Oxidant-Dependent Thermoelectric Properties of Undoped ZnO Films

Article pubs.acs.org/cm

Oxidant-Dependent Thermoelectric Properties of Undoped ZnO Films by Atomic Layer Deposition Hyunho Kim, Zhenwei Wang, Mohamed N. Hedhili, Nimer Wehbe, and Husam N. Alshareef* Materials Science and Engineering, King Abdullah University of Science and Technology (KAUST), Thuwal, 23955-6900, Saudi Arabia S Supporting Information *

ABSTRACT: Extraordinary oxidant-dependent changes in the thermoelectric properties of undoped ZnO thin films deposited by atomic layer deposition (ALD) have been observed. Specifically, deionized water and ozone oxidants are used in the growth of ZnO by ALD using diethylzinc as a zinc precursor. No substitutional atoms have been added to the ZnO films. By using ozone as an oxidant instead of water, a thermoelectric power factor (σS2) of 5.76 × 10−4 W m−1 K−2 is obtained at 705 K for undoped ZnO films. In contrast, the maximum power factor for the water-based ZnO film is only 2.89 × 10−4 W m−1 K−2 at 746 K. Materials analysis results indicate that the oxygen vacancy levels in the water- and ozone-grown ZnO films are essentially the same, but the difference comes from Zn-related defects present in the ZnO films. The data suggest that the strong oxidant effect on thermoelectric performance can be explained by a mechanism involving point defect-induced differences in carrier concentration between these two oxides and a selfcompensation effect in water-based ZnO due to the competitive formations of both oxygen and zinc vacancies. This strong oxidant effect on the thermoelectric properties of undoped ZnO films provides a pathway to improve the thermoelectric performance of this important material.

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INTRODUCTION

Despite the superior performance of Bi2Te3 and its alloys, practical application such as a thermoelectric generator is still limited due to the near-RT operating range, thermal instability, and the high price of the raw materials. Recently, research on metal oxide thermoelectrics has accelerated due to their advantages such as thermal stability, nontoxicity, and abundance. Although thin films of ZnO have been investigated by various deposition methods for decades, there are few reports on the thermoelectric properties of ZnO films deposited by atomic layer deposition (ALD). These reports are mostly on the thermoelectric properties of Al-doped ZnO films by ALD.6,7 The Karppinen group has reported a unique inorganic−organic superlattice based on ALD ZnO and molecular layer deposition (MLD) hydroquinone, which can significantly reduce the thermal conductivity of the film.8−10 To the best of our knowledge, only one study has been reported on the electrical properties of ALD ZnO films using different oxidants; however, the electrical characterization was limited to RT.11 Thus, in this study, we focus on the role of ALD oxidizing agents and corresponding defect chemistry on the thermoelectric properties of undoped ALD ZnO films. By the term “undoped ZnO,” we mean that no foreign atoms or dopants (e.g., Al, Ga, In, etc.) have been intentionally

The field of energy harvesting has been attracting significant interest due to the limited availability of traditional energy sources and to help reduce global CO2 emissions. Among all existing energy harvest methods, the thermoelectric effect provides a direct solution for reversible conversion between thermal energy and electrical energy.1 The efficiency of thermoelectric performance is determined by a dimensionless figure of merit, ZT = σS2T/κ, where σ is electrical conductivity, S is Seebeck coefficient, T is absolute temperature, and κ = κL + κE is thermal conductivity, which consists of a lattice contribution (κL) and electron contribution (κE).2 According to the definition of ZT, a high thermoelectric power factor (σS2) with low thermal conductivity is essential to achieving outstanding thermoelectric performance. However, all three parameters that determine the ZT are interrelated; thus extensive research had been exerted to optimize the ZT, including a nanostructuring approach for suppressing the κL and controlling the carrier density by external doping.3 For the latter case, ionized dopants are often helpful to increasing overall thermoelectric efficiency, although σ and κL may simultaneously increase due to the increased carrier concentration.4 Studies on thermoelectric materials have accelerated since the discovery of high ZT ∼ 2.4 at room temperature (RT) from p-type Bi2Te3/Sb2Te3 superlattices by Venkatasubramanian et al.5 © 2017 American Chemical Society

Received: October 31, 2016 Revised: February 26, 2017 Published: February 27, 2017 2794

DOI: 10.1021/acs.chemmater.6b04654 Chem. Mater. 2017, 29, 2794−2802

Article

Chemistry of Materials

Figure 1. Schematic illustration of the ALD mechanism for the growth process of (a) ZnO-W and (b) ZnO-O. All films were grown on a glass substrate, first, cycle temperature dependent thermoelectric properties of ZnO-O and ZnO-W. (c) Electrical conductivity and (d) Seebeck coefficient. Drastically different properties of ZnO films are observed depending on the oxidant used (water vs ozone) during ZnO growth. Heating cycle and subsequent cooling cycle are presented with filled and open symbols, respectively.

introduced to the ZnO films. We find that the oxidant used during ALD reaction for undoped ZnO can increase the power factor by nearly a factor of 2. We further study the defect chemistry in the films by multiple methods, and a point-defect model is proposed to explain the observed changes in thermoelectric performance.

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Electrical contact was established by pin-contact of four Pt electrodes using metal springs. Carrier concentration and Hall mobility were measured by a Hall-effect measurement system (Lake Shore 7700A) using the van der Pauw technique at RT under a magnetic field range of ±5, 10, 15, and 20 kG. Ohmic contact was achieved by physical soldering of small indium dots at the corners of the sample and lead wires from a printed circuit board for the measurement (see the Supporting Information for details) Characterization. X-ray diffraction (XRD) patterns were collected with a Bruker diffractometer (D8 advance, AXS system, Germany) with Cu Kα radiation, λ = 1.5406 Å. Surface morphologies were observed by an FE-SEM (Nova Nano 630, FEI, USA) and MFP-3D AFM system (Asylum Research, USA) in the tapping mode. PL spectra were measured with a micro-Raman spectrometer (LabRAM Aramis, Horiba, Japan) equipped with a ThorLabs LMU-15x-NUV lens and a temperature controllable stage (THMS600, Linkam, UK), using a He−Cd laser with wavelength of 325 nm for excitation and liquid nitrogen for the cooling. UV−vis transmission spectra were measured with a UV−vis spectrometer (Evolution 600, Thermo Scientific, USA) with respect to the glass substrate. XPS studies were carried out in an Axis Ultra DLD spectrometer (Kratos Analytical, UK) equipped with a monochromatic Al Kα X-ray source (hν = 1486.6 eV) operating at 150 W, a multichannel plate, and delay line detector under a vacuum of 1−10−9 mbar. Binding energies were referenced to the C 1s binding energy of adventitious carbon contamination, which was taken to be 284.8 eV. The data were analyzed with commercially available software, CASAXPS. O 1s and Zn 2p core levels were fitted with a mixture of Gaussian (70%)− Lorentzian (30%) (GL30) function after Shirley type background subtraction. The depth profiling experiments were acquired on a Dynamic SIMS instrument (Hiden analytical, UK) operated under ultrahigh vacuum conditions, typically 10−9 Torr, using an argon ion

EXPERIMENTAL SECTION

Thin Film Deposition. Nanograined ZnO thin films were deposited on a glass substrate (Eagle 2000 AMLCD glass, Samsung Corning, Korea) at 200 °C using a Savannah ALD system (Cambridge Nanotech, USA). A 1 inch × 1 inch sized glass substrate was cleaned with a commercial glass cleaner, acetone, isopropanol, and deionized water followed by drying with high purity (5N) nitrogen gas. The ZnO ALD reaction sequence was (1) under a constant N2 flow at 20 sccm, (2) diethylzinc (Sigma-Aldrich, USA) dose for 0.015 s, (3) zinc precursor reaction and purge for 10 s, (4) either deionized water dose for 0.015 s (ZnO-W) or ozone dose for 0.2 s (ZnO-O), and (5) oxygen precursor reaction and purge for 10 s. This sequence constructs one ALD cycle, and 900 cycles were used to deposit ZnO films in this study. The growth rate is 1.4−1.5 Å per cycle under this deposition condition. Ozone was supplied by an LG-7 Corona Discharge ozone generator (Del Ozone, USA). The differences in these well-established growth processes are illustrated in Figure 1a,b, respectively.12,13 Transport Properties. ZnO samples were cut into 1 cm × 1 cm size. Electrical conductivity and the Seebeck coefficient were measured under a dynamic flow of Ar/H2 reducing ambient in the temperature range of 300−787 K by the linear four-probe and temperature differential methods, respectively, using a commercially available thermoelectric tester RZ2001i (Ozawa Science Co. Ltd., Japan). 2795


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Chemistry of Materials beam. Throughout the sputtering process, the selected ions ascribed to oxygen, zinc oxide, and silicon were sequentially collected using a MAXIM spectrometer equipped with a quadrupole analyzer.

In addition to thermoelectric properties, drastic changes were also observed in the Hall measurement results of both ZnO-O and ZnO-W films before and after thermal cycling, as shown in Table 1. It is worthwhile to note that the RT electrical

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RESULT AND DISCUSSION Thermoelectric and Transport Properties. The firstcycle temperature dependent thermoelectric properties of ZnOW and ZnO-O samples are shown in Figure 1c,d. Typically, the samples take one or two thermal cycles to show constant and repeatable results, but the irreversible changes that take place in the electrical conductivity and thermopower during the first thermal cycle carry significant information about the nature of the ZnO films. A thermal cycle here includes heating and subsequent cooling of the sample to do thermoelectric measurements. The thermoelectric measurements were performed in an Ar/H2 atmosphere from RT to 787 K. As shown in Figure 1c, the as-grown ZnO-O films were relatively resistive with an electrical conductivity of 21.6 S cm−1 at RT. During the first heating step, a typical semiconducting behavior was observed up to around 450 K, where conductivity increases continuously with temperature. In the temperature range between 450 and 600 K, the ZnO-O films showed a rapid increase in electrical conductivity and a simultaneous decrease in absolute value of Seebeck coefficient, which we attribute to the formation of oxygen vacancies (VO) during thermal cycling in the Ar/H2 reducing ambient atmosphere.14 By the end of the first thermal cycle, the conductivity of the ZnO-O films has increased by a factor of 12 (from 21.6 to 268 S cm−1). As we discuss in detail later, it turns out that both carrier concentration and carrier mobility increase at the end of the first thermal cycle, which explains the increase in conductivity of ZnO-O films. Consequently, the Seebeck coefficient of ZnOO films decreased at the end of the first thermal cycle (Figure 1d). In contrast, the ZnO-W films exhibited a more complex behavior. The conductivity of the as-grown ZnO-W films were 184.2 S cm−1 at RT, which is nearly 1 order of magnitude higher than the conductivity of the as-grown ZnO-O films. During the first thermal cycle, the ZnO-W films showed degenerate behavior, where the conductivity decreased with increasing temperature up to 450 K. The Seebeck coefficient does not follow the trend that would be expected if the conductivity change was simply due to changes in carrier concentration over this same temperature range. Instead, it shows a linear temperature dependence over the same temperature range, suggesting that the decrease in conductivity of the ZnO-W films is possibly due to microstructural rearrangement rather than changes in carrier concentration. A slight increase in conductivity was observed in the temperature range between 450 and 550 K, where the formation of VO takes place. From 550 K, the ZnO-W films show a rapid and significant decrease in conductivity with a simultaneous increase in the absolute value of Seebeck coefficient, which is the opposite of the trend found in ZnO-O films. At the end of the first thermal cycle, the conductivity of the ZnO-W films has decreased by a factor of 6 (from 184.2 to 30.2 S cm−1). As we will discuss in detail later, we believe that this significant drop in the conductivity of the ZnO-W films after the first thermal cycle can be attributed to the formation of acceptor defects such as zinc vacancies (VZn) or oxygen interstitials (Oi), which reduce free electron concentration. Since the measurement was performed in a reducing atmosphere, the latter possibility (Oi) can be excluded.

Table 1. Thermoelectric and Transport Properties of ALD ZnO Films at RTa sample ZnO-O asgrown ZnO-W asgrown ZnO-O annealed ZnO-W annealed

nb (× 1019/ cm3)

μb (cm2/V·s)

−158.8

1.75

7.7

184.2

−107.9

4.13

27.8

268.2

−110.6

5.76 × 10−4

4.87

34.4

30.2

−204.9

2.82 × 10−4

0.69

27.1

σb (S/ cm)

S (μV/ K)

21.6

PFc (W/m·K2)

a

Electrical conductivity (σ), Seebeck coefficient (S), power factor (PF), carrier density (n), and Hall mobility (μ). bValues from the Halleffect measurement. cPower factor was calculated from the conductivity and Seebeck coefficient at 705 K, from thermoelectric properties measurement.

properties of our as-grown films match well with previous studies on ALD ZnO films in the literature.11,15,16 However, the carrier concentration of our ZnO-O and ZnO-W films shows another interesting trend. For example, the as-grown ZnO-O films had a carrier concentration of 1.75 × 1019 cm−3, while the as-grown ZnO-W films had a carrier concentration of 4.73 × 1019 cm−3. Furthermore, the Hall mobility of the as-deposited ZnO-O films was 7.7 cm2 V−1 s−1, but the as-grown ZnO-W films had a carrier mobility of 27.8 cm2 V−1 s−1. These measurements help explain why the as-grown ZnO-W films exhibit higher conductivity at RT (182.4 S cm−1) than asgrown ZnO-O films (21.6 S cm−1). However, at the end of the first thermal cycle, the carrier concentration and mobility of the two ZnO films changed dramatically (Table 1). For example, the carrier concentration of the ZnO-O films increased to 4.87 × 1019 cm−3 after thermal stabilization. Simultaneously, the Hall mobility of the same ZnO-O films increased to 34.4 cm2 V−1 s−1. Hence, the origin of the dramatic increase of the RT electrical conductivity of ZnO-O films after the thermal cycle is now clear. The electrical conductivity of ZnO-O films increased 12.4 times (21.8 to 268 S cm−1), which can be attributed to the 2.8 times increase in carrier concentration and 4.4 times increase in Hall mobility. In contrast, the carrier concentration of the ZnO-W films decreased to 0.69 × 1019 cm−3, while the Hall mobility remained nearly constant at the end of the first thermal cycle. Therefore, the RT conductivity of the ZnO-W films actually decreased by 6.1 times at the end of the first thermal cycle. This conductivity drop is mostly attributed to the diminished carrier concentration. The main point is that the carrier mobilities of the annealed ZnO-O and ZnO-W films become comparable, which suggests that the large difference in RT conductivity between annealed ZnO-O and ZnO-W films is most likely attributed to the formation of different majority point defects. As mentioned earlier, once these films are thermally cycled once (twice at the most), steady state performance, presumably with equilibrium defect concentration, sets in. Both annealed ZnO-O and ZnO-W films show repeatable data as their first cooling cycle results, and their thermoelectric and transport properties remain constant irrespective of how many more 2796


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Figure 2. Temperature dependent thermoelectric properties of ZnO thin films for the first, second, and third thermal cycle. Relatively large irreversible changes with the opposite trend have been observed during the first thermal cycle, while the films show repeatable results after the second measurement cycle. (a) ZnO-O electrical conductivity, (b) ZnO-O Seebeck coefficient, (c) ZnO-O power factor, (d) ZnO-W electrical conductivity, and (e) ZnO-W Seebeck coefficient. The difference in the power factor between annealed ZnO-O and ZnO-W is shown in f, where it is clear that a dramatic improvement in the power factor of ozone-deposited films can be observed compared to the water based one. Heating and the subsequent cooling cycle are presented with filled and open symbols, respectively.

previous reports on ALD ZnO films with different oxidants.11,18 Cross-section SEM micrographs of as-grown and annealed ZnO-O films are shown in Figure 3c,d, respectively. It is noted that as-deposited ZnO-O films show columnar structures. After the first thermal cycle, the ZnO-O films develop larger grains, and the XRD peak intensity increases (Figure 3g), both of which may explain the increased Hall mobility of ZnO-O films after annealing. In comparison, the micrographs of as-grown and annealed ZnO-W films are shown in Figure 3e,f, respectively, and the XRD pattern is shown in Figure 3g. The typical polycrystalline structure is clearly observed, without any obvious increase in the grain size or XRD peak intensity after thermal cycling. Figure 3h−k show atomic force microscope (AFM) micrographs of all ZnO samples. The root-mean-square roughness value for as-deposited and annealed ZnO-O samples are 2.48 and 2.76 nm, respectively. The corresponding values for ZnO-W samples are 2.86 and 2.66 nm. Both ZnO-O and ZnO-W films show low surface roughness, which is expected from the self-limiting mechanism of the ALD process. The different crystal structure of ZnO-O and ZnO-W films can be considered as a factor affecting the Hall-mobility, depending on the electron transport direction. In the wurtzite structure, relative displacement of cations and anions from the ideal structure induces a spontaneous polarization parallel to the c-axis.19 As a result, the effective mass of the electron (m*) is highly dependent on the electron transport direction of the anisotropic ZnO crystal.20 Our results are actually consistent with the concept that c-axis oriented ZnO-O films have larger mobility (lower m*) perpendicular to the c-axis, compared with the polycrystalline ZnO-W films. It is noted that m* is

thermal cycles they further go through. To show this thermal stabilization effect, the corresponding temperature dependent thermoelectric properties for three continuous thermal cycles are plotted in Figure 2. For annealed ZnO-O films, temperature dependent conductivity shows degenerate behavior in which the conductivity decreases with temperature, and the corresponding Seebeck coefficient is quite linear over the temperature range of 350−700 K. In comparison, annealed ZnO-W films show mild nondegenerate behavior over the same temperature range. The power factor of the annealed films is shown in Figure 2f. The highest power factor of annealed ZnO-W films is only 2.89 × 10−4 W m−1 K−2 at 746 K. On the other hand, annealed ZnO-O films show a remarkably reversible power factor that reaches 5.76 × 10−4 W m−1 K−2 at 705 K (nearly double the water based ZnO films). Furthermore, this power factor is much higher than the best previously reported power factor of undoped ZnO film (2.5 × 10−4 W m−1 K−2 at 800 K, prepared by sol−gel method).17 Given the drastic differences in the electrical properties and the thermoelectric power factor based on the oxidant used, it is important to understand the mechanism behind this difference. Therefore, a series of materials analysis studies were carried out in order to explain this interesting behavior. Microstructure and Surface Morphology. Figure 3a,b show scanning electron microscope (SEM) micrographs of the as-grown ZnO-O and ZnO-W, respectively. The top view images show that the ZnO-O films have circular grains, while the ZnO-W films have elongated grains pointing to a different growth orientation. These microstructural differences match 2797


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Figure 3. SEM images of (a) as-grown ZnO-O and (b) as-grown ZnO-W. Cross-section SEM images of (c) as-grown ZnO-O, (d) annealed ZnO-O, (e) as-grown ZnO-W, and (f) annealed ZnO-W. (g) X-ray diffraction patterns of ALD ZnO thin films. AFM images of (h) as-grown ZnO-O, (i) annealed ZnO-O, (j) as-grown ZnO-W, and (k) annealed ZnO-W. Scale bars: (a,b) 500 nm, (c,d,e,f) 200 nm, and (h,i,j,k) 400 nm.

Figure 4. (a,b) XPS spectra of the O 1s core level of ZnO-O and ZnO-W, respectively. (c) Area fraction of oxygen components based on the fitting results. (d) Estimated relative atomic fraction of ZnO films based on XPS results. The arrow shows extraordinary change.

perpendicular to the amorphous glass substrate. Since the c-axis direction is the lower mobility (larger m*) direction in ZnO, the enhanced c-axis orientation after annealing may have helped

reciprocally proportional to the Hall mobility. After annealing, the ZnO-O film becomes more c-axis oriented, which means a larger number of crystallites now have the c-axis direction 2798


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Figure 5. (a,b) D-SIMS depth profile of ZnO-O and ZnO-W, respectively. Dashed line indicates the interface between ZnO film and glass substrate. (c,d) Secondary ion signal comparison of O − and ZnO −, respectively.

to increase the in-plane Hall mobility of ZnO-O.21 In contrast, neither the microstructure nor the XRD patterns of ZnO-W films changed much after annealing. Thus, the drastic difference in power factor between the ZnO-O and ZnO-W films cannot be completely attributed to microstructural changes. Likewise, the issue of polarization and higher effective mass along the c-axis of ZnO cannot induce such a big change in conductivity, since the mobility of the stabilized ZnO-W and ZnO-O are comparable.22 However, a large drop in the carrier concentration is observed in the ZnOW film after the first thermal cycle, which is the opposite trend found in ZnO-O. Thus, we conclude that concentration and type of point defects that dominate each one of these films must be different, leading to the differences in carrier concentration, conductivity, Seebeck coefficient, and hence the power factor. It should be noted that there is slight variation between the measured electrical conductivity values from two different measurement systems, but only for the ZnO-W samples. The variation is possibly due to the inhomogeneously distributed point defects in ZnO-W films, since water is a less efficient oxidant compared to ozone. The question we must now answer is, which point defects are responsible for the observed changes in carrier concentration shown in Table 1. In the case of ZnO-O films, where an increase in carrier concentration was observed after the first thermal cycle, electron donating defects (such as oxygen vacancies) must have been formed. In the case of ZnO-W films, the drop in carrier concentration after thermal cycling can be understood (as we will show) by the competitive formation of VZn (acceptor) and VO (donor) defects. We have used multiple techniques to confirm this hypothesis.

X-ray Photoelectron Spectroscopy (XPS). XPS survey spectra from all ZnO films are presented in Figure S1, showing that Zn, O, and C element are detected. Figure 4a,b show the XPS spectra for the O 1s core level of ZnO-O and ZnO-W films, respectively. The O 1s core level was fitted into three components, namely OI, OII, and OIII, centered at 529.9, 530.9, and 532.0 eV, respectively. 23 The dominant O I peak corresponds to oxygen in the ZnO lattice without oxygen vacancy, whereas the OII peak denotes oxygen in the ZnO lattice with oxygen vacancy. Additionally, the OIII peak could be assigned to the surface attached oxygen such as hydroxyl (−OH) and carbonyl (−COOH) groups.23 It is interesting to note that both as-grown ZnO-O and ZnO-W films have a similar amount of VO despite their very different carrier concentrations. The area fractions of each oxygen component were estimated based on the fitting results and are shown in Figure 4c. In the case of as-grown films, [OII/(OI + OII + OIII)] is 17.8% for ZnO-O and 16.2% for ZnO-W. Considering that as-grown ZnO-W shows a 2.3 times higher carrier concentration even with a lesser amount of VO, these excessive carriers in ZnO-W most likely originate from the zinc interstitials (Zni) rather than VO. For annealed films, [OII/(OI + OII + OIII)] is increased to 24.0% for ZnO-O and 24.6% for ZnO-W. This indicates the formation of VO from both films, in agreement with the 4.9 times increase of the carrier concentration of ZnO-O after annealing. However, the carrier concentration of ZnO-W decreased by 5.9 times even after the formation of VO. Considering that both annealed films have a similar level of VO, simultaneously formed VZn seems to be the most possible reason for the decreased carrier density in ZnO-W films after thermal cycling. Figure 4d shows the estimated atomic fraction 2799


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Figure 6. (a) Low temperature photoluminescence and (b) room temperature UV−vis spectra of all ZnO films. Corresponding Tauc plots for optical band gap analysis are shown as an inset in b.

of ALD ZnO films based on the peak area of O 1s, Zn 2p, and C 1s. Relative content of oxygen in ZnO-O films decreased after annealing, which is most likely due to the formation of VO. However, it increased after annealing in the case of ZnO-W (marked with arrow). This does not necessarily mean an increase of oxygen content, because the formation of VO has already been confirmed by XPS O 1s spectra. Therefore, the most logical explanation is that the ZnO-W film rapidly loses the zinc atoms during annealing (i.e., formation of VZn) along with the simultaneous formation of VO, where the former reaction is dominant. Dynamic Secondary Ion Mass Spectrometry (D-SIMS). D-SIMS depth profiles of ZnO-O and ZnO-W samples are shown in Figure 5a,b, respectively. All ZnO films are homogeneous through the whole eroded depth since O− and ZnO− detected ions exhibit stable signals until reaching the glass interface described by the sharp increase of the silicon signal. For the easier comparison, O− and ZnO− signals ascribed to all ALD ZnO films are plotted together in Figure 5c,d, respectively. For ZnO-O, it is clearly seen that only the O− signal decreases after annealing, while the ZnO− signal remains at the same level. The decrease of O− signal can be explained by the spontaneous formation of VO during annealing. The constant ZnO− signal, despite the fact that it is attributed to both oxygen and zinc species, seems to be more sensitive to the zinc content than that of oxygen. In contrast, ZnO-W exhibits a decrease of both O− and ZnO− signals after annealing, which is in agreement with the previously mentioned competitive formation of VZn and VO during annealing. Since the O− signal of the as-grown ZnO-W is higher than that measured for the ZnO-O films, it is reasonable to apply the same assumption suggesting that the ZnO− signal is more affected by the zinc content. Considering the higher carrier density of the as-grown ZnO-W compared to the as-grown ZnO-O film, the enhancement of both O− and ZnO− signals of the as-grown ZnO-W is another evidence suggesting that Zni is the predominant origin of the excessive carriers in as-grown ZnO-W films, which has been experimentally confirmed to be a dominant native shallow donor in ZnO by Look et al.24

Optical Characterization. Figure 6a shows the photoluminescence spectra of all ZnO films in the wavelength range from 350 to 700 nm, measured at 77.25 K. Three distinct peaks located at approximately 540 nm (green), 410 nm (violet), and 370 nm (UV) are observed. Peak intensity was normalized with respect to the strong UV peak which originated from the nearband-edge (NBE) emission. The violet emission is attributed to the shallow donor level to the valence band transition and the conduction band to the shallow acceptor level transition. A slight increase in the green emission was observed in the annealed ZnO-O films. Interestingly, a significant increase in the same emission was noticed in annealed ZnO-W film. The origin of green emission in ZnO is still a controversial issue but has been attributed to either oxygen vacancy or donor-toacceptor transition. According to a theoretical calculation reported by Janotti et al., VO does not show any possible electron transition for the green luminescence.25 Electron transition from VZn to the valence band is also a possible origin of the green emission; however, the corresponding luminescence center at 2.5 eV (∼500 nm) is not well matched to our study. The observed green emission near 540 nm is likely to be related to more than one defect level. Zhao et al. reported that green emission centered at ca. 540 nm diminished after implanting of either zinc or oxygen into the bulk ZnO sample, which supports donor to acceptor transition as the origin of the observed green emission.26 It is important to note that the neutral state of VO is located at ∼0.86 eV below the conduction band and neutral VZn is placed at ∼0.215 eV above the valence band.27−30 For the typical band gap energy of ZnO, 3.37 eV, the corresponding electron transition energy from VO to VZn is 2.295 eV (wavelength of ca. 540 nm), which agrees well with the observed green emission near 540 nm. We believe that the enhanced green emission is another evidence for the competitive formation of both VO and VZn in ZnO-W during annealing. UV−vis transmittance spectra are shown in Figure 6b, and the corresponding Tauc plots are shown as an inset. All transmittance spectra of ZnO films rapidly increased at around 370 nm of incident wavelength, which is consistent with the 2800


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Figure 7. (a) Schematic illustration of point defect model for ALD ZnO films. (b) Comparison of thermoelectric power factor of recent literature on ZnO thin films by multiple preparation methods.

for enhancing the thermoelectric performance of other materials.

NBE in the PL spectra. The as-grown ZnO-W shows higher average transparency in the visible range, and no obvious change was detected after annealing. In the case of ZnO-O, the transmittance slightly decreased after annealing. Optical band gaps (Eg,opt) had been estimated based on the Tauc plot. The Eg,opt of ZnO-O increased from 3.26 to 3.28 eV after annealing, while the Eg,opt of ZnO-W decreased from 3.29 to 3.28 eV. It is also noticed that a blue (red) shift occurred in the ZnO-O (ZnO-W) films, indicating an increase (decrease) of the carrier concentration due to the Burstein−Moss effect,31 which is in good agreement with the electrical measurements that we discussed in a previous section. To recap, the oxidant dependent thermoelectric and transport properties of ALD ZnO films can be explained with the help of the point defect model shown in Figure 7a. For the as-grown films, the excessive carrier density found in ZnO-W is mostly likely attributed to the additional Zni compared to ZnOO. During the thermal cycling in reducing ambient atmosphere, VO is created in both ZnO-O and ZnO-W films, which is confirmed by XPS study. However, the decrease in carrier concentration of ZnO-W suggests that another acceptor defect must be simultaneously created in ZnO-W. Considering that the films were annealed in an oxygen deficient ambient atmosphere, formation of VZn is most likely to be the origin of acceptor defects rather than Oi. The competitive formation of VO and VZn in ZnO-W is the possible cause of reduced n-type carriers, even after annealing in a reducing ambient atmosphere, through the self-compensation mechanism found in wide band gap semiconductor materials.32 This drastic differences in defect chemistry between ZnO-O and ZnO-W can explain the oxidant effect on the thermoelectric properties of ALD films. Figure 7b shows a literature review of the recent thermoelectric studies on undoped and doped ZnO thin films by multiple deposition methods.14,17,22,33−38 One may notice that there are a limited number of studies on the thermoelectric properties of undoped ZnO films compared to doped ZnO films. A strong oxidant effect is clearly seen for undoped ALD ZnO films. The power factor of undoped ALD ZnO films grown using ozone as an oxidant is comparable to some of the substitutionally doped ZnO films. This result shows the importance of ALD oxidant, which could potentially be used

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CONCLUSION In summary, we have demonstrated a large thermoelectric power factor in undoped ZnO films of 5.76 × 10−4 W m−1 K−2 at 705 K by controlling the oxidant species during atomic layer deposition. Ozone-based ZnO films showed nearly double the power factor of water-based ZnO films, a fact we show is mostly related to the type and concentration of point defects present in the films. This result is likely due to the insufficient oxidation power of the deionized water (compared to ozone), which results in a higher content of native Zni that leads to selfcompensation of carriers at elevated temperature by competitive formation of VO and VZn. This defect-induced drastic difference in carrier concentration is the root cause of the large difference in thermoelectric power factor of these oxides. Further improvements in thermoelectric performance of ALD ZnO are expected to be achieved by additional substitutional doping or structure engineering.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.6b04654. Further information on XPS spectra for survey and Zn 2p core levels, sample preparation for Ohmic contact, and electrical and thermoelectric property measurement (PDF)

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Hyunho Kim: 0000-0003-2381-9716 Husam N. Alshareef: 0000-0001-5029-2142 Author Contributions

All authors have given approval to the final version of the manuscript. 2801


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Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS Research reported in this publication was supported by King Abdullah University of Science and Technology (KAUST). REFERENCES

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