The Optimum Fabrication Condition of p-Type Antimony Tin Oxide Thin Films Prepared by DC Magnetron Sputtering

Dang, Huu Phuc; Luc, Quang Ho; Le, Tran; Le, Van Hieu

doi:https://doi.org/10.1155/2016/7825456

Journal of Nanomaterials

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Abstract Introduction Results and Discussion Conclusions Acknowledgments References Copyright Related Articles

Research Article | Open Access

Volume 2016 | Article ID 7825456 | https://doi.org/10.1155/2016/7825456

The Optimum Fabrication Condition of p-Type Antimony Tin Oxide Thin Films Prepared by DC Magnetron Sputtering

Huu Phuc Dang,¹Quang Ho Luc,²Tran Le,³and Van Hieu Le⁴

Academic Editor: Zhaoyao Zhan

Received14 Mar 2016

Accepted08 May 2016

Published23 May 2016

Abstract

Transparent Sb-doped tin oxide (ATO) thin films were fabricated on quartz glass substrates via a mixed (SnO₂ + Sb₂O₃) ceramic target using direct current (DC) magnetron sputtering in ambient Ar gas at a working pressure of 2 × 10⁻³ torr. X-ray diffraction (XRD), X-ray photoelectron spectroscopy (XPS), Hall-effect, and UV-vis spectra measurements were performed to characterize the deposited films. The substrate temperature of the films was investigated in two ways: (1) films were annealed in Ar ambient gas after being deposited at room temperature or (2) they were deposited directly at different temperatures. The first process for fabricating the ATO films was found to be easier than the second process. The deposited films showed p-type electrical properties, a polycrystalline tetragonal rutile structure, and their average transmittance was greater than 80% in the visible light range at the optimum annealing temperature of 500°C. The best electrical properties of the film were obtained on a 10 wt% Sb₂O₃-doped SnO₂ target with a resistivity, hole concentration, and Hall mobility of 0.55 Ω·cm, 1.2 × 10¹⁹ cm⁻³, and 0.54 cm²V⁻¹s⁻¹, respectively.

1. Introduction

Optoelectronic devices are found in many areas of society, from simple household appliances and multimedia systems to communications, optical monitoring, and medical instruments. The majority of optoelectronic devices are LEDs, laser diodes, photodiodes, and solar cells. Transparent conducting oxides (TCOs), which required as electrodes in these above optoelectronic devices, are made of Sn-doped In₂O₃ (ITO) [1], Al-doped ZnO (AZO) [2], Ga-doped ZnO (GZO) [3], Sb-doped SnO₂ (ATO) [4], or F-doped SnO₂ (FTO) [5], which have low resistivities (10⁻³–10⁻⁴ Ω·cm) and high transparencies (80–90%). However, research on p-type TCOs has only recently begun.

The first candidate is p-type delafossite; CuY_1−xCa_xO₂, CuScO_2+x, CuCr_1−xMg_xO₂, and CuScO₂, as well as AgFeO₂ [6–11] …, have low resistivities but a present limitation in stability and have low transparencies, narrow band gaps, and difficult compositions. The second candidate is ZnO; however, it is difficult to produce p-type ZnO due to self-compensation effects [12–14]. Compared with ZnO, SnO₂ is a wide band gap semiconductor (3.6–4 eV) [15, 16] with good chemical, mechanical, and thermal stabilities [17–19]. Additionally, tin can easily be doped by group III metals because tin has a valence of 4.

Recently, the p-type doping of SnO₂ thin films was studied using metal elements in group III. For example, Al [20, 21] and Ga [22] have been incorporated into the structure as the acceptor impurities. Later, Ji et al. [23] and Ni et al. [24] also obtained p-type SnO₂ thin films doping In and Sb. Besides, Ni et al. [25] found that the incorporation of group IIA atoms (Zn) was also suitable for doping SnO₂, and Mao et al. [26] reported an In-Ga codoping method to realize p-type SnO₂ thin films with a high mobility.

The previous research has been focused on thin films synthesized by RF magnetron sputtering [20, 22, 24–30], the sol-gel [31, 32], spay-pyrolysis [33], or PLD [34]. Among these deposition methods, RF magnetron sputtering method is widely used because of high deposition rate, good adhesion, and easy control of the electrical properties of the films by controlling processing parameters such as sputtering power, pressure, and temperature, particularly postannealing temperatures. However, no report explained clearly how to achieve the p-type by annealing and why thin films were not directly prepared with the various temperatures as well as why the thin films were deposited at 200°C after postannealing. The results of the previous research have not yet solved this problem, and, furthermore, a DC magnetron sputtering process has not yet been studied. Because the ionic radius of the Sb³⁺ ion (0.76 Å) is similar to that of Sn⁴⁺ (0.71 Å), we studied optical, electrical properties and the crystal structure of Sb-doped SnO₂ films deposited on quartz substrates by DC magnetron sputtering from ceramic targets made of a mixture of SnO₂ and Sb₂O₃. The parameters investigated were the substrate temperature during fabrication, the annealing temperature, and the percentage of the Sb₂O₃ dopant.

2. Experimental Techniques

SnO₂:Sb (ATO) thin films were deposited on quartz substrates from a ceramic target composed of a mixture of (SnO₂ + Sb₂O₃) by DC magnetron sputtering in ambient Ar gas at a working pressure of 2 × 10⁻³ torr with the Sb₂O₃ (wt%) varying from 0% to 15%, on a Univex 450 system. The quartz glass substrates were ultrasonically cleaned in 10% NaOH/acetone/deionized water solution for 35 min. Before deposition, the target was presputtered for 15 min to remove any contaminants on the target surface. The sputtering power and the target to sample distance were 15 W and 7 cm, respectively.

The substrate temperature of the films was investigated in two ways: (1) the films were deposited directly at different temperatures from room temperature to 500°C and heating and cooling rate 20°C/min, respectively, and (2) they were annealed in ambient Ar gas after being deposited at room temperature with the same heating and cooling rate as part 1.

The thicknesses of the thin films were measured using the SCOUT program by fitting the modeled spectrum to the experimental spectrum. By the controlling the deposition time, the thickness of SnO₂ and ATO films were about 400 nm. The structure of the ATO films was studied using an X-ray diffractometer (D8 ADVANCE) operated at 40 kV and 40 mA with Cu–Kα radiation. The optical transmittance was measured in the wavelength range 200–1100 nm using a UV-vis Jasco V-530. The film resistivity was measured at room temperature using the four-point probe method. A Hall measurement system (HMS 3000), which operated at room temperature, was used to measure the carrier concentration and the mobility. X-ray photoelectron spectroscopy (XPS) was measured at room temperature on an ESCALAB and the I-V characteristics were measured using a semiconductor testing system (Keithley 2400).

3. Results and Discussion

3.1. Structural Properties

Figures 1 and 2 show the XRD patterns of the SnO₂ and the ATO films (15% ATO films deposited from ATO target with 15% wt Sb₂O₃) deposited directly on quartz substrates at different temperatures. In particular, the temperature was applied directly to the films during the fabrication process; however, S6 and A6 films were deposited at 500°C followed by annealing at 600°C because the evaporation rate from substrates is greater than the condensation rate on substrates at temperatures above 500°C. Figure 1 shows that all films made at room temperature had an amorphous structure; the films started to crystallize at 200°C with a SnO₂ tetragonal rutile structure (JCPDS Number 41-14445). The films had better crystalline quality at approximately 400°C marked by the appearance of the (110) peak, and the intensity of the peak sharply increased for the films deposited at 500°C and annealed at 600°C. Figure 2 shows that ATO films started to crystallize at 200°C with a SnO₂ tetragonal rutile structure and no peaks from antimony or antimony oxide phases appear in the XRD patterns, indicating that antimony could be incorporated into the tin oxide lattice. Several authors also showed all SnO₂ films doping atoms such as Ta [35], In-Ga [26], Sb [24, 34], Ga [22], Zn [28], and In [31] are in rutile phase without any secondary phases of dopant atoms or their oxides, indicating the incorporation of dopants into host lattice. Figure 2 also shows that there is the appearance of Sn₂O₃ phase (101) for films deposited at 500°C and the films deposited at 500°C followed by annealing at 600°C had one more Sn₂O₃ phase . In particular, the SnO₂ phase (110) disappeared and phase (101) appeared for ATO 500°C followed by annealing 600°C while it contrasted to SnO₂ films. This could also be the substitution of Sn⁴⁺ by Sb³⁺. In fact, the report [36, 37] had also the same results that the (110) surface transformed to (101) surface which is rich in Sn²⁺ atoms (i.e., rich oxygen vacancy) as Sn⁴⁺ were substituted by Sn²⁺.

The above results indicate that the ATO films can be made p-type through the annealing step after deposition at room temperature. Therefore, to avoid evaporating the material on the substrate, which occurs when the film is deposited at a higher temperature, the film can be deposited at room temperature and then annealed at various temperatures.

The XRD patterns of the as-deposited SnO₂ thin films and those annealed below 550°C are shown in Figure 3. The films become polycrystalline and are in the rutile phase with (110), (101), and (211) orientations. When the SnO₂ films were annealed at 550°C and 600°C, the Sn₂O₃ phase appeared and can contribute to the positive carriers, which result in the p-type films. It shows clearly in electrical properties (Table 2).

Figure 4 shows that the peak intensity of the ATO films is lower than those of the SnO₂ films under the same conditions, indicating that the dopant Sb³⁺ ions could replace the Sn⁴⁺ ions sites. When annealed above 500°C, the Sb³⁺ are activated to substitute Sn⁴⁺ significantly and there is appearance of Sn₂O₃ phase which is not recorded at SnO₂ films of the same fabricated conditions. The energy released from the replacement of Sn⁴⁺ by Sb³⁺ also contributes to the oxygenation of Sn₂O₃ to SnO₂ when annealed at 600°C. This oxygenation does not occur at SnO₂ films of the same fabricated conditions. The ATO films have lower resistivities (Table 2) with the presence of the Sn₂O₃ phase at an annealing temperature of 550°C. However, to confirm whether the positive contribution of Sn³⁺ is significant, the ATO films were annealed for 2 hours.

Figure 5 shows that the reflection intensity for (110), (101), and (211) increased versus annealing time, showing the high crystalline nature of the ATO films. Sn₂O₃ phase (101) disappeared at 500°C and appeared at 600°C for 2 h compare to 1 h. Besides, ATO annealed at 550°C for 2 h had one more Sn₂O₃ phase being (030). All transformation from Sn₂O₃ phase to SnO₂ phase and vice versa could be the attribution to films energy released from the substitution of Sn⁴⁺ by Sb³⁺ as mentioned in the discussion of Figure 4. Next, we investigated the effect of the doping percentage on the structure of the ATO films.

Figure 6 illustrates the XRD patterns of the as-deposited samples annealed for 2 h at 500°C with various Sb₂O₃-doping percentages. The 5% and 10% ATO films do not show the presence of a Sn₂O₃ phase as 15% ATO films. Thus, the number of Sb³⁺ ions substituting for Sn⁴⁺ ions was enough to supply the energy required for the oxidation process of Sn₂O₃ transforming to SnO₂ at minimum doping percentages of 5 wt% Sb₂O₃.

3.2. Electrical Properties of ATO

As seen in Table 1, all of the SnO₂ films deposited at various temperatures including sample S6 exhibit n-type conductivity due to oxygen vacancies and tin interstitials. This result has been also given in [38, 39] at a certain deposition temperature. On the other hand, the samples A5 and A6 of 15% ATO films showed p-type conductivity. Clearly, antimony atoms substituted for tin sites at 500°C temperature and this substitution was recorded at the deposition temperatures of 600 and 700°C as mentioned in [34]. In particular, the carrier concentration increased from 10¹⁵ to 10¹⁶, and the carrier mobility decreased from 19.3 to 3.49 after annealing at 600°C following the direct deposition at 500°C. The decrease in the carrier mobility is attributed to the increase in ionized acceptors as mentioned in [24]. Therefore, it is obvious that the annealing method plays an important role in the electrical properties of the ATO films.

As seen in Table 2, the SnO₂ and ATO films were nonconductive at room temperature and were n-type at 200°C; however, the ATO films had much lower resistivities compared to the SnO₂ films. This effect might also be explained as follows: creating donors such as oxygen vacancies, tin interstitials as mentioned in [23, 25, 28, 31, 32, 38, 40], and unintentional hydrogen dopants as in agreement with previous theoretical and experimental studies [41–44] compensate by the Sn²⁺ acceptors [29, 32] at room temperature, leading to nonconductive SnO₂ and ATO films, while Sb appear as a result of the dissociation of Sb₂O₃ at 200°C, which supplies more negative carriers for the ATO films leading to a resistivity of the ATO films lower than that of the SnO₂ film. This dissociation of Sb₂O₃ shows clearly at optical properties (Figure 10). The SnO₂ films annealed at a temperature in the range of 300–500°C are nonconductive because Sn⁺² is oxidized to Sn⁴⁺ [45], the oxygen vacancies are filled with the oxygen present in the ambient air, and unintentional hydrogen donors are driven out of films [43, 46]. The 15% ATO films annealed at 300°C are n-type, whereas, in the range of 400–600°C, the 15% ATO films show p-type conductivity because the Sb³⁺ substitutes for Sn⁴⁺. Several authors have also observed this phenomenon that the conduction type of In-Ga co- [26], Sb- [24], Ga- [27], Zn- [25, 28], or In-doped SnO₂ [23] changed from n-type to p-type at a certain annealing temperature. The SnO₂ films annealed at 550°C are p-type, which is attributed to the dissociation from SnO₂ to Sn₂O₃ being in agreement with results of structure analyzing. However, the resistivity of SnO₂ film is much higher than that of the 15% ATO film annealed at 550°C. It was found that hole carriers were formed by Sb³⁺ replacing Sn⁴⁺ rather than by Sn³⁺ replacing Sn⁴⁺. Indeed, the hole concentration of the SnO₂ films decreased at 550°C because the Sn₂O₃ phases were oxidized and this also increased the mobility. Contrarily, the hole concentration of the 15% ATO films first increased with increases in the annealing temperature, and the highest hole concentration was obtained after annealing at 550°C for 1 h. As the annealing temperature was further increasing, the hole concentration value started to decrease because of disappearing of acceptor Sn³⁺ and native donors being able to appear. The 15% ATO films annealed at 550°C for 1 h had low resistivities (0.83 Ω·cm) and high hole concentrations (4.4 × 10¹⁸ cm⁻³). However, in the following section, we increased the annealing time of the ATO films to confirm the effect of the Sn³⁺ on the electrical properties.

As we see in Table 3, the resistivity increased slightly due to decreasing hole mobility as samples were annealed at 550°C for 2 h compared to 1 h regardless existence of Sn₂O₃ phase in films. Contrarily, the resistivity decreases as samples were annealed at 500°C for 2 h compared to 1 h without existing Sn₂O₃ phase in films. In conclusion, the contribution of the hole concentration is primarily from the substitution of Sn⁴⁺ by Sb³⁺, while the contribution from Sn³⁺ is minor. The other issue, we determined the effect of the doping percentage on the electrical properties.

As seen in Table 4, the resistance first decreases with increasing Sb concentration, and the lowest resistance is obtained at a doping concentration of 10 wt%. The 10% ATO films had the lowest resistivity (0.55 Ω·cm) and the highest hole concentration (1.2 × 10¹⁹ cm⁻³). As the wt% Sb₂O₃ was further increased, the resistance value started to increase. Our present observations suggest that beyond a certain doping concentration, the doping atoms do not occupy the lattice sites but instead produce some types of defects.

3.3. Optical Properties

At room temperature, both SnO₂ and ATO films have absorption edge in visible light at 431 nm (2.87 eV) for SnO₂ films and 350 nm for ATO films; this indicates existence of defect level Sn²⁺ [47] which served as an acceptor; electron at acceptor energy level Sn²⁺ absorbed light energy so that it transited to conduction band (Figure 7); this transited energy differ from exciton transited energy. At 200°C, the absorption edge significantly shifts to the shorter wavelength side; at above 200°C, this shifts was slight. This is due to acceptor Sn²⁺ transform to Sn⁴⁺, particular at 500°C band gap energy is exciton absorbed energy (Figures 7 and 8).

Figure 9 presents the transmittance spectra of SnO₂ films with different annealing temperatures. At room temperature, the transmittance spectra of SnO₂ films were taken from Figure 7. Above 200°C, acceptor Sn²⁺ were supplied energy to transform to Sn⁴⁺ so that the average transmittance in the visible region increased gradually above 80%; furthermore, the absorption edge gradually shifts to the shorter wavelength side. Meanwhile, Figure 10 shows that the ATO films have lower average transmittances compared to the SnO₂ films at temperatures below 400°C, which indicates the decomposition of Sb₂O₃ to form Sb which has discrete distribution in films, and this causes a decrease in the transmittance due to their absorption. It does not occur at ATO films deposited at 200°C. The transmittance in the visible region increased to 80% at 500°C. This effect occurs because, with increasing temperature, the doping Sb enters the film from interstitial sites and substitutes into other positions.

Figure 11 presents the spectral transmittance of 5%, 10%, and 15% ATO films with the same thickness annealed at 500°C for 2 h. The average transmissions of the 5%, 10%, and 15% ATO films in the visible region (300–800 nm) were approximately 80%. All these results are promising for diverse applications in the field of transparent electronics.

3.4. XPS Analysis

Figures 12 and 13 show that the results of the XPS spectra of the O 1s and the Sn core-level of the ATO films with different Sb concentrations were annealed at 500°C for 2 h after being deposited at room temperature and SnO₂ films were annealed at 550°C for 1 h after being deposited at room temperature. All peak positions were corrected for surface charge by the C 1s peak at 284.5 eV. The O 1s spectrum can be deconvoluted into three peaks for the SnO₂ (530, 531.5, and 532.5 eV) using XPSEAK41 software, and the profile of each peak was taken as a Lorentzian-Gaussian mixed function, respectively. For SnO₂ films, deconvolutions contained the 3 peaks of O 1s at 530, 531, and 532.5 eV. Two peaks at 530 and 531.5 eV can be attributed to O²⁻ ions in the rutile structure of the Sn⁴⁺ ion array [48] and the O²⁻ ions in oxygen-deficient regions [49]. The peak at 532.5 eV is assigned to the adsorbed oxygen species [50]. The above deconvolutions of all ATO films contained the 3 peaks of O 1s and 1 peak at 529.5 eV due to the contribution of the Sb³⁺ species [50]. This peak was attributed to the substitution of Sn⁴⁺ by Sb³⁺. It showed clearly that the peak for the Sn 3d_3/2 of ATO films shifted to a lower binding energy compared to that of SnO₂ films. The atomic percentages and compositional ratios of all the samples were calculated by the ratio between the integrated areas of O_531.5 and O₅₃₀ peaks and the integrated areas of the 2Sn and 1.5Sb peaks, and the results are listed in Table 5. In ideal stoichiometric SnO₂ films, the [O₅₃₀ + O_531.5]/(2[Sn] + 1.5[Sb]) ratio must have been 1, but the results in Table 5 give 1.427, indicating that valence state of Sn was not 4 perfectly and including lower valence state such as Sn³⁺. These results were in agreement with the existence of acceptor Sn³⁺ discussed in structure and electrical properties (3.1 and 3.2). For ATO, this ratio varied from 0.221–0.866, which indicates the existence of numerous oxygen deficiencies (the ratio was 1 for stoichiometric ATO films). The oxygen deficiencies prove substitution of Sn⁴⁺ by Sb³⁺. Moreover, the O 1s peak of 10 wt% Sb₂O₃ was asymmetric proving the highest substitution of Sn⁴⁺ by Sb³⁺, resulting in the highest concentration of holes (relative to the lowest ratio).

(a)

(b)

(c)

(d)

(a)

(b)

(c)

(d)

3.5. Properties of p-Type ATO/n-Type Si Heterojunction

To confirm the fabrication of p-type conductive ATO films, a p-type 10% ATO was selected as a p-type layer due to its high hole concentration of 1.2 × 10¹⁹, and they were fabricated on an n-type Si substrate. The current-voltage characteristics of the p-type ATO/n-type Si structure are plotted in Figure 14. Indium electrodes were placed both on the p-type and the n-type layers, as shown in the upper left inset. The lower right inset shows the I-V curves for a pair of In contacts on the p-type and n-type material layers. The linear I-V characteristics indicate good ohmic contact behavior. It is found that the I-V curve of the heterojunction p-type ATO/n-type Si is nonlinear. This characteristic clearly indicates the formation of p-n junction at the interface. Figure 14 also shows a forward turn-on voltage of 1.09 V.

4. Conclusions

In summary, p-type transparent conducting ATO films were prepared by DC magnetron sputtering. It was found that an annealing temperature of 500°C for 2 h was necessary for high p-type conductivity. The best electrical properties of the film were obtained from 10 wt% Sb₂O₃-doped SnO₂ target with a resistivity, hole concentration, and Hall mobility of 0.55 Ω·cm, 1.2 × 10¹⁹ cm⁻³, and 0.54 cm²V⁻¹s⁻¹, respectively. The XPS analysis indicated that Sb exists in the form of Sb³⁺ in the film and acts as an acceptor. The ATO films show polycrystalline tetragonal rutile structure and their average transmittance is above 80% in the visible light range. The p-n heterojunction diode composed of p-ATO/n-Si exhibits good I-V characteristics, thereby, demonstrating that a p-type ATO film and its fabrication technology might have a promising future in practical applications.

Competing Interests

The authors declare that there are no competing interests regarding the publication of this paper.

Acknowledgments

This research was funded by the Vietnam National University (VNU-HCM), Ho Chi Minh City, under Grant no. C2014-18-27.

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Copyright © 2016 Huu Phuc Dang et al. This is an open access article distributed under the Creative Commons Attribution License, which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

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