Composition Controlled Superlattice InGaO3 (ZnO) m Thin Films by

Sep 8, 2010 - Ga Ordering and Electrical Conductivity in Nanotwin and Superlattice-Structured Ga-Doped ZnO. Sang-Won Yoon , Jong-Hyun Seo , Tae-Yeon ...
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DOI: 10.1021/cg100924a

Composition Controlled Superlattice InGaO3(ZnO)m Thin Films by Thickness of ZnO Buffer Layers and Thermal Treatment

2010, Vol. 10 4638–4641

Dong Kyu Seo, Bo Hyun Kong, and Hyung Koun Cho* School of Advanced Materials Science and Engineering, Sungkyunkwan University, 300 Cheoncheon-dong, Jangan-gu, Suwon, Gyeonggi-do 440-746, Korea Received July 13, 2010; Revised Manuscript Received August 20, 2010

ABSTRACT: Single crystal InGaO3(ZnO)m compound films with a superlattice structure for thermoelectric applications were fabricated on sapphire substrates with epitaxial ZnO buffer layers. Using a buffer layer with an appropriate thickness allows us to form InGaO3(ZnO)m films with complete superlattice structures along the c-axis at annealing temperatures below 1000 °C. The crystal phases of the InGaO3(ZnO)m films annealed at 900 °C were assigned to InGaO3(ZnO)1 and InGaO3(ZnO)2, depending on the thickness of the buffer layer and annealing time. When compared to the as-grown films, these InGaO3(ZnO)m films with superlattice structures exhibited both a lower electrical resistivity and higher Seebeck coefficient, due to the quantum confinement effect, resulting in improved thermoelectric properties.

1. Introduction Oxide semiconductors have attracted much attention, due to their high carrier mobility and high transmittance in the visible range. These properties suggest that multicomponent oxide (MO) semiconductors are a better choice of material than Si and organic materials for use as thin-film-transistors (TFTs) for next-generation displays with a large size and light weight, since the MO thin films have highly uniform amorphous phases and sufficient carrier mobility, despite their low growth temperature.1 Due to their increased scientific and technological importance, various MO semiconductors, such as InGaZnO, InZnO,2 InSnZnO,3 AlSnZnO,4 etc., have been actively investigated using magnetron sputtering, atomic layer deposition,5 and pulsed laser deposition.6 In addition, these MOs are among the most promising materials for solid-state gas and photosensor applications, since their electrical properties are highly sensitive to the preparation and environmental conditions.6 Therefore, from this point of view, the interest in transparent MOs with semiconducting properties is continuously increasing. A novel application field involving the use of semiconducting MOs is sustainable energy-resource thermoelectrics, operating in the high temperature region.7 Compared to conventional Bi based compounds, oxide thermoelectric materials have several advantages, including their high chemical and thermal stability and the controllability of their electrical conductivity, as well their being nontoxic green elements. Unfortunately, most previous research on oxide thermoelectrics has been limited to p-type semiconducting ones,8,9 whereas only a few n-type oxide based thin films have been studied, such as SrTiO3.10 Interestingly, it has been proposed that thermoelectric thin films with a periodic superlattice structure offer a powerful methodology to dramatically enhance the figure-of-merit of the thermoelectric (ZT) by reducing their thermal conductivity and increasing their Seebeck coefficient, which is achieved through the increase in the interfacial phonon scattering and quantum confinement effect, respectively.11,12 Thus, p-type Bi2Te3/Sb2Te3 and n-type TiO2/SrTiO3 superlattices were artificially produced by the *Corresponding author. Telephone: þ82 31 299 4733. Fax: þ82 31 290 7410. E-mail: [email protected]. pubs.acs.org/crystal

Published on Web 09/08/2010

elaborate control of the thickness, because the uniform periodicity of the superlattice is very important and leads to a significant enhancement in the thermoelectric figure-of-merit (ZT).13,14 Noticeably, Nomura et al. reported that single crystal MO InGaZnO thin films have a naturally formed superlattice structure along the c-axis following thermal treatment at high temperature (g1400 °C) and fabricated single crystal InGaZnO TFTs with high mobility.15,16 However, these single crystal InGaZnO films require unreasonably high annealing temperatures, and consequently, the use of relatively inexpensive commercial substrates is impossible. Until now, there have been few studies on the formation of the InGaZnO superlattice structure at reduced temperatures on commercial transparent substrates for thermoelectric applications. In this article, single crystal InGaZnO thin films with long-range periodicity were fabricated on sapphire substrates by sputtering growth using epitaxial ZnO buffer layers and thermal annealing at temperatures below 1000 °C. In addition, the ZnO fraction in the InGaO3(ZnO)m films was conveniently controlled by adjusting the thickness of the ZnO buffer layers and annealing time. 2. Experimental Section The oxide films in the present study were prepared on c-sapphire substrates by radio frequency magnetron sputtering. Prior to the deposition of the InGaZnO films, ZnO was grown as a buffer layer with thicknesses ranging from 10 to 200 nm. In particular, the growth temperature and argon/oxygen flow rate of the buffer layers were set to 700 °C and 20/10 SCCM, respectively, in order to form single crystal ZnO films with a low defect density epitaxial to the sapphire substrates, which was confirmed by the phi scan of X-ray diffraction (XRD) showing 6-fold symmetry. This was followed by the growth of the MO InGaZnO films at 700 °C with an argon/oxygen flow rate of 28/2 SCCM. The 4-in. high-purity InGaZnO sintered with a ratio of In/Ga/Zn = 1:1:1 was used as a target material for RF sputtering. To induce the crystallization of the InGaZnO thin films, postannealing was performed at 900 °C for 3-9 h under an air atmosphere. The bare sapphire substrates were loaded on the surface of the thin films to prevent the evaporation of the elements of the films. The variation in the crystal phases of the InGaZnO films was examined by XRD. The electrical and optical properties of the films were also characterized by a Hall-effect measurement system and UV-vis transmittance measurements, respectively. r 2010 American Chemical Society

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3. Results and Discussion Figure 1 shows the XRD results of the InGaZnO films grown on single crystal ZnO buffer layers with thicknesses of 10, 50, 100, and 200 nm, where all of the samples were thermal-annealed at 900 °C for 3 h; for comparison, the XRD spectrum of the as-grown sample with a 50 nm thick buffer layer is shown in Figure 2. Before thermal annealing, the XRD

Figure 1. Logarithmic plots of (a) θ-2θ scan XRD patterns for the as-grown InGaZnO/ZnO film and the InGaO3(ZnO)m films, with buffer layers having different thicknesses annealed at 900 °C for 3 h, and (b) narrow scan XRD patterns for the thermal annealed samples.

Figure 2. Logarithmic plot of θ-2θ scan XRD patterns for the InGaO3(ZnO)m film with a 50 nm thick buffer layer annealed at 900 °C for 9 h.

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pattern of the as-grown InGaZnO films grown at 700 °C on single crystal ZnO represents the partially crystallized film in the amorphous matrix. It can be seen that the peaks corresponding to only ZnO and sapphire substrate are clearly discriminated and patterns with a shoulder shape are faintly observed in the 2θ range of 20-30°. On the contrary, all of the samples thermal-annealed at 900 °C for 3 h show additional diffraction peaks from the InGaZnO films, as shown in Figure 1a. This means that the InGaZnO films are strongly crystallized by the thermal annealing. Unlike the other samples with several strong additional peaks, the InGaZnO film with the thickest ZnO buffer layer (200 nm) shows additional weak peaks. In particular, the diffraction patterns generated by thermal annealing in the InGaZnO films with the 10, 50, and 100 nm thick ZnO buffer layers exhibit a periodic arrangement with a constant separation. The distance between the periodic peaks (indicated by blue circles) in the InGaZnO films with the 10 and 50 nm thick ZnO buffer layers is estimated to be Δ(2θ) ∼ 10.3°, while it is much smaller in the sample with the 100 nm buffer layer (indicated by red triangles) and corresponds to Δ(2θ) ∼ 7.9°. Such a periodic distance between the diffraction peaks is generally due to the formation of superlattice structures with a regular atomic arrangement in one direction. However, the relationship between the InGaO3-ZnO related phases [InGaO3(ZnO)m] and diffraction peaks is likely to be complex, due to the very similar peak positions, despite the different ZnO contents, based on the JCPDS cards. To identify the exact InGaO3(ZnO)m phases in the superlattice structure, the narrow region diffraction patterns for the thermal annealed samples were investigated, as shown in Figure 1b. The dotted vertical lines in Figure 1b correspond to m = 1, 2, 3, 4, 5, 6, 7, and 15 from the left to right-hand sides. It can be clearly seen that the thermal annealing induces the formation of InGaO3(ZnO)1 superlattice phases in the samples with the 10 and 50 nm thick ZnO buffer layers and an InGaO3(ZnO)2 superlattice phase in the sample with the 100 nm thick buffer layer. On the other hand, the sample with the thickest buffer layer of 200 nm forms the InGaO3(ZnO)3 phase during annealing, whereas there is no sign of the formation of a significant superlattice. Consequently, the blue circles in Figure 1a correspond to the (003), (006), (009), ... planes of the InGaO3(ZnO)1 phases and the red triangles correspond to the (002), (004), (006), ... planes of the InGaO3(ZnO)2 phases. Interestingly, the ZnO peaks in the sample with the thinnest ZnO buffer layer of 10 nm disappear after thermal annealing, which means that the entire InGaZnO and ZnO films are completely intermixed in the InGaO3(ZnO)1 superlattice phase. However, increasing the thickness of the ZnO buffer layers results in increasing the

Figure 3. (a) Transmission spectra for the as-grown InGaZnO/ZnO film and the InGaO3(ZnO)m film with a 50 nm thick buffer layer annealed at 900 °C for 9 h. (b) R2 (R: absorption coefficient) versus E plot.

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intensity of the ZnO diffraction peaks in the annealed samples. This indicates that the thick buffer layers retained their ZnO phase near the sapphire substrate and the top region crystallized into the InGaO3(ZnO)m phase. Therefore, thermal annealing induces the crystallization of the InGaO3(ZnO)m films and the intermixing between the InGaZnO film and the ZnO buffer at the same time. Thus, the increase in the m value in the InGaO3(ZnO)m films with the thicker ZnO buffer layers is due to the sufficient supply of the ZnO from the buffer layer into the InGaZnO. Unlike in previous reports, where a high annealing temperature of above 1400 °C was required,16 our finding explains that the InGaZnO films can be crystallized into single crystal InGaO3(ZnO)m at temperatures below 1000 °C using the single crystal ZnO buffer layers. Considering the fact that the XRD peak intensity of the InGaO3(ZnO)m compound decreases with increasing thickness of the buffer layer, the crystallinity of the thermal annealed films exhibits the best quality in the samples, with buffer layers having thicknesses of 10 and 50 nm. Figure 3a shows the XRD spectra of the InGaZnO film with the 50 nm buffer layer after thermal annealing at 900 °C for 9 h, and further data are shown in Table 1. As compared with the sample annealed for 3 h (Figure 1), Figure 3a illustrates that the distance between the periodic superlattice peaks is significantly reduced from Δ(2θ) ∼ 10.3 to 7.9° and the peak with the highest intensity was shifted toward a higher angle (∼1°) by the longer thermal annealing. This result illustrates the increase of the m value to 2, that is, the change in the superlattice phase from InGaO3(ZnO)1 to InGaO3(ZnO)2 compound. In addition, it is observed that the ZnO diffraction peak at ∼34.48° completely disappears, due to the formation of single phase InGaO3(ZnO)2. Further annealing does not change the composition of the InGaO3(ZnO)m Table 1. Electrical Properties and Seebeck Coefficients of the As-Grown InGaZnO/ZnO Film and the InGaO3(ZnO)m Film with a 50 nm Thick Buffer Layer Annealed at 900 °C for 9 h carrier conc (cm-3) as-dep -1.3  l018 9h -3.5  l019

mobility (cm2/V 3 s) 31.0 3.3

resistivity (Ω 3 cm) 0.158 0.054

Seebeck coefficient (μV/K) -158 -192

Seo et al.

compounds. The intensity of the InGaO3(ZnO)2 diffraction peak at 31.5° is significantly improved by a factor of 5  105, compared with that of the InGaO3(ZnO)2 film annealed for 3 h. We believe that the enhanced diffraction intensity for the sample annealed for 9 h is due to the well ordered superlattice structure of the InGaO3(ZnO)2 compounds resulting from the longer thermal annealing and the appropriate thickness of the buffer layer. Summarizing the above-noted observations, we can conclude that the quaternary InGaZnO thin films are well crystallized by the assistance of the ZnO buffer layers and the thermal annealing at 900 °C and that the final composition of the InGaO3(ZnO)m compound is determined by the combination of the thickness of the ZnO buffer layers and the thermal annealing time. To study the electrical properties of the InGaO3(ZnO)m compounds formed by thermal annealing, we employed a Hall-effect system to measure the electrical resistivity and carrier concentration at room temperature. The as-grown InGaZnO/ZnO film with a 50 nm thick buffer layer showed an electrical resistivity of 0.158 [Ω 3 cm] and a carrier concentration of 1.3  1018 [cm-3]. On the other hand, the sample annealed for 9 h, showing the complete disappearance of the ZnO buffer layer, exhibits a lower electrical resistivity (0.054 [Ω 3 cm]), which is attributed to the increase in carrier concentration. A decrease in the Seebeck coefficient (S) was generally expected in the samples with increasing carrier concentration.17 However, the single crystal InGaO3(ZnO)m compound formed by the 50 nm thick ZnO buffer layer and thermal annealing at 900 °C for 9 h shows an enhanced S value of -192 μV/K (at 345 K), in spite of the increase in its carrier concentration, compared to that of the as-grown sample (-158 μV/K). Such an increase in the Seebeck coefficient is due to the quantum confinement induced by the superlattice structure, resulting in a larger electronic density of states.18 The optical band gap of InGaO3(ZnO)m compounds has been reported to range from 3.5 to 3.8 eV for the single crystal and to be ∼3.2 eV for the amorphous phase.19,20 Figure 3b shows the optical transmission spectra of the as-grown InGaZnO/ZnO film and the single crystal InGaO3(ZnO)m film (50 nm thick buffer layer) formed by thermal annealing at 900 °C for 9 h. It is well-known that the optical transparency

Figure 4. Schematic models for the formation of the superlattice InGaO3(ZnO)m films, with ZnO buffer layers having different thicknesses during thermal annealing: (a) thin buffer; (b) thick buffer.

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of oxide films is closely related to their surface roughness, band gap, and crystalline quality. The thermally annealed InGaO3(ZnO)m film consisting of a complete single phase shows a considerably improved optical transparency of above 90% in the visible and infrared regions, while the as-grown sample has an average transmittance of 70%. The thermal annealing at 900 °C for 9 h slightly roughens the surface morphology (not shown here). Nevertheless, the enhancement in the optical transmission of the sample with the rough surface is attributed to the improvement of its crystalline characteristics by the formation of a superlattice structure with a well ordered atomic arrangement. For the as-grown sample, the incident photons entering into the oxide films were actively scattered in the boundaries of the crystallites embedded in the amorphous matrix, but a well ordered atomic arrangement, such as a superlattice structure, suppresses the light scattering in the single crystallite films, resulting in good optical transparency. The band gaps of the as-grown sample and the InGaO3(ZnO)2 film are estimated to be 3.9 and 4.1 eV, respectively, resulting in the increase in the optical band gap through the formation of the single crystal compound. 4. Conclusions In conclusion, to fabricate single crystal InGaO3(ZnO)m thin films with a superlattice structure periodic along the c-axis for use in transparent thin film thermoelectric devices, we proposed a novel method using a ZnO buffer layer with an appropriate thickness and thermal treatment. Unlike in the previous studies using a high temperature process (g1400 °C), we fabricated single crystal InGaO3(ZnO)m films at temperatures below 1000 °C, which allows the use of sapphire substrates. In addition, it is possible to conveniently control the composition of the InGaO3(ZnO)m phases by adjusting the thickness of the ZnO buffer layer and the annealing time. Like the models shown in Figure 4, the initial InGaZnO films on the thin ZnO buffer layer changed into the InGaO3(ZnO)1 single phase by annealing them for a short time at 900 °C. However, the use of a thicker buffer layer induced the formation of a mixture of phases including ZnO and InGaO3(ZnO)m and required additional thermal treatment for the formation of the single phase InGaO3(ZnO)m compound. All of the observed single phase InGaO3(ZnO)m compounds showed a single crystal and well ordered superlattice structure. Further increasing the thickness of the buffer layer disturbed the formation of the superlattice structure, due to the excess supply of ZnO elements.

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Acknowledgment. This work was financially supported by the Pioneer Research Center Program through the National Research Foundation of Korea (Grant No. 2009-0078876 and Grant No. 2010-0002231) funded by the Ministry of Education, Science and Technology (MEST). This research was also financially supported by the Ministry of Knowledge Economy (MKE) and Korea Industrial Technology Foundation (KOTEF) through the Human Resource Training Project for Strategic Technology.

References (1) Kim, M. K.; Jeong, H. J.; Lee, H. J.; Ahn, T. K.; Shin, H. S.; Park, J. S.; Jeong, J. K.; Mo, Y. G.; Kim, H. D. Appl. Phys. Lett. 2007, 90, 212114. (2) Jung, Y. S.; Seo, J. Y.; Lee, D. W.; Jeon, D. Y. Thin Solid Films 2003, 445, 63. (3) Choi, K. H.; Nam, H. J.; Jeong, J. A.; Cho, S. W.; Kim, H. K.; Kang, J. W.; Kim, D. G.; Cho, W. J. Appl. Phys. Lett. 2008, 92, 223302. (4) Cho, D.; Yang, S.; Byun, C.; Shin, J.; Ryu, M. K.; Ko, S. H.; Hwang, C. S.; Chung, S. M.; Cheong, W. S.; Yoon, S. M.; Chu, H. Y. Appl. Phys. Lett. 2008, 93, 142111. (5) Carcia, P. F.; McLean, R. S.; Reilly, M. H. Appl. Phys. Lett. 2006, 88, 123509. (6) Fung, T. C.; Chuang, C. S.; Nomura, K.; Shieh, H. D.; Hosono, H.; Kanicki, J. J. Inf. Display 2008, 9, 4. (7) Ohta, H.; Seo, W. S.; Kumoto, K. J. Am. Ceram. Soc. 1996, 79, 2193. (8) Wang, A. S.; Guo, S. V.; Chen, K.; Li, Q.; Soukiassian, A.; Schlom, D. G.; Katz, M. B.; Pan, X. Q.; Winnie, W. N.; Vaudin, M. D.; Xi, X. X. Appl. Phys. Lett. 2009, 94, 022110. (9) Ohta, H.; Mizutani, A.; Sugiura, K.; Hirano, M.; Hosono, H.; Koumoto, K. Adv. Mater. 2006, 18, 1649. (10) Ohta, H.; Kim, S.; Mune, Y.; Mizoguchi, T.; Nomura, K.; Ohta, S.; Nomura, T.; Nakanishi, Y.; Ikuhara, Y.; Hirano, M.; Hosono, H.; Koumoto, K. Nat. Mater. 2007, 6, 129. (11) Kim, W.; Zide, J.; Gossard, A.; Klenov, D.; Stemmer, S.; Shakouri, A.; Majumdar, A. Phys. Rev. Lett. 2006, 96, 045901. (12) Capinski, W. S.; Cardona, M.; Katzer, D. S.; Maris, H. J.; Ploog, K.; Ruf, T. Physica B 1999, 263, 530. (13) Zheng, Z.; Xiao, B.; Chhay, B.; Zimmerman, R. L.; Edwards, M. E.; Ila, D. Surf. Coat. Technol. 2009, 203, 2682. (14) Wang, Y.; Lee, K. H.; Ohta, H.; Koumoto, K. J. Appl. Phys. 2009, 105, 103701. (15) Nomura, K.; Ohta, H.; Ueda, K.; Kamiya, Y.; Orita, M.; Hirano, M. J. Appl. Phys. 2004, 95, 10. (16) Nomura, K.; Ohta, H.; Ueda, K.; Kamiya, T.; Hirano, M.; Hosono, H. Science 2003, 300, 1269. (17) Snyder, G. J.; Toberer, E. S. Nat. Mater. 2008, 7, 105. (18) Hicks, L. D.; Dresselhaus, M. S. Phys. Rev. B 1993, 47, 12727. (19) Lee, K. W.; Heo, K. Y.; Kim, H. J. Appl. Phys. Lett. 2009, 94, 102112. (20) Ohta, H.; Huang, R.; Ikuhara, Y. Phys. Status Solidi (RRL) 2008, 2, 105.