Improving Photovoltaic Performance of ZnO Nanowires Based Colloidal Quantum Dot Solar Cells via SnO2 Passivation Strategy

Introduction

Colloidal quantum dots (CQDs) have attracted immense attention due to their applications in the field of optoelectronic devices such as lasers (Hoogland et al., 2006), light-emitting diodes (Wood et al., 2009), and photovoltaic devices due to their bandgap tunability and solution processing (Mcdonald et al., 2005; Nozik et al., 2010; Zhang et al., 2012; Kagan et al., 2016). Colloidal quantum dot solar cells (CQDSCs), a promising contender for new-generation solar cells, have gained more and more attention (Brown et al., 2011; Zhang et al., 2014b, 2016, 2018; Carey et al., 2015; Wang et al., 2016; Hori et al., 2018). CQDs exhibit unique optical and electrical properties, such as an adjustable absorption spectrum and efficient multiple exciton generation (MEG) (Nozik, 2005). To date, the record power conversion efficiency (PCE) value of PbS-based CQDSCs is more than 12% (Xu et al., 2018).

In CQDSCs, CQDs work as an active (or a light absorbing) layer, and wide bandgap semiconductors (e.g., ZnO, TiO₂, etc.) are employed as an electron transport layer (ETL). In this architecture, a depletion region was formed near the ETL/CQD interface, which plays a very important role in charge separation and extraction (Choi et al., 2009; Willis et al., 2012). To absorb all incident light, the required thickness of the CQD layer is about 1 μm (Wang et al., 2015; Rekemeyer et al., 2016), but the total of depletion width (about 300 nm) and carrier diffusion length (about 100 nm) of PbS CQDSCs is presently only several hundred nanometers (Zhitomirsky et al., 2013). Thus, thickness of the PbS CQD layer is usually limited to about 500 nm even though the light absorption is not enough. Bulk heterojunction (BHJ) architecture is an effective design for solving this absorption-extraction compromise. The BHJ architecture of CQDSCs is beneficial for extending the depletion region and increasing the thickness of the CQD absorption layer (Barkhouse et al., 2011; Kramer et al., 2012), which can improve both optical absorption and charge collection. One-dimensional ZnO nanowires (NWs) have been used in BHJ CQDSCs (Wang et al., 2013; Chang et al., 2015; Rekemeyer et al., 2016; Zhang et al., 2017), and the record PCE of ZnO NW/PbS CQD-based CQDSCs has reached 9.6 % (Rekemeyer et al., 2016). However, this strategy would result in large interfacial charge recombination at the ZnO NW/CQD interface because of the high surface defect density of ZnO NWs (Ehrler et al., 2013; Choi et al., 2017; Cheng et al., 2018; Ding et al., 2018), in turn reducing the PCE of the CQDSCs. Thus, reducing the surface defect density of ZnO NWs is of critical importance for improving the photovoltaic performance of CQDSCs.

Surface passivation of ETL has proved to be an effective way to eliminate surface defects of ETL and inhibit interfacial recombination at the ETL/light absorber interface in various photovoltaic devices (Hori et al., 2018; Zang et al., 2018). In our previous work, thin amorphous TiO₂ was used as a passivation layer to cover the surface of ZnO NWs, and the interfacial charge recombination at the ZnO NW/CQD interface was attenuated in the CQDSCs (Chang et al., 2015). However, that work did not clarify which defect species of ZnO NWs was suppressed. Furthermore, how the variation of the surface defect state of ZnO NWs will influence the performance of CQDSCs has not been fully revealed.

Recently, SnO₂ has attracted attention not only as an ETL but also as a surface passivation material in solar cells, owing to its unique properties such as wide bandgap (~3.6 eV), low chemical reactivity, high conductivity and good durability in ambient environment (Liu et al., 2016b; Khan et al., 2017; Wang et al., 2017). Herein we report the preparation of ZnO NW-based PbS CQDSCs and improve the PCEs of the CQDSCs from 5.6 to 7.8% through a thin amorphous SnO₂ surface passivation strategy. The type of defect states of the ZnO NWs was revealed, and we found that the surface defect density of ZnO NWs was reduced after being passivated by SnO₂. Due to the enhanced charge separation and reduced interfacial charge recombination efficiencies, the short-circuit current (J_sc), open-circuit voltage (V_oc) and fill factor (FF) of SnO₂ passivated ZnO (ZnO@SnO₂) NW-based CQDSCs largely improved more than that of the ZnO NW-based device, which corresponds to prolonging the effective carrier lifetime (τ_eff) in the former due to SnO₂ passivation. In addition, ZnO@SnO₂ NW-based CQDSCs can still work efficiently after being stored in air for more than 250 days. Our results indicate that SnO₂ surface passivation is an efficient way to improve the performance of ZnO NW-based CQDSCs.

Materials and Methods

Synthesis of PbS CQDs

PbS CQDs were synthesized by our previous reported method (Zhang et al., 2017). Briefly, a mixed solution of 6 mmol PbO, 15 mmol oleic acid (OA) and 50 ml 1-ocatadecene (ODE) was vigorously stirred and vacuum degassed at room temperature for 30 min and at 100°C for 2 h, respectively. Finally, a clear lead oleate solution was obtained. Under nitrogen flow, 10 mL hexamethyldisilathiane (TMS)/ODE mixed solution (contained 3 mmol TMS) was rapidly injected into the already synthesized lead oleate solution at 100°C. After that, when the temperature of PbS colloid solution reduced to 75°C, a CdCl₂ precursor solution [contained 1 mmol CdCl₂, 0.1 mmol tetradecylphosphonic acid (TDPA), and 3 mL oleylamine (OLA)] was added into the above PbS solution. Finally, PbS CQDs were collected and purified by a toluene/acetone/methanol solution cleaning and centrifugation process. The PbS CQD precipitate was dried under an N₂ flow and finally dispersed in octane (60 mg mL⁻¹).

Growth of ZnO@SnO₂ NWs

Firstly, ZnO NWs were grown on FTO substrates by a similar method as shown in the literature (Chang et al., 2015). A Sn²⁺ precursor solution was prepared by dissolving 0.2 M SnCl₂ in ethanol and stirring at room temperature for about 30 min. Then, this Sn²⁺ precursor solution was dropped onto ZnO NW substrate and spun-cast at 4,000 rpm for 60 s. Finally, the passivated-ZnO NW substrate was heated at 150°C for 10 min and annealed at 400°C for 30 min in air.

Fabrication of PbS CQDSCs

The PbS CQD active layer was deposited onto the as-prepared NW substrate by a layer-by-layer method as shown in previously (Chang et al., 2015; Nakazawa et al., 2019). After the desired thickness of the PbS CQD layer was obtained, a thin layer of Au (~100 nm) was deposited onto the PbS CQD layer by thermal evaporation as the metal electrode.

Characterization

Field-emission scanning electron microscope (FESEM) (JEOL, JSM-6340) and high-resolution transmission electron microscope (HRTEM) (JEOL, JEM- 2100F) techniques were applied to measure the lengths of the ZnO NW and thickness of SnO₂ passivation layer. X-ray photoelectron spectroscopy (XPS) (JEOL, S4 JPS-90MX) was used to measure the chemical state of elements, and photoelectron yield spectroscopy (PYS) spectra (Bunkoukeiki, BIP-KV205) were applied to determine the energy levels of NWs and CQDs. Photoluminescence (PL) spectra (JASCO, FP-6500) and UV-vis absorption spectra (HITACHI, U-3900H) were applied to measure the optical properties of NWs. The Hall effect was measured by using the Van der Pauw method at room temperature (TOYO ResiTest 8300). The photocurrent density-voltage (J-V) measurements of PbS CQDSCs were conducted using a Peccell solar simulator PEC-L10 with a Keithley 2400 source meter. The transient photovoltage (TPV) decay measurements were performed by using a Nd:YAG laser (wavelength 532 nm) with a pulse duration of 5 ns and a pulse frequency of 4 Hz.

Results and Discussion

Surface Passivation of ZnO NWs

After ZnO NW growth, SnO₂ layers were deposited onto the ZnO surface by spin-coating SnO₂ precursor containing SnCl₂. The SnO₂ layers were obtained through the hydrolysis and oxidation reaction of SnO₂ precursor in the following chemical equation (Marikkannan et al., 2015; Wang et al., 2017; Lu et al., 2018):

$\begin{array}{l}\text{SnC}{\text{l}}_2+2{\text{H}}_2\text{O}\to \text{S}\text{n}{(\text{O}\text{H})}_2+2{\text{H}}^{+}+2\text{C}{\text{l}}^{-}\\ \text{ Sn}{(\text{OH})}_2\to \text{S}\text{n}\text{O}+{\text{H}}_2\text{O}\\ 2\text{SnO }+{\text{O}}_2\to 2\text{S}\text{n}{\text{O}}_2\end{array}$

As shown in the above equation, the SnCl₂ which is absorbed on the surface of ZnO NWs firstly changed to Sn(OH)₂ through the hydrolysis reaction during or after the spin-coating process. It is worth noting here that the hydroxyl groups on the ZnO surface may benefit this reaction. Then, after annealing Sn(OH)₂ in open atmosphere, the SnO₂ layers were formed on the surface of ZnO NWs. Thus, the absorbed oxygen on the ZnO surface may be consumed when SnO₂ layers are formed. Figure 1 shows a cross-sectional SEM image and TEM image of ZnO NWs with and without SnO₂ passivation. It can be clearly seen that the length of ZnO NWs is approximately 1 μm (see Figure 1A), and the morphology of NWs has no change after SnO₂ passivation (see Figure 1B). Figures 1C,D show the TEM images of ZnO NWs and ZnO@SnO₂ NWs, respectively. It can be found that the average diameter of ZnO NWs is about 40 nm and a thin amorphous SnO₂ layer coated the surface of ZnO NWs with a thickness of about 2 nm.

FIGURE 1

Figure 1. The cross-sectional FESEM images of bare ZnO NWs (A) and ZnO@SnO₂ NWs (B). TEM images of bare ZnO NWs (C) and ZnO@SnO₂ NWs (D). The inset photograph is the HR-TEM image of NWs. TEM image of ZnO@SnO₂ NWs shows the surface of ZnO NWs covered with SnO₂ layer.

We employed XPS to investigate the composition and chemistry state of the surface of ZnO NWs and ZnO@SnO₂ NWs. As shown in Figure 2A, a characteristic peak of Sn 3d_5/2 at 486.1 eV, which belongs to SnO₂, was detected in ZnO@SnO₂ NWs (Khan et al., 2017). It indicates that SnO₂ was deposited on the surface of ZnO NWs. The result is well-consistent with the TEM images. In Figure 2B, after fitting, the O 1s peak of ZnO NWs and ZnO@SnO₂ NWs can be divided into three peaks. The binding energy peaks assigned to the O²⁻ state in Metal oxides (O_M), oxygen vacancy or defects (O_D) and chemisorbed oxygen or hydroxyl groups (O_C) are observed at 529.5, 530.9, and 531.8 eV, respectively (Zhang et al., 2014a; Wang et al., 2017). We define the area of these peaks derived from O_M, O_D and O_C as S_M, S_D and S_C. The ratio of S_D/S_M and S_C/S_M roughly corresponds to the weight of each surface defect species in ZnO NWs (Azmi et al., 2016; Yang et al., 2018b). After calculation, we found that the ratio of S_D/S_M was almost the same (from 0.61 to 0.59) before and after SnO₂ passivation, but the ratio of S_C/S_M was significantly decreased (from 0.14 to 0.05) after SnO₂ passivation. These results indicate that the density of surface defects which corresponds to O_C was reduced; this is mainly because of the reduction of chemisorbed oxygen species or hydroxyl groups on the surface of ZnO NWs after forming the SnO₂ layer.

FIGURE 2

Figure 2. XPS spectra of ZnO NWs and ZnO@SnO₂ NWs on glass: (A) Sn 3d and (B) O 1s with three resolved O bonding components (blue, green and pink lines).

PL Spectra of ZnO NWs and ZnO@SnO₂ NWs

To further reveal the effect of the SnO₂ passivation layer, we measured PL spectra of ZnO NWs and ZnO@SnO₂ NWs at room temperature. The PL spectra of ZnO NWs and ZnO@SnO₂ NWs were fitted by using Gaussian functions for understanding the origins of various luminescent components as shown in Figure 3. After fitting, the PL spectra can be divided into seven peaks (Peak A-G) and the centers of those seven peaks are at 3.28 (Peak A), 3.21 (Peak B), 2.98 (Peak C), 2.62 (Peak D), 2.30 (Peak E), 2.17 (Peak F), and 2.03 eV (Peak G). The narrow emission peak in the ultraviolet region (Peak A) corresponding to the intrinsic near-band-edge transition of ZnO was observed in both ZnO NWs and ZnO@SnO₂ NWs as well. The origin of the violet-to-blue (Peak B, Peak C, and Peak D) region emission is attributed to vacancies or interstitials of zinc, and these defect species were reported to form the shallow level trapping states in ZnO (Djurišić et al., 2010; Vempati et al., 2012; Ding et al., 2018). Notably, the emission ranging from green (Peak E) to orange (Peak G) is assigned to oxygen-related defects that form deep level trapping states in ZnO, especially at the surface. Green (Peak E), yellow (Peak F), and orange (Peak G) emissions are caused by oxygen vacancy, hydroxyl groups, and excess oxygen, respectively (Studenikin et al., 1998; Djurišić et al., 2010; Panigrahy et al., 2010; Vempati et al., 2012; Zhang et al., 2014a). We can estimate the approximate surface defect density of ZnO from these emissions (Choi et al., 2017; Yang et al., 2018b). In Figure 3A, the PL intensity of green-to-orange (from Peak E to Peak G) emission clearly decreased after SnO₂ passivation on the ZnO NWs. This result indicates that the surface defect density of ZnO NWs was reduced after SnO₂ passivation. Furthermore, comparing the area of each PL emission spectrum related to different peaks of ZnO NWs and ZnO@SnO₂ NWs, it can be found that percentages of Peak E, Peak F, and Peak G are decreased after SnO₂ passivation (as shown in Tables S1 and Tables S2). This is consistent with the result of XPS measurement, whereby the adsorbed oxygen species and hydroxyl groups on the surface of ZnO@SnO₂ NWs are lower than those of ZnO NWs. Therefore, SnO₂ passivation has the potential to suppress the charge recombination through deep level defect states of ZnO NWs, which originate from the absorbed oxygen and hydroxyl groups.

FIGURE 3

Figure 3. PL spectra with several fitted emission peaks (from Peak A to Peak G): (A) ZnO NWs (B) ZnO@SnO₂ NWs.

Energy Band Structure of ZnO@SnO₂ NWs

Surface passivation by SnO₂ not only has the potential to suppress interfacial recombination at the ZnO/CQD interface, but also affects the charge injection from CQDs to ZnO, so it is necessary to understand and select a favorable band alignment of ZnO and CQDs. First, we measured the absorption spectra of the ZnO NWs, ZnO@SnO₂ NWs, and PbS CQDs (see Figure S1). The optical band gaps (E_g) of both ZnO NWs and ZnO@SnO₂ NWs are about 3.29 eV (see Figure 4A), and the E_g of PbS CQDs is 1.25 eV (first exciton absorption peak is at 979 nm) (see Figure S2). Secondly, the valence band maximum (VBM) of ZnO NWs, ZnO@SnO₂ NWs, and PbS CQDs was measured by PYS technique (see Figure 4B and Figure S3). The VBM values of ZnO NWs, ZnO@SnO₂ NWs, and PbS CQDs, which are determined by the intersection of the tangent line and the baseline of the PY spectra, are −7.58, −7.59, and −5.14, respectively. Moreover, the conduction band minimum (CBM) of these three samples can be calculated (ZnO NWs: −4.29 eV, ZnO@SnO₂ NWs: −4.30 eV, PbS CQDs: −3.89 eV) from the E_g and VBM values. Finally, the energy band diagram including the Fermi level (E_F) of the ZnO NWs, ZnO@SnO₂ NWs and PbS CQDs was obtained, as shown in Figure 4C. We calculated E_F of the three samples from their carrier density, which were obtained from Hall effect measurement (as shown in Table S3). Herein, the effective mass of bulk ZnO (m_e ≈ 0.24 m₀) and PbS (m_h ≈ 0.075 m₀) were used for calculation (Brus, 1984; Prastowo et al., 2018). It can be seen that the E_F of ZnO@SnO₂ NWs (−4.38 eV) is a little higher than that of ZnO NWs (−4.46 eV), which is mainly due to the larger carrier density of ZnO@SnO₂ NWs. This indicates that the density of the deep level defect state in ZnO@SnO₂ NWs, of which the trapped electrons cannot be thermally excited at around room temperature, has decreased compared to that in ZnO NWs. In addition, the upshift of E_F leads to an increasing of the depletion region width and V_oc in the CQDSCs; there is a possibility that ZnO@SnO₂ NWs have a more significant band bending (Liu et al., 2016a; Choi et al., 2017; Yang et al., 2018a), which may lead to an efficient charge separation at the ETL/CQD interfaces. Therefore, ZnO@SnO₂ NW-based CQDSCs have the potential to achieve a higher PCE than ZnO NW-based devices because of the improvement in charge separation efficiency and suppression of ZnO/CQD interfacial recombination.

FIGURE 4

Figure 4. (A)Tauc plots obtained from the absorption spectra of ZnO NWs and ZnO@SnO₂ NWs. (B) PYS spectra of ZnO NWs and ZnO@SnO₂ NWs. (C) Schematic energy band diagrams of ZnO NWs, ZnO@SnO₂ NWs, and PbS CQDs.

Characterization and Photovoltaic Performance of ZnO@SnO₂ NWs Based CQDSCs

Figure 5 shows the schematic diagram of the architecture of CQDSCs and the J-V curves of the CQDSCs; their corresponding performance parameters are shown in Table 1. Cross-section SEM imaging of the ZnO@SnO₂ NW/PbS CQD/Au CQDSCs is shown in Figure S3. Compared to the ZnO NW-based device, ZnO@SnO₂ NW-based CQDSCs significantly exhibit superior performance. The deposition cycle of the SnO₂ layer was optimized and the device based on 2 deposition cycle of SnO₂ exhibits the best performance (see Figure S4). As expected from the band alignment in the CQDSCs, the open-circuit voltage (V_oc) was improved due to the E_F upshift of ZnO@SnO₂ NWs and the reduced deep level defect density of ZnO. Meanwhile, enhancement of short-circuit current (J_sc) and fill factor (FF) of ZnO@SnO₂ NW-based CQDSCs can be explained by the improvement of charge separation and collection efficiencies, which were caused by the larger band bending and increased depletion region width. IPCE spectra of these devices supported this result (see Figure S5). Particularly, the FF of the ZnO@SnO₂ NW-based CQDSCs improved about 19% compared to that of the ZnO NW-based device, which confirms that the charge recombination was considerably suppressed. It is mainly caused by the reduction of surface defects on ZnO NWs and the reduced leakage current after SnO₂ passivation. Consequently, the PCE of ZnO@SnO₂ NW-based CQDSCs (7.8%) was enhanced about 39% more than that of the ZnO NW-based device (5.6%).

FIGURE 5

Figure 5. (A)Schematic diagram of the architecture of ZnO@SnO₂ NW-based CQDSCs. (B) The J-V curves of ZnO NW CQDSCs with or without SnO₂ passivation.

TABLE 1

Table 1. Performance detail of the ZnO NW-based CQDSCs with and without SnO₂ passivation layer^a.

To investigate the detail of the charge recombination mechanism in ZnO NW CQDSCs, we have evaluated the diode ideality factor (n) of two types of CQDSCs (Figure 6A). As a reliable parameter of solar cells, the value of n can reflect the charge recombination process in CQDSCs. When interfacial recombination or trap-assisted recombination take over a majority in the recombination process, the value of n should be theoretically larger than unity (1 < n < 2) (Zhang et al., 2018). The light-intensity (P_light) dependence of the V_oc has the following relation:

$\begin{array}{rc}{V}_{\text{oc}}\propto \frac{nkT}{q}& (1)\end{array}$

where n is the diode ideality factor, k is the Boltzmann constant, T is the temperature, and q is elementary charge (Cowan et al., 2010; Gao et al., 2014; Ding et al., 2018). By fitting the V_oc vs. light intensity plotted points, the values of n are determined to be 1.73 and 1.21 for ZnO NW-based CQDSCs and ZnO@SnO₂ NW-based devices, respectively (see Figure 6A). Interfacial recombination or trap-assisted recombination correlated with the deep level defect density in CQDSCs (Ding et al., 2018). In this work, the only difference between the two types of device is the presence or absence of SnO₂ passivation, so these results strongly confirm that ZnO/CQD interface charge recombination via deep level defect states of the ZnO NWs was suppressed after SnO₂ passivation. To further examine the charge recombination process in CQDSCs, we have measured transient photovoltage (TPV) decay. Figure S6 shows the TPV decay curves of ZnO NW CQDSCs with and without SnO₂ passivation. TPV decay of ZnO@SnO₂ NW-based CQDSCs is slower than that of ZnO NW-based CQDSCs, which confirms that the charge recombination in ZnO@SnO₂-based devices was weakened. We also investigated the effective carrier lifetime (τ_eff) in CQDSCs, and τ_eff can be defined by the following equation:

$\begin{array}{ll}{\tau }_{\text{eff}}=-\frac{\left(\frac{kT}{q}\right)}{\left(\frac{d{V}_{\text{oc}}}{dt}\right)}& (2)\end{array}$

Here, k is the Boltzmann constant, T is the temperature, and q is elementary charge (Zaban et al., 2003; Zhang et al., 2016). As shown in Figure 6B, it shows that the τ_eff of the CQDSCs was significantly increased after SnO₂ passivation, especially at high voltage regions (over 0.25 V). This result demonstrates that the charge recombination rate, which is the reciprocal of the carrier lifetime, of the ZnO@SnO₂ NW-based device decreased in the fast decay process. This result further confirms that SnO₂ passivation can significantly attenuate the charge recombination at the ZnO NW/CQD interface, and then enhance the performance of the devices, such as V_oc, FF, and PCE.

FIGURE 6

Figure 6. (A) Light-intensity dependence of V_oc for CQDSCs with and without SnO₂ passivation. The solid black and red lines are the fitting curves. (B) τ_eff of CQDSCs with or without SnO₂ passivation that was calculated from the TPV decay curves of the devices which were shown in Figure S6.

For CQDSCs, stability is important as well as PCE. Thus, the long-term stability of those devices was also evaluated, resulting in them being stored under ambient condition over half year. Both ZnO NW- and ZnO@SnO₂ NW-based CQDSCs exhibit excellent stable properties after being stored in air for 250 days, as shown in Figure 7. Therefore, it means that SnO₂ passivation can improve the PCE of CQDSCs while it does not impair the long-term stability of devices.

FIGURE 7

Figure 7. Stability evaluations of ZnO NWs and ZnO@SnO₂ NWs based CQDSCs.

Conclusions

In summary, we used SnO₂ to passivate the surface of ZnO NWs and investigated the influence of SnO₂ passivation on the defect species of ZnO and the photovoltaic performance of the ZnO NW/PbS CQDSCs for the first time. We found that SnO₂ passivation can reduce the deep level defect density of the ZnO NWs, which is attributed to the absorbed oxygen and hydroxyl groups on the surface of ZnO NWs, and in turn reduces the surface recombination and leads to an upward shift of E_F. By SnO₂ passivation, the PCE of CQDSCs was improved from 5.6 to 7.8%. The SnO₂ passivation layer can effectively attenuate the charge recombination at the ZnO NW/CQD interface and increase the effective carrier lifetime in CQDSCs. More importantly, SnO₂ passivation enhanced the PCE of CQDSCs, while it does not impair the long-term stability of devices. Our study demonstrates that reducing the interfacial recombination at the ETL/CQD interface by using a wide bandgap semiconductor as a passivation layer is an effective way to enhance the performance of CQDSCs.

Supporting Information

Fitting data of PL spectra; optical absorption spectra of ZnO NWs, ZnO@SnO₂ NWs, and PbS CQDs; PYS spectrum of PbS CQDs treated with CTAB; Absorption and IPCE spectra of devices; SnO₂ deposition cycle dependent photovoltaic performances of QDSCs; SEM image of CQDSCs; light-intensity dependence of J_sc for CQDSCs with and without SnO₂ passivation; TPV decay curves of ZnO NWs CQDSCs with and without SnO₂ passivation.

Author Contributions

SO, YZ, and QS: conceived and designed the experiments; SO, HY, YK, and MH: performed the experiments; SO, YZ, TT, KK, and QS: analyzed the data; KY, SH, RW, and QS: contributed reagents, materials, analysis tools; SO: wrote the paper; YZ and QS: corrected the paper.

Funding

This work was supported by the Japan Science and Technology Agency (JST) CREST and PRESTO programs, Beijing Advanced Innovation Center for Future Urban Design, Beijing University of Civil Engineering and Architecture (Grant UDC2018031121), and the MEXT KAKENHI grant (Grands 26286013, 17H02736).

Conflict of Interest Statement

The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Supplementary Material

The Supplementary Material for this article can be found online at: https://www.frontiersin.org/articles/10.3389/fenrg.2019.00011/full#supplementary-material

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