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Methodology of Thermoelectric Power Factor Enhancement by Controlling Nanowire Interface Takafumi Ishibe, Atsuki Tomeda, Kentaro Watanabe, Yoshinari Kamakura, Nobuya Mori, Nobuyasu Naruse, Yutaka Mera, Yuichiro Yamashita, and Yoshiaki Nakamura ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b13528 • Publication Date (Web): 10 Oct 2018 Downloaded from http://pubs.acs.org on October 13, 2018

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ACS Applied Materials & Interfaces

Methodology of Thermoelectric Power Factor Enhancement by Controlling Nanowire Interface

Takafumi Ishibe a), Atsuki Tomeda a), Kentaro Watanabe a), Yoshinari Kamakura b), Nobuya Mori b), Nobuyasu Naruse c), Yutaka Mera c), Yuichiro Yamashita d), and Yoshiaki Nakamura a) * a) Graduate

School of Engineering Science, Osaka University, 1-3 Machikaneyama-cho, Toyonaka, Osaka 560-8531, Japan

b)

Graduate School of Engineering, Osaka University, 2-1 Yamada-oka, Suita, Osaka 565-0871, Japan c) Faculty

d) National

of Medicine, Shiga University of Medical Science, Otsu, Shiga 520-2192, Japan

Institute of Advanced Industrial Science and Technology, 1-1-1 Umezono, Tsukuba, Ibaraki 305-8563, Japan E-mail: [email protected] Telephone numbers: +81-6-6850-6315

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ABSTRACT The simultaneous realization of low thermal conductivity and high thermoelectric power factor in materials has long been the goal for the social use of high-performance thermoelectric modules. Nanostructuring approaches have drawn considerable attention because of the success in reducing thermal conductivity. On the contrary, the enhancement of thermoelectric power factor, namely the simultaneous increase of Seebeck coefficient and electrical conductivity, has been difficult. We propose a method for the power factor enhancement by introducing coherent homoepitaxial interfaces with controlled dopant concentration, which enables the quasi-ballistic transmission of high-energy carriers. The wavenumber of the high-energy carriers is nearly conserved through the interfaces resulting in simultaneous realization of high Seebeck coefficient and relatively high electrical mobility. Here, we experimentally demonstrate the dopant-controlled epitaxial interface effect for the thermoelectric power factor enhancement using our “embedded-ZnO nanowire structure” having high-quality nanowire interfaces. This presents the methodology for substantial power factor enhancement by interface carrier scattering.

KEYWORDS Nanowire, ZnO, thermoelectric material, thermoelectric power factor, carrier transport, phonon


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INTRODUCTION Thermoelectric materials can generate green energy from wasted heat originating from automobile, factory, etc.1,2 Thermoelectric material performance is quantified by a dimensionless figure of merit: ZT = S2σT/κ, where T is absolute temperature, σ is electrical conductivity, κ is thermal conductivity, and S is Seebeck coefficient. The interdependence among the three material properties, namely σ, S, and κ has made it difficult to enhance ZT1,2. Therefore, the simultaneous realization of low κ and high thermoelectric power factor (S2σ) in materials has long been the goal for the social use of high-performance thermoelectric modules. Recently, nanostructuring approaches have drawn much attention because of the success in reducing κ while maintaining σ3-9. However, the improvement in ZT is limited because the reduction in κ is limited to the amorphous value. Therefore, it is demanded actively to increase S2σ, not only to maintain its bulk value, along with the reduction in κ. Although increasing S2σ has been difficult in most cases, many efforts have been reported: increasing density of states (DOS) such as band convergence10, resonant scattering11,12, and lowdimensional effect13, as well as controlling the scattering processes such as energy filtering effect14-17. As a promising universal method without depending on the materials, the energy filtering effect has been studied for the S increase in theoretical approaches 14,15,17, where higher (lower) energy carriers than the energy barrier (EB) are transmitted (reflected) through (at) the introduced interfaces. However, in general, the S reduction often occurred in realistic nanostructured materials owing to defect scattering mechanism20 because the uncontrolled interfaces include various defects such as point defects (dangling bonds, etc.) and misfit dislocations that scatter high-energy carriers. Therefore, the energy filtering effect is not promising for the realistic enhancement of S2. For the building of universal S2enhancement technique controlling the scattering processes, we propose a novel enhancement method of S2 by introducing the “dopant-controlled epitaxial (DCE) interface” that connects two crystal structures coherently with almost no lattice distortions or misfit dislocations, where the dopant control at the interfaces (Fermi energy (EF) control) creates the energetic barrier for carriers. Unlike the aforementioned uncontrolled interface with high-energy carrier scattering, ACS Paragon Plus Environment

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the DCE interface with a controlled barrier enables the quasi-ballistic transmission of high-energy carriers through the interfaces with almost no scattering, where the wave function of the high-energy carriers is coherently connected (Figure 1(a)), resulting in the simultaneous realization of high Seebeck coefficient and relatively high electrical mobility. Ubiquitous material, ZnO, expected for use in transparent thermoelectric material applications, is suitable for investigating the scientific effect based on the large number of studies and information. Here, we experimentally demonstrate the aforementioned DCE interface effect that enhances S2 in our “embedded-ZnO nanowire (NW) structure (ENS),”21 as shown in Figures 1(a) and (b) in addition to the κ reduction via the NW interface phonon scattering. This study presents the novel guideline for substantially enhancing S2 in interface carrier scattering methods.

RESULTS AND DISCUSSION Optical transmittance and Seebeck coefficient characteristics Figure 1(c) shows the typical plan-view scanning electron microscope (SEM) image of the ENS surface. The observed bright hexagonal structures correspond to the top of the ZnO NWs as shown by the schematic of the cross-sectional ENS in the inset of Figure 1(c). Figure 1(d) displays the transparent ENS formed on quartz glass with the transmittance characteristic shown in Figure 1(e). The written characters (“Nakamura Lab.”) are clearly shown in Figure 1(d) because of the high optical transmittance of the ENS under the visible range (Figure 1(e)). This demonstrates the application potential of the ENS as a transparent-device material. To compare with the ENS, ZnO films without NWs were formed on the ZnO buffer layers (Film A) and directly on the substrates (Film B). Figure 2(a) shows the S of the ENS (inverted triangles), Film A (circles), and B (squares) as a function of  with the results of ZnO materials reported by other groups22-28. A remarkable fact is that the ENS showed a larger S than the films without NWs (Film A and B) at the same . However, it is difficult to conclude that this S enhancement is due to the NWs because S also depends on film properties, namely, the crystallinity and ACS Paragon Plus Environment

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the electron concentration (n). To remove the effect of the n, the S values of the ENS (inverted triangles), Film A (circles), and B (squares) are shown as a function of n with the following calculated S curves in Figure 2(b). The carrier transport properties can be derived from the Boltzmann transport equation under the relaxation time approximation and the parabolic band approximation. S can be expressed as

S

1 qT





0

 f ( E )  dE E     f ( E )  0  ( E ) g ( E ) E  E dE

 ( E ) g ( E ) E ( E  E F ) 

(1),

where q is the elementary charge, (E) is the energy-dependent carrier relaxation time, g(E) is the DOS, and f(E) is the Fermi–Dirac distribution function. We simply define (E) as the expression:(E) = AEr; A is an energy-independent coefficient and r is a scattering parameter. Based on Eq. (1), the S curves were calculated as shown by dotted and solid lines in Figure 2(b), where scattering parameters were used as 1.5 (ionized impurity scattering)29 and -0.5 (short-range defect scattering)20, respectively. It is difficult to discuss scattering mechanisms completely due to large experimental data variability in n - S graph. However, the ENS clearly shows a larger S than the films without NWs, even when the influence of n on S is considered. Furthermore, to discuss the effect of film crystallinity on the thermoelectric properties, we evaluated the film crystallinity by the full width at half maximum (FWHM) in an X-ray rocking curve around the 0002ZnO diffraction peak. We investigated the relationship between S and crystallinity (FWHM in rocking curve), and found a larger S in the ENS compared to Film A and B at the same n of ~4×1018 cm-3 and the same FWHM. This proved that S enhancement comes from the NW structure, neither from the dependence on n nor from the film crystallinity (see Figure S1 in the Supporting Information).

Structural characteristics of ENS ACS Paragon Plus Environment

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We performed the cross-sectional transmission electron microscope (TEM) observation of the ENS as shown in Figure 3(a). Figure 3(b) is an enlarged image of the broken square region in Figure 3(a). The solid square region in Figure 3(b) near the interface of the embedded NW parts is magnified in Figure 3(c). The TEM image in Figure 3(c) revealed three crystal parts: (1) core and (2) shell parts of the NW, and (3) filling parts among the NWs as shown in the ENS schematic in Figure 3(g). The fast Fourier transformation (FFT) patterns of the shell parts (Figure 3(e)) are the same as those of the core parts (Figure 3(d)), revealing the epitaxial relationship between the core and shell parts due to the lateral growth of the shell parts from the core ones (see Figure S2 in the Supporting Information). This indicates the coherent core/shell homointerfaces. On the other hand, FFT pattern of the filling parts (Figure 3(f)) was different from ones of NW parts (=core and shell parts) in Figures 3(d) and (e), indicating that the crystal orientations in the filling parts were different from those of the NW parts. The filling parts among the NWs were considered to have formed on the underlying buffer layers at the perpendicular direction to the substrate surfaces. Furthermore, at the epitaxial core/shell homointerface, the Zn concentration was lower than that of the other parts, as shown by the solid arrow in the electron dispersive X-ray (EDX) spectroscopy profile (Figure 3(h)), while no Zn concentration reduction was observed at the other undefined interfaces such as the interfaces between filling parts. The Zn interstitials are known to act as shallow donors30, resulting in the n-type semiconductor behavior in the present undoped ENS. Therefore, this low Zn concentration at the core/shell epitaxial interfaces indicates the concentration modulation of the Zn interstitial dopants, which can create an energetic barrier at the epitaxial interfaces. Namely, the core/shell interfaces are the aforementioned DCE interfaces. We can control a proportion of these core/shell epitaxial interfaces in the ENS by changing the NW areal density. We formed the ENS samples with different NW areal density: low density (LD) of ~ 1.0×109 cm-2 and high density (HD) of > ~ 4.0×109 cm-2. In the ENS with HD NWs (HDENS), the core/shell NWs are directly connected to the neighboring NWs easily because the diameter of the core/shell NWs (~ 120 nm) shown in Figure 3(b) is almost equivalent to the distance between the NW centers (~ 140 nm), which reduces the proportion of the filling parts. Therefore, in the HDENS, the ACS Paragon Plus Environment

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proportion of the DCE interfaces (core/shell epitaxial interfaces) is larger than that of the ENS with LD NWs (LDENS) with many undefined polycrystalline interfaces in the filling parts.

Quasi-ballistic transmission of high-energy carriers through the DCE interfaces For a detailed discussion about the carrier transport in the ENS with DCE interfaces, the electron Hall mobilities () were shown as a function of n in Figure 4(a) with the semi-empirical curve31 and the results of ZnO materials reported by other groups22,25-28,32,33. At an n of ~1018-19 cm-3, the HD and LDENSs showed different dependencies (dotted and dashed lines, respectively) in Figure 4(a), which arises from the differences in the NW areal density in the ENS samples (see Figure S2 in the Supporting Information and methods section). As n increases, decreases in the LDENS (triangles), while  increases peculiarly in the HDENS with the DCE interfaces (diamonds). This increasing trend is similar to that of grain boundary (GB) scattering. According to a previous work34, the GB is negatively charged by the interface electron trap, resulting in EB for strong GB scattering34. At a carrier concentration of n0 when the trapping sites are completely filled, EB becomes the largest and  becomes the lowest. When n is becoming larger than the n0 value, EB decreases because of the screening effect by the positively charged ionized donors, resulting in the increase of . The n0 value is known to exceed 1018 cm-3 from previous works35. This trend with n0 was confirmed in Films A and B, which detail is described in Figure S3 in the Supporting Information. However, in the present HDENS with the DCE interfaces, the order magnitude of n0 with the smallest  is assumed to be ~ 1017 cm-3 as shown in Figure 4(a). In addition, the aforementioned GB scattering was reported only at the polycrystalline interfaces. Therefore, this GB scattering can be ruled out as the mechanism of our trend in the present ENS with the DCE interfaces (less polycrystalline interfaces). We assumed a N+ (high-doped crystal) -N- (lowdoped interface)- N+ junction due to the modulated dopant concentration near the DCE interfaces shown in Figure 1(a), where EB decrease with increasing n can also be theoretically explained due to the ACS Paragon Plus Environment

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screening effect (see Section S4 in the Supporting Information). This EB unique to the DCE interfaces can bring the large aforementioned S enhancement effect shown in Figure 2(b) while exhibiting almost the same  as the bulk one (n of ~ a few 1019 cm-3 in Figure 4(a)) through the mechanism of the S2 enhancement proposed in Figure 1(a). Here, the eigen wavefunctions of the energetic carriers almostcoherently spread over the two epitaxially connected crystals with DCE interface (core and shell in NWs). Namely, the key point of our proposed enhancement mechanism (DCE interface effect) is the quasi-ballistic transmission of high-energy carriers through homoepitaxial interfaces with almost no GB scattering, in addition to the prevention of low-energy carrier conduction, bringing substantial power factor enhancement in interface carrier scattering methods. First, let’s prove the quasi-ballistic transmission of high-energy carriers through homoepitaxial interfaces in the aforementioned DCE interface mechanism by investigating the carrier transport properties in an HDENS with DCE interfaces. The  values of the HDENS with DCE interfaces were measured in the low temperature range without the phonon influences

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as shown in Figure 4(b). It is

known that the temperature dependence of  with carriers going beyond energy barrier EB can be described by Seto’s formula34 as follows:  EB    kT 

   0 exp 

(2),

where 0 is the intragrain mobility and k is the Boltzmann constant. In the low temperature range, the possible scattering mechanisms are generally considered to be ionized impurity scattering, short-range defect scattering, and GB (interface) scattering, which have the relaxation times of ion, SD, and grain, respectively. In ionized impurity scattering mechanism, ion can be expressed by the Brooks–Herring formula29:


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ACS Applied Materials & Interfaces 2  8md LD E   16 2 2md 32   8md LD 2 E  2      ion ( E )  E log 1   2 2  2 Z0 q 4 Ni 8 m L E    1 d D    2

1

(3),

where  is the static permittivity of the material, md is the DOS effective mass, Ni is the ionized impurity concentration, Z0 is the charge in the impurity in units of q,  is the reduced Planck constant, and LD is the Debye length. md was obtained to be 0.3m0 (m0 is the free electron mass) from the results of optical transmittance (see Section S5 in the Supporting Information). In short-range defect scattering, SD can be expressed as20

 SD ( E ) 

 4 2 ( m d ) 3 / 2 N nU n

2

E



1 2

(4),

where Nn is the nonionized defect concentration and Un is the short-range potential of the defects. In this study, the Un of Bi2Te3-xSex (1×10-46 J m3) was used as the example value20. In GB scattering, grain can be expressed as

 grain ( E )  L

md  12 E 2

(5),

where L is the effective grain size. The 0 in Eq. (2) were calculated by Boltzmann transport equation, where the relaxation time was expressed as 1/ion(E) +1/SD(E) +1/grain(E) through Matthiessen’s rule (see Section S6 in the Supporting Information). We fitted the experimental  data in the HDENS using the calculated  with three parameters, Nn, L, and EB. For simplicity, the experimental  data between 100 and 200 K were used owing to the almost constant values of n (defined as Ni) in this temperature range (see Section S7 in the Supporting Information). In this fitting analysis, we assumed that Z0 was unity. The data was best-fitted when Nn was 6.4 ± 0.9×1019 cm-3, L was 181 ± 40 nm, and EB was 21 ± 0.7 meV; the fitting curve is shown in the inset of Figure 4(b). Then, the relaxation rate in GB (interface) scattering was almost ignorable compared with the other two scattering (see Section S6 in the ACS Paragon Plus Environment

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Supporting Information). This proved the quasi-ballistic transmission of carriers with higher energy than EB through the DCE interfaces without GB (interface) scattering in HDENS.

Seebeck coefficient enhancement by the DCE interface To investigate the EB–induced S enhancement in the DCE interface mechanism where only high-energy carriers transmitted through the interfaces with EB, the S values of HDENSs with n of ~ 7×1018 cm-3 at room temperature (RT) were measured in a low temperature range of 170–320 K (Figure 5(a)). Film A showed the usual trend: increase in S with increasing T. However, the ENS showed an anomalous trend: with increasing T, the S value decreased above 260 K while S increased in the temperature region of 170-260 K. This anomalous trend (S decrease with T increase) can be explained by EB reduction due to some type of screening effect in DCE interface mechanism (like the n -  property). We simulated EB in this N+-N--N+ regions (see Section S4 in the Supporting Information) and measured n-T dependence (see Section S7 in the Supporting Information), proving the EB decrease with T increase above 260 K, that is such a screening effect by n increase that accompanies the T increase (Figures 5(c) and (d)). Here, we check whether the DCE interface mechanism is reasonable by discussing the carrier transport properties quantitatively. In the case that the energy relaxation length is comparable to interface spacing, S is rewritten by using the multiplication of the intragrain relaxation time 0(E) and the transmission coefficient through the interface T(E) 16,

S

1 qT





0

 f ( E )  dE E     f ( E )   0 ( E )T ( E ) g ( E ) E   dE 0 E  

 0 ( E )T ( E ) g ( E ) E ( E  EF ) 



(6).

Although T(E) practically becomes a complex function, the simplest approximation is performed for comprehension; T(E) is assumed to be a step function as follows: ACS Paragon Plus Environment

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1 ( E  E B ) T (E)   0 ( E  E B )

(7).

The temperature dependence of S in Figure 5(b) was fitted using the theoretical S calculated from Eqs. (6) and (7), with a fitting parameter EB, where the Nn and L values obtained in the -T data were used (6.4×1019 cm-3 and 181 nm, respectively). The fitting curve was shown by the solid line in Figure 5(b) and EB was estimated to be 22 ± 1.1 meV. This value is almost equivalent to the EB value obtained in the –T data (21 ± 0.7 meV). This consistency of EBs acquired from different measurements ( and S) is a direct evidence of the DCE interface mechanism proposed in Figure 1(a).

Thermoelectric power factor and thermal conductivity in the ENS The HDENS with DCE interfaces showed a higher S2 than the films at the same  as shown in Figure 6. To optimize the doping condition, we prepared the ENS by filling the Al-doped and Ga-doped ZnO. The Al-doped ENS showed the highest S2 value of ~ 7.0 Wcm-1K-2 in the ZnO films, which is about twice larger than the highest S2values ever reported (3.9 Wcm-1K-2)28 in ZnO-based films. This anomalous S2enhancement proved the DCE interface effect with the quasi-ballistic transmission of carriers with higher energy than EB, bringing substantial power factor enhancement in interface carrier scattering methods. This effect is easily expected to have material universality. For confirmation, this was proved by the simulation experiment (see Section S8 in the Supporting Information). The further enhancement is expected at the heavy doping concentration equivalent to the one of bulk ZnO showing the maximum value of S2 For  of the ENS, we embedded NWs completely to obtain flat surface of ENS (surface roughness of ±25 nm). The  of the ENS with NW areal density of ~8.0×108 cm-2 was measured to be 17.5±0.7 Wm-1K-1,37,38 where the error includes the thickness variability coming from the surface roughness of ENS. This value was lower than the measured  of Film A (21.3 Wm-1K-1). In the future ACS Paragon Plus Environment

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work, further  reduction is expected by optimizing the parameters (NW areal density and doping concentration). Here, for reference, we estimated the ZTs of HDENS and Film A at RT to be 0.012 and 0.003, respectively, although the measurement direction of electrical properties is different from that of

. Therein, the maximum S2 values of 7.0 Wcm-1K-2 (2.3 Wcm-1K-2) and measured  of 17.5 Wm1K-1

(21.3 Wm-1K-1) were used in HDENS (Film A). This demonstrates that the ENS is a promising

option for the realization of a vital goal: S2 enhancement and  reduction, based on our strategy in Figure 1(a).

CONCLUSIONS We demonstrated that the DCE interfaces in ENS enabled quasi-ballistic transmission of carriers with higher energy than EB, causing anomalous S2enhancement. A lower concentration of Zn interstitial dopant at the homoepitaxial core/shell interfaces was observed, which indicates the DCE interface formation. T dependence of  proved the quasi-ballistic transmission of high-energy carriers through homoepitaxial interfaces. An EB of ~20 meV was separately acquired by fitting the data of S and

 as a function of T, which showed the direct evidence of the DCE interface mechanism. As a result, the doped ENS showed the highest S2 value of ~7.0 Wcm-1K-2 in the ZnO films. This presents the methodology for substantial power factor enhancement by interface carrier scattering.

EXPERIMENTAL METHODS Sample preparation Substrates of size 2 mm × 10 mm × 0.3 mm were cut from various substrates (undoped Si(111) wafers of resistivity 1000–2000 Ω∙cm, thermally oxidized Si(001), and quartz glass) and introduced into a high vacuum chamber at a base pressure of approximately 1×10-6 Pa. The detailed formation procedure ACS Paragon Plus Environment

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has been previously reported21. The ZnO buffer layers were formed on the substrates at 673 K under an oxygen pressure of 1 Pa for 1 h using PLD. The ZnO targets were developed by pressing the ZnO powders (99.999 % purity) at 8 MPa and sintering at 1373 K in air for 24 h. The targets were placed 3 cm away from the substrates and irradiated with 10 Hz pulsed laser (ArF excimer laser with 193 nm wavelength). After the formation of the buffer layers, ZnO NWs were grown on ZnO buffer layers/substrates using the physical vapor transport method, where the wetting layers may be formed at the first growth stage. The Zn powder source (99.999 % purity) and ZnO buffer layers/substrates were placed in a quartz tube of 50-mm diameter that was placed in the furnace. The Zn powder and ZnO buffer layers/substrates were positioned at the center of the furnace and at approximately 15 cm downstream of the furnace center, respectively. Ar gas flowing at 400 sccm was used as the carrier gas, whereas O2 gas flowing at 1.5 sccm was used as the reactive gas. The temperature at the position of the source material was increased to 1053 K in 25 min and cooled to RT in 40 min. As a result, ZnO NWs grew on the ZnO buffer layers/substrates. A density of ZnO nanowires was changed by controlling O2 flow rate in PVT method.39 Subsequently, to increase the wetting layer resistivity to remove the influence on the thermoelectric performance, the samples were annealed at 873 K under O2 flow for 10 min. Zn interstitials, acting as donor, were also eliminated at the nanowire interface by the annealing under O2 flow. Thus, the doping level was modulated near the nanowire interface.40 Next, we embedded the NWs with ZnO by depositing ZnO at 673 K under an oxygen pressure of 0.005–0.2 Pa for 2.5 h in a vertical direction on the substrate to prevent the void formation unlike the ENS with voids formed by oblique deposition 21. The embedded thicknesses were ~150 nm (for electrical properties) and ~500 nm (for thermal conductivity). For a high carrier concentration of ~1019-20 cm-3, Al or Ga doping (0.1-0.5 at%) was performed by the deposition of Al- or Ga-doped ZnO. For reference, we further fabricated ZnO films on the ZnO buffer layers/substrates (Film A) and ZnO films on the substrates (Film B) under the same conditions of PLD for embedding the NWs. Film A and B exhibited higher and lower crystallinity (see Section S1 in the Supporting Information), respectively. ACS Paragon Plus Environment

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Structure characterization The surface and cross-sectional morphologies of the samples were observed by scanning electron microscopy (SEM) with a 10-keV electron beam. The crystal orientation was measured by Xray diffraction (XRD) with a Cu Kα line (wavelength: 0.15418 nm). Transmission electron microscopy (TEM) observations were performed with a 200-keV electron beam incident in the Si direction.

Thermoelectric measurement The electrical conductivity and carrier concentration of the samples were measured by the van der Pauw method and the Hall effect measurements, respectively. The Seebeck coefficient was measured by Seebeck coefficient measurement system (ADVANCE-RIKO Inc., ZEM-3). These electrical properties were measured after checking ohmic contact. We measured the electrical properties of the ZnO buffer layers/substrates annealed in O2 for 10 min and confirmed that the resistance of the ZnO buffer layers/Si(111) was sufficiently high, namely the electrical current did not flow in ZnO buffer layers on the substrates. To demonstrate the enhancement effect of S2 due to the introduction of the NWs, as a reference, we measured Film A with higher crystallinity (on the ZnO buffer layers/substrates) and Film B with lower crystallinity (on the substrates).

Energy barrier simulation The potential distribution (x) was obtained by numerically solving Poisson equation: d  d      q nx   N D (x) , dx  dx 






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where  is the dielectric constant of ZnO (8.340; 0 is the vacuum permittivity), q is the elementary charge, and ND+ is the donor concentration. The electron concentration n(x) was calculated from (x) using the Fermi-Dirac statistics. Donors are assumed to be entirely ionized, and their concentrations were ranged from ND,h+ = 2 × 1018 cm-3 to 8 × 1018 cm-3. To represent the potential energy barriers at the core/shell interface, the lightly doped thin layers with a thickness of 4 nm were inserted periodically with intervals of 30 nm. Their concentrations were set to be ND,l+ = ND,h+ - ND+, where ND+ = 1.8 × 1018 cm-3 was assumed to be independent of ND,h+. The detailed calculation is shown in Section S4 in the Supporting Information.

Thermal transport measurement The thermal conductivity of the specimens was measured using a front-detection/front-heating type pulsed-light-heating time-domain thermoreflectance method (TDTR)37. As additional preprocessing on the specimens, 100-nm-thick Mo was deposited on the ZnO surface as the reflective layer using DC magnetron sputtering because ZnO was transparent to both the pump laser beam and the probe laser beam. A pump laser beam of wavelength 1550 nm, pulse duration 0.5 ps, repetition rate 20 MHz, modulation frequency 200 kHz, and spot radius 35 μm was irradiated on the Mo surface. A probe laser beam of wavelength 775 nm, pulse duration 0.5 ps, repetition rate 20 MHz, and spot radius 15 μm was focused on the same spot as that from the pump laser. A lock-in amplifier was employed to detect the phase component of the thermoreflectance signal, hereafter called the lock-in phase signal. A lock-in phase signal was recorded as a function of the delay time of the probe laser pulse to the pump laser pulse up to 50 ns. The laser spot radii were much larger than the sample film thickness such that the heat deposited on the Mo surface can be assumed to diffuse one dimensionally in the cross-plane direction. The thermal conductivity of the specimens was determined by fitting a simulated transient lock-in phase signal38 based on the one-dimensional heat conduction equation to an experimental lock-in phase signal.


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ASSOCIATED CONTENT Supporting Information. The Supporting Information is available free of charge on the ACS Publication website at DOI: Thermoelectric properties, nanowire formation process, energy barrier simulation, optical transmittance property, material universality.

AUTHOR INFORMATION Corresponding Author E-mail: [email protected] Author Contributions T. I. and A. T. prepared the samples and carried out thermoelectric experiments. Y. N. is a principal investigator of this work. T. I. and Y. N. discussed the physics and wrote the manuscript. Y. K. and N. M. helped T. I. carry out the simulation of energy barrier. N. N. and Y. M. carried out the TEM experiment. Y. Y. carried out the thermal conductivity measurement and analyzed the data. Note The authors declare no competing financial interest.

ACKNOWLEDGEMENTS This work was supported in part by the JST CREST program. A part of this work was also supported by a Grant-in-Aid for Scientific Research A (Grant No. 16H02078), a Grant-in-Aid for Exploratory Research (Grant No. 15K13276), and for the JSPS Research fellow (17J00328). A part of ACS Paragon Plus Environment

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PLD in this work was supported by “Nanotechnology Platform Project (Nanotechnology Open Facilities in Osaka University)” of Ministry of Education, Culture, Sports, Science and Technology, Japan (No. S17-OS-0025).


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(26) Jood, P.; Mehta, R. J.; Zhang, Y.; Peleckis, G.; Wang, X.; Siegel, R. W.; Tasciuc, T. B.; Dou, S. X.; Ramanath, G. Al-Doped Zinc Oxide Nanocomposites with Enhanced Thermoelectric Properties. Nano Lett. 2011, 11, 4337-4342. (27) Loureiro, J.; Neves, N.; Barros, R.; Mateus, T.; Santos, R.; Filonovich, S.; Reparaz, S.; Torres, C.; Wyczisk, F.; Divay, L.; Martinsa, R.; Ferreira, I. Transparent Aluminium Zinc Oxide Thin Films with Enhanced Thermoelectric Properties. J. Mater. Chem. A 2014, 2, 6649-6655. (28) Saini, S.; Mere, P.; Honda, H.; Matsumoto, K.; Miyazaki, K.; Molina, L.; Hopkins, P. E.; Influence of Postdeposition Cooling Atmosphere of Thermoelectric Properties of 2% Al-Doped ZnO Thin Films Grown by Pulsed Laser Deposition. J. Electron. Mater. 2015, 44, 1547-1553. (29) Chattopadhyay, D.; Queisser, H. J. Electron Scattering by Ionized Impurities in Semiconductors. Rev. Mod. Phys. 1981, 53, 745-768. (30) Ellmer, K. Electrical Properties. In Transparent Conductive Zinc Oxide.; Ellmer, K.; Klein, A.; Rech, B., Eds.; Springer: Berlin, 2008; pp 38-39. (31) Ellmer, K. Resistivity of Polycrystalline Zinc Oxide Films: Current Status and Physical Limit. J. Phys. D 2001, 34, 3097-3108. (32) Minami, T.; Sato, H.; Ohashi, K.; Tomofuji, T.; Tanaka, S. Conduction Mechanism of Highly Conductive and Transparent Zinc Oxide Thin Films Prepared by Magnetron Sputtering. J. Cryst. Growth 1992, 117, 370-374. (33) Lorenz, M. E.; Kaidashev, M.; Wenckstern, H.; Riede, V.; Bundesmann, C.; Spemann, D.; Benndorf, G.; Hochmuth, H.; Rahm, A.; Semmelhack, H. C.; Grundmann, M. Optical and Electrical Properties of Epitaxial (Mg,Cd)xZn1-xO, ZnO, and ZnO(Ga,Al) Thin Films on C-plane Sapphire Grown by Pulsed Laser Deposition. Solid-State Electron. 2003, 47, 2205-2209.


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FIGURE CAPTIONS Figure 1 (a) Energy band structure of DCE interfaces with almost no point defects or misfit dislocations, where the wave function is coherently connected through the interfaces, leading to the quasi-ballistic transmission of the carriers with higher energy than energy barrier (EB) through DCE interfaces. Dopant concentration at the DCE interfaces is modulated lower than in the other region, leading to the control of Fermi energy (EF) at DCE interfaces, resulting in the formation of EB. The high-quality DCE interfaces bring higher Seebeck coefficient owing to the lack of defect scattering and relatively high electrical mobility. (b) Design concept of ENS with DCE interfaces. Carrier and phonon scatterings are controlled by crystal interfaces. Light is transmitted through ENS under the visible range. (c) Plan-view SEM images of ENS formed on thermally oxidized Si substrates. Cross-sectional view of the schematic is shown in the inset. (d) Photograph of ENS. (e) Dependence of transmittance of ENS on quartz glass on the optical wavelength.

Figure 2 Seebeck coefficients of ENS (inverted triangles), Film A (circles), and B (squares) as a function of (a)  and (b) n. Solid and open marks describe the samples with intrinsic and extrinsic donors, respectively. Black marks indicate other works (bulk22-25, nanocomposites26, and films27,28). Al-(Ga-) doped and Al and Ga co-doped ZnO are defined as AZO (GZO) and AGZO, respectively. In (a), trends of ENS and films are marked by the eye guides of the dashed lines. In (b), dotted and solid lines describe the calculated S based on Eq. (1), where scattering parameters were used as 1.5 (ionized impurity scattering) and -0.5 (short-range defect scattering), respectively.

Figure 3 (a) Cross-sectional TEM image of ENS. (b) Enlarged TEM image of the broken square in (a). (c) Enlarged TEM image of the solid square in (b). FFT patterns of (d) region A, (e) B, and (f) C in (c). Solid arrows in (d) and (e) indicate 1100 , 1101 , 1101 , 1100 , and 1101 spots. (g) Schematic of core and ACS Paragon Plus Environment

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shell parts of NW, and filling parts among NWs. (h) EDX profile at the region indicated by the solid blue line from P to Q in (b). Solid triangles and circles indicate Zn and O signal intensity relating to atomic concentration, respectively. Zn concentration is decreased at the epitaxial core/shell interface, indicating the formation of DCE interfaces with controlled dopants.

Figure 4 (a)  of each sample as a function of n. For reference, the solid line indicates the semiempirical curve obtained by fitting the experimental data for single-crystalline ZnO31. (b)  of HDENS as a function of T (100 K to 300 K). Inset in (b) shows the enlarged graph of (b) (100 K to 200 K). Experimental data is fitted as shown by the dotted curve in the inset of (b), where the fitted parameters of Nn (nonionized defect concentration), L (grain size), and EB (energy barrier) were 6.4 ± 0.9×1019 cm-3, 181 ± 40 nm and 21 ± 0.7 meV, respectively. Black marks indicate other works (bulk22,25, nanocomposites26, and films27,28,32,33).

Figure 5 (a) S of HDENS (solid diamonds) and Film A with n of ~ 4×1018 cm-3 at RT (solid circles) as a function of T (170 K to 320 K). ENS showed an anomalous trend; with increasing T, S decreased above 260 K while S increased at the temperature region of 170-260 K. (b) Enlarged graph in (a) (170 K to 220 K). Experimental data is fitted as shown by the dotted curve with the fitting parameters of EB (22 ± 1.1 meV). The reference curve of S without energy barrier (EB = 0 meV) (broken line). (c) Simulation result of the n dependence of EB. (d) Schematic of the energy barrier which decreases with increasing T. EC, EF, and Ev indicate the conduction band minimum, Fermi energy, and valence band maximum, respectively.


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Figure 6 S2 of each sample as a function of . Vertical dotted line points the maximum S2 value in the ZnO-based films ever reported27. Black marks indicate other works (bulk22-25, nanocomposites26, and films27,28).


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FIGURES

Figure 1

Figure 2


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Figure 3


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Figure 4

Figure 5


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Figure 6


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