geometry (Nanometrics HL5500PC). Optical spectra with wavelengths between 200 and 2000 nm were taken by a conventional optical spectrometer (Hitachi U-4100).

3. RESULTS Figure 2 shows the electrical properties of the ITO films determined by Hall effect measurements at room temperature

(3) −

(4)

where the residual O2 reacts with the path (eq 1 → 2) and stable O2 reacts with the path (eq 3 → 4 → 2). In order to irradiate O− ions, a positive bias voltage is applied to the substrate. Because ionized species are delayed during time without plasma, a pulsed switching process for the plasma is needed to maintain the concentration of O− ions. 2.2. Thin Film Growth. Polycrystalline ITO films with a thickness of 50 nm were grown on nonalkali glass substrates (Corning Eagle XG) at a substrate temperature of 200 °C. A sintered mixture of In2O3 and SnO2 (SnO2 concentration of 5 wt %) was used as a source material. The flow rates of Ar gas introduced into the deposition chamber and a plasma gun were 25 and 40 sccm, respectively. The flow rate of oxygen (O2) gas into the chamber was 12 sccm. The total pressure during the film growth was 0.3 Pa. The typical growth rate was 3.6 nm/s. The details of the RPD equipment for thin film deposition were already illustrated in ref 18. 2.3. Postirradiation of O− Ions. After the deposition, we carried out a postirradiation of O− ions at 250 °C, as illustrated in Figure 1. A zirconium metal ingot was inserted at the position of the ITO source to avoid evaporation from the source position. A pulsed Ar/O2 mixture plasma was generated with a frequency of 60 Hz and an on/ off duty ratio of 50%. When the plasma was switched off, a voltage of 100 V was applied to the polycrystalline ITO films with the same frequency. The flow rates of Ar gas introduced into the deposition chamber and a plasma gun were 105 and 40 sccm, respectively. The flow rate of oxygen (O2) gas into the chamber was 15 sccm. 2.4. Evaluation of Physical Properties. X-ray diffraction (XRD) and X-ray reflectivity (XRR) measurements were performed with a Rigaku ATX-G diffractometer having an X-ray source of Cu Kα (wavelength of 0.15405 nm) to identify the structural properties of the ITO films. The mass density (dm), surface roughness (rsur), and roughness at the interface between the ITO films and the substrates (rif) of the ITO films were determined by the XRR measurement 1

Figure 2. (a) Electrical resistivity ρ, (b,d) carrier density ne, and (c,e) Hall mobility μH for polycrystalline ITO films as a function of O−-ionirradiation time tirr. The plots of ne and μH have different ranges of tirr with (b,c) 0 ≤ tirr ≤ 15 min and (d,e) 15 ≤ tirr ≤ 120 min. Closed symbols denote the data of as-deposited polycrystalline ITO films.

as a function of O−-ion-irradiation time (tirr). A series of tirr and the resulting ne are listed in Table 1. As-deposited ITO films had a low electrical resistivity (ρ) of 1.5 × 10−4 Ω·cm, as shown in Figure 2a, with an ne of 9.3 × 1020 cm−3 (Figure 2b). With increasing tirr, ρ increased, and ne decreased monotonically. Figure 2d shows that ne became as small as 1.2 × 1019 cm−3 as a result of irradiation. This indicates that 98.7% of the original free carriers were suppressed due to the O−-ion irradiation. On the other hand, we found interesting behavior B


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ACS Applied Electronic Materials Table 1. Relation between Irradiation Time tirr and Corresponding Carrier Density ne tirr/min ne/1020 cm−3

0 9.26(3)

2 7.07(11)

5 5.34(6)

15 2.05(7)

30 0.67(10)

60 0.37(9)

120 0.12(1)

Figure 3. XRD measurement results of polycrystalline ITO films after O−-ion irradiation. (a) 2θ/ω profiles, (b) lattice parameter a, (c) peak intensity ratios of (134)/(222) (triangles) and (440)/(222) (diamonds), and (d) peak widths of (222) peaks. (b−d) are plotted as a function of irradiation time tirr. Closed symbols in (b−d) denote the data for as-deposited polycrystalline ITO films.

in the variation of μH. With increasing tirr up to 2 min, ne decreased to 7.1 × 1020 cm−3, and μH increased to more than 50 cm2·V−1·s−1, as respectively shown in Figure 2b,c. With further increasing tirr, however, μH started to decrease and then abruptly decreased at tirr > 20 min (see Figure 2d,e). This unexpected tendency of μH will be discussed later on the basis of a carrier-scattering model. Figure 3 shows XRD measurement results plotted against tirr. In Figure 3a, three typical peaks of In2O3, (222), (440), and (134), can be observed. No additional phase nor precipitation such as SnOx were detected for these films. The shape of these spectra was almost identical irrespective of tirr. Figure 3b−d show the lattice parameter a, the integrated intensities of the (440) and (134) peaks, I(440) and I(134), relative to that of the (222) peak, and the peak width of 2θ of the (222) peak (Δ2θ(222)), respectively. a tended to decrease with increasing tirr, and the difference among the values of a was less than 0.1% over the entire range of tirr. From Figure 3c, at tirr = 2 min, the (110) orientation became slightly dominant. Taking into account the fact that the ratio of I(134) to I(222) was larger than that of I(440) to I(222) at any given tirr, we found that the O−-ion irradiation only slightly affected the orientation relationship. Figure 3d shows that Δ2θ(222) at tirr = 120 min was slightly increased by 17% compared with that of the asdeposited ITO films. The dm, rsur, and rif of the ITO films yielded from XRR measurement results are plotted in Figure 4. The XRR profiles were successfully fitted based on a simple two-layer model, that is, ITO on glass substrates without any density distribution. The dm values remain constant at 7.17 g/ cm3 within an error of ±0.3% irrespective of tirr. The rif also changed little. On the other hand, rsur increased abruptly by

Figure 4. (a) Mass density (dm), (b) surface roughness (rsur), and (c) interface roughness (rif) of polycrystalline ITO films after O−-ion irradiation derived from XRR measurement results.

20% at tirr = 120 min. This would possibly be due to an etching by O−-ion irradiation. Note that tirr values for all the samples were less than 1.0 nm, which is much smaller than the thickness of a monolayer of ITO. The above findings lead to the conclusion that the O−-ion irradiation did not significantly C


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ACS Applied Electronic Materials

Figure 5. Optical absorption coefficient spectra α for O−-ion-irradiated polycrystalline ITO films: wide spectra for the films with different tirr, (a) 0 ≤ tirr ≤ 15 min and (b) 15 ≤ tirr ≤ 120 min; spectra near the band edge for different tirr, (c) 0 ≤ tirr ≤ 15 min and (d) 15 ≤ tirr ≤ 120 min.

affect the crystallinity nor the surface morphology of the ITO films, whereas it decreased ne markedly. Figures 5−7 summarize the optical properties of the ITO films. The complex refractive index (N) was derived from the

Figure 6. Apparent band-gap energy Eg of O−-ion-irradiated polycrystalline ITO films derived from their optical spectra. At ne of more than 5 × 1020 cm−3, the value of Eg is saturated, owing to the absorption edge of a nonalkali glass substrate.

Figure 7. Real part of refractive index spectra (n) for O−-ionirradiated polycrystalline ITO films with different turr, (a) 0 ≤ t irr ≤ 15 min and (b) 15 ≤ t irr ≤ 120 min.

optical transmittance and reflectance spectra by using Fresnel’s equation to eliminate the effects of multiple reflections. N is defined by the equation N = n − ik, where the real part of N, namely n, is the normal refractive index, and the imaginary part, k, is called the extinction coefficient. Figures 5a,b and 5c,d show the absorption coefficient α (= 4πk/λ, where λ denotes wavelength) in the entire range of λ ranging from 200 to 2000 nm and near the band edge, respectively. From Figure 5a,c, we found a monotonic decrease in α in the near-infrared region at any given λ and a red-shift of the absorption edge with increasing tirr up to 5 min (ne > 5 × 1020 cm−3). This is explained very well by the decrease in ne caused by the O−-ion irradiation. On the other hand, with further increasing tirr, the shape of α hardly changed as shown in Figure 5d, whereas the shape of α in the near-infrared region exhibited complicated behavior as shown in Figure 5b. Figure 6 plots the apparent band-gap energy Eg as a function of ne. Basically, an increase in ne led to an enhancement in Eg. Figure 7a,b shows n as

functions of λ ranging from 300 to 900 nm. For the ITO films with an ne of more than 5.3 × 1020 cm−3, corresponding to a tirr of less than at 5 min, n became less than 2 at a λ larger than 500 nm, as shown in Figure 7a. This indicates that a tirr within 15 min, corresponding to ne = 2.1 × 1020 cm−3, can be used as a factor to control n in the visible and near-infrared λ regions: for example, between 1.88 and 2.05 at λ = 500 nm. This is probably caused by the effect of free-carrier absorption.20,21 There was no effect of a long irradiation of tirr of more than 30 min, which resulted in an ne of less than 1 × 1020 cm−3, on n in the above λ regions.

4. DISCUSSION As shown in the previous section, the O−-ion irradiation greatly modified both the electrical and optical properties with almost no significant change in the crystal structure. ne is the D


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ACS Applied Electronic Materials dominant factor limiting the optical properties of as-deposited and O−-ion-irradiated ITO films. In the following, we discuss the relationship among O− irradiation and the changes in ne with μH as shown in Figures 2b,d. Note that the atomic concentration of Sn in the source material was 1.4 × 1021 cm−3 (4.6 at%), which corresponds to the maximum dose concentration of Sn atoms. To date, we have no data for the residual concentration of Sn atoms in the resulting ITO films. On the other hand, as shown in Figures 3 and 4, XRD and XRR results of the ITO films showed little difference among the samples. This implies that the O−-ion irradiation caused no effect on the total amount and the arrangement of each host In atom and Sn dopant if we consider that the response from an X-ray is mainly limited by heavy atoms with a high atomic scattering factor. A comparison between the tirr dependencies of ne and μH makes it possible to understand the phenomena that occurred in the ITO films upon O−-ion irradiation, together with the conventional ionized impurity scattering mechanism, as discussed below. Let us begin by assuming that there is no net charge in an ITO semiconductor. This condition is referred to as charge neutrality. This also means that the total density of positively charged particles (holes and ionized donors) must equal that of negatively charged particles (electrons and ionized acceptors). With a notation similar to Kröger−Vink notation,22 this relation can be written as ph + [SnIn +] + [VO+] = [Oi−] + ne

Figure 8. Possible partial crystal structures of O−-ion-irradiated ITO films at various stages of the irradiation along with plots of ne and μH: [A] before the irradiation, [B] at an early stage of the irradiation (tirr = 2 min), and [C] at a stage after a long O−-ion irradiation. The possible chemical reactions from [A] to [B] and from [B] to [C] are also shown.

Table 2. Ionic Radii of In3+ and Sn4+ for Various Coordination Numbers23

(5)

where ph is the density of positive holes, and the square brackets denote the concentrations of ionized donors and acceptors in the ITO film. SnIn+ and VO+ represent a positively charged Sn ion at an In site (singly ionized) and a positively charged vacancy (relative to a perfect lattice) at an O site (singly ionized), respectively. Oi− denotes a negatively charged O ion at an interstitial site or VSO site (singly ionized). In this case, ne is very much larger than ph in eq 5. Note that doubly charged defects have resulting deep levels in the band gap that lead to strong lattice distortion in the neighborhood. Because there were no significant changes in the crystal structure for the O−-irradiated ITO films as shown in Figures 3 and 4, we conclude that all of ionized point defects in these films are singly charged. Figure 8 summarizes information on possible arrangement of In−Sn−O networks for the ITO films with various tirr, based on the relationship between ne and μH. The notations of the species (In, Sn, O(OO), VO+, VSO, and OVSO−) are specified in images of the crystal structure. The notations OO and OVSO− denote an O atom on an O (VO+) site and VSO site, respectively, in Kröger−Vink style.22 With increasing tirr, up to 2 min, that is, from [A] to [B] in Figure 8, we found a decrease in ne together with an increase in μH. If we take into account the fact that the ionized impurity scattering mechanism implies that μH is a decreasing function of ne, this indicates the suppression of ionized impurity scattering caused by a decrease in the number of the scattering centers, VO+, by 2.2 × 1020 cm−3 (i.e., from VO+-donor defects to OO upon the O−irradiation).8 Here, as shown in eq 5, we assume that the charge state of VO+ defects with a concentration of 2.2 × 1020 cm−3 is +1 (singly ionized), as described above. With further increasing tirr, that is, from [B] to [C] in Figure 8, μH decreased with decreasing ne. Note that, as shown in Table 2, the ion radius of In3+ for six-coordination (0.080 nm) is almost the

ion

coordination number

radius/nm

In3+

6 8 6 8

0.080 0.092 0.069 0.081

Sn4+

same as that of Sn4+ for eight-coordination (0.081 nm).23 This means that excess O atoms can be incorporated into the structural vacancies around Sn atoms to have a larger coordination number, such as seven or eight, without any structural change. The lack of the lattice-parameter variation and the crystal deformation upon the O−-ion irradiation was proven by the fact that XRD patterns with almost the same profiles were obtained irrespective of tirr as shown in Figure 3a−c, because this effect also maintains the point symmetry of the filled VSO sites. This behavior is illustrated in the crystal structure at the stage [C] of Figure 8. The excess O atoms incorporated into the In2O3 matrix will combine with singly ionized SnIn+ donors to energetically favor the formation of Sn−O complexes, such as [Sn2O] and [Sn2O4]. These complexes act as neutral scattering centers with a large scattering cross section,5,11 which results in a reduction of the carrier mobility of ITO films. In such ITO films at the stage [C], from eq 5, an increase in tirr resulted in a decrease in ne together with an increase in [Oi−] (in this case, [OVSO−]) and a very small constant value of [VO+]. Next, we briefly investigated the optical absorption at low ne. As shown in Figure 6, the observed Eg behavior should be reasonably consistent with that predicted from a trend in a previous report.20 The Eg of single crystal In2O3 is reported24 to be as high as 3.75 eV. Therefore, to consider the absorption edge of ITO films subjected to O−-ion irradiation, both a E


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ACS Applied Electronic Materials band-gap-narrowing effect25 ΔEg and a BursteinMoss shift26 ΔEBM must be taken into account. ΔEBM can be given by26,27

Author Contributions

Y.F. performed O− irradiation by RPD and measured all the physical properties. M.M. deposited ITO films by RPD. Y.F. and T.Y. wrote this manuscript. All authors reviewed the manuscript.

ΔE BM = (1 + m*e /m*h )(h2 /8m*e(3ne /π )2/3 − 4kBT ) (6)

Notes

where m*e, m*h, h, kB, and T are the electron- and holeeffective masses, Planck’s constant, Boltzmann’s constant, and the absolute temperature, respectively. From eq 6, the critical carrier concentration necri is defined by that for which ΔEBM (ne = necri) is 0. If the relations of m*e = 0.4m0 and m*h = 1.0m0 are assumed,28 where m0 is the mass of a free electron, and T is room temperature of 300 K, we calculated necri to be 3.7 × 1019 cm−3. This means that for the ITO films with an ne of less than 3.7 × 1019 cm−3, the band-gap-narrowing effect is dominant. Figure 5 shows the characteristics of α curves of the ITO films. We found the trade-off behavior of α curves between the fundamental absorption mode with a sharp increase in α toward the ultraviolet range and the absorption mode with α increasing slowly from various scattering centers such as ionized impurities, phonons, and neutral impurities, and freeelectron absorption toward infrared spectral range. For ITO films having the highest ne of 9.3 × 1020 cm−3, very strong freeelectron absorption limits λ-dependent α in the visible spectral range. An increase in λ from almost 500 nm to the infrared spectral range leads to an increase in α. In the visible λ ranging from 380 to 780 nm, with decreasing ne to 6.7 × 1019 cm−3, α tends to decrease at any given λ, whereas α increases at any given λ clearly with further decreasing ne to 1.2 × 1019 cm−3. This is probably due to a sharp increase in the concentration of SnIn+−Oi− complexes. Figure 5b,d clearly shows a red-shift of α for the ITO films with decreasing ne from 2.1 × 1020 to 6.7 × 1019 cm−3 and from 6.7 × 1019 to 1.2 × 1019 cm−3. In the λ range from 354 nm, corresponding to 3.5 eV, to 400 nm, a decrease in ne reduces α at any given λ. This ne-dependent α completely agrees with Eg behavior limited by ne.

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors gratefully acknowledge Dr. Nobuaki Takahashi, Mr. Toshiyuki Sakemi, and Mr. Hisashi Kitami of Sumitomo Heavy Industries, Ltd., for their helpful discussions during this study. This study was supported by collaboration with Sumitomo Heavy Industries, Ltd (Patent: WO2017/ 014278A1, 2017-01-26). ABBREVIATIONS

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REFERENCES

ITO, tin (Sn)-doped In2O3; TCO, transparent conducting oxide; VSO, structural defect(s); RPD, reactive plasma deposition; XRD, X-ray diffraction; XRR, X-ray reflectivity

(1) Ginley, D. S.; Bright, C. Transparent conducting oxides. MRS Bull. 2000, 25, 15−18. (2) Hamberg, I.; Granqvist, C. G. Evaporated Sn-doped In2O3 films: Basic optical properties and applications to energy-efficient windows. J. Appl. Phys. 1986, 60, R123−R159. (3) Calnan, S.; Tiwari, A. N. High mobility transparent conducting oxides for thin film solar cells. Thin Solid Films 2010, 518, 1839− 1849. (4) Buchholz, D. B.; Ma, Q.; Alducin, D.; Ponce, A.; Jose-Yacaman, M.; Khanal, R.; Medvedeva, J. E.; Chang, R. P. H. The Structure and properties of amorphous indium oxide. Chem. Mater. 2014, 26, 5401− 5411. (5) Gonzalez, G. B.; Mason, T. O.; Quintana, J. P.; Warschkow, O.; Ellis, D. E.; Hwang, J.-H.; Hodges, J. P.; Jorgensen, J. D. Defect structure studies of bulk and nano-indium-tin oxide. J. Appl. Phys. 2004, 96, 3912−3920. (6) Shigesato, Y.; Paine, D. C. Study of the effect of Sn doping on the electronic transport properties of thin film indium oxide. Appl. Phys. Lett. 1993, 62, 1268−1270. (7) Fan, J. C. C.; Bachner, F. J.; Foley, G. H. Effect of O2 pressure during deposition on properties of rf-sputtered Sn-doped In2O3 films. Appl. Phys. Lett. 1977, 31, 773−775. (8) Bierwagen, O.; Speck, J. S. High electron mobility In2O3 (001) and (111) thin films with nondegenerate electron concentration. Appl. Phys. Lett. 2010, 97, No. 072103. (9) Yamada, N.; Yasui, I.; Shigesato, Y.; Li, H.; Ujihira, Y.; Nomura, K. Donor Compensation and Carrier-Transport Mechanisms in Tindoped In2O3 Films Studied by Means of Conversion Electron 119Sn Mössbauer Spectroscopy and Hall Effect Measurements. Jpn. J. Appl. Phys. 2000, 39, 4158−4163. (10) Parida, B.; Gil, Y.; Kim, H. Highly Transparent Conducting Indium Tin Oxide Thin Films Prepared by Radio Frequency Magnetron Sputtering and Thermal Annealing. J. Nanosci. Nanotechnol. 2019, 19, 1455−1462. (11) Frank, G.; Köstlin, H. Electrical Properties and Defect Model of Tin-Doped Indium Oxide Layers. Appl. Phys. A: Solids Surf. 1982, 27, 197−206. (12) Shigesato, Y.; Paine, D. C.; Haynes, T. E. Study of the effect of ion implantation on the electrical and microstructural properties of tin-doped indium oxide thin films. J. Appl. Phys. 1993, 73, 3805− 3811. (13) Takahashi, Y.; Nitobe, K.; Uramoto, J.; Momose, T.; Ishimaru, H. Aluminum oxide thin film deposition by reactive ion plating using

5. CONCLUSION We have investigated the structural, electrical, and optical properties of polycrystalline ITO films and showed that ne, α, and n can be controlled by O−-ion irradiation. The film deposition and the generation and irradiation of O− ions were performed by RPD. We found that the O−-ion irradiation decreases the ne of the ITO films significantly, with little rearrangement of host In atoms and Sn dopants and with no generation of additional phase or precipitation. We identified the relationship between ne and μH during the O−-ion irradiation as a consequence of an annihilation of VO-donor defects and a subsequent filling of the VSO to form Sn−O complexes. Our key finding is that the state-of-the-art technology of O−ion irradiation is very effective in tailoring the physical properties of ITO films in a wide range while retaining the crystal structure of the films.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]; Phone: +81887-57-2180; Fax: +81-887-57-2181. ORCID

Yutaka Furubayashi: 0000-0001-8136-2553 Tetsuya Yamamoto: 0000-0002-7317-1252 F


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