Article pubs.acs.org/journal/apchd5
White Light Emission from Black Germanium Satish Laxman Shinde,† Thang Duy Dao,† Satoshi Ishii,† Li-Wei Nien,† Karuna Kar Nanda,*,§ and Tadaaki Nagao*,†,‡ †
International Center for Materials Nanoarchitectonics (MANA), National Institute for Materials Science (NIMS), Tsukuba, Ibaraki 305-0044, Japan ‡ Department of Condensed Matter Physics Graduate School of Science, Hokkaido University, Kita-10 Nishi-8 Kita-ku, Sapporo 060-0810, Japan § Materials Research Centre, Indian Institute of Science, Bangalore 560012, India S Supporting Information *
ABSTRACT: We demonstrate the nearly perfect absorption and quantum dots-mediated enhanced visible light emissions from defect engineered Ge nanopyramids or black germanium. High-resolution 3D photoluminescence (PL) imaging of the pyramid structure elucidated the position dependency of defects and their emission: Stronger photoluminescence yield was observed at the nanopyramid tips, which is correlated to the efficient light nanofocusing at the tips where increased defect density and roughness at the interface between Ge and surface oxide (GeO2) also takes place. Furthermore, the visible light emissions from this GeO2/Ge interface can be enhanced ∼15-fold when CdTe quantum dots (QDs) are adsorbed on the GeO2/Ge system. The enhanced luminescence of our structure can be attributed to the extraordinary light harvesting property of pyramid structure; strong antireflection effect, pronounced defect formation at the nanopyramid tips, and anomalous resonant energy transfer between GeO2 defects and CdTe QDs. The proposed methodology can be applied to other nanostructured wide bandgap materials to turn them into solar light harvesters and bright white light-emitting phosphors. KEYWORDS: germanium oxide, defects, photoluminescence, visible light emission, quantum dots
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defect containing WBG oxide phosphors has very low absorption efficiency for visible light, making them impractical to coat on blue LED to generate white light. Therefore, compatibility of strong absorption (for UV/blue light) together with efficient PL are major requirement for the practical phosphor materials. Germanium dioxide (GeO2) is a WBG oxide with fascinating electronic and optical properties as well as high thermal stability. It is considered as a promising dielectric material for Ge gate stack formation in Ge microelectronics,20 optical waveguides, optoelectronic devices, and catalysts.21−24 The most stable phase of GeO2 is the α-quartz (trigonal) phase, which has an indirect bandgap of 5.4−5.9 eV. 12 GeO2 nanostructures have been reported to have violet and blue
hite light-emitting-diodes (LEDs) have become the major light sources in the household and industrial lighting markets. Through the development of LEDs in the near-ultraviolet and blue wavelengths, solid-state lighting (SSL) is considered as a major promising technology for general illumination.1−5 The most prominent example of SSL is gallium-nitride-based white LEDs, where the narrowband ultraviolet/blue emissions are converted into white light with broader and longer wavelengths by the application of inorganic phosphor materials.4−7 Various oxide- and sulfide-based phosphor materials, quantum dots (QDs), and organic dyes have been developed, and several strategies have been proposed to generate white light in combination with ultraviolet/blue LEDs.8−11 An alternative to rare earth or transition metal-based phosphors is defects containing wide-bandgap (WBG) oxides such as silica, zirconates, aluminates, Ga2O3, ZnO, GeO2, and so on. These materials have recently been recognized for their advantages in generating white light.12−19 However, most of the © 2017 American Chemical Society
Received: March 6, 2017 Published: June 14, 2017 1722
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Figure 1. SEM images of (a−c) etched Ge wafer and (d−f) etched Ge wafer annealed at 400 °C for 2 h at different magnifications, respectively. Raman spectra of etched (g) and annealed (h) Ge wafers at different positions, respectively.
luminescence,23−26 and this is believed to be due to the disorders at the GeO2/Ge interfaces and oxygen vacancy levels present in GeO2.14−18 However, due to its wide band gap nature, the absorption of GeO2 in the visible region is rather limited. It is known that by the formation of nanostructrures at the Ge and Si surfaces, high density of impurity defects and structural defects in the lattice are formed, and the light absorption are significantly increased in the visible to nearinfrared region.27,28 There exists several reports on preparing various structures of GeO2 and its variants by physical methods, including synthesizing whiskers, nanotubes, nanorods, and nanocones by laser ablation,19 physical evaporation and carbothermal method,25,26 vapor phase reaction,29 and thermal evaporation.13 Here, we report near-unity perfect absorption and defectmediated visible light emissions from black germanium (GeO2/ Ge nanopyramids) that strongly enhanced by the presence of CdTe QDs. The facile, cost-effective, wet-chemical etching technique developed to tune the GeO2/Ge nanopyramids geometry to increase the surface-to-volume ratio and to achieve near-perfect absorption for trapping the incoming light. The defect concentration was optimized in parallel with the surface morphology and annealing conditions in an ambient (air/inert) environment. The GeO2/Ge nanopyramids were observed to be highly luminescent compared to flat and other GeO2/Ge nanostructures (flat-tipped pyramids, nanoparticles, columner, porous) and showed the highest emissions at the tip of pyramid, which had the highest roughness, defect density, as well as the highest field intensity. In parallel to this finding, we observed that the CdTe QDs, which are inherently narrowband orange light emitters, enhances the visible light emissions when they are loaded on WBG GeO2 nanostructures. Together with the perfect absorptivity arising from the nanomorphology and the electromagnetic field confinement due to the high refractive index of Ge, we clarified that the emissions are strongly
enhanced by the interfacial resonant energy transfer from CdTe QDs to the defect levels in GeO2 based on our spectroscopic characterizations and electromagnetic simulations. The results of this study present a novel strategy to realize hybrid materials that exhibit strong broadband photoluminescence (PL) by engineering interface defects and interface nanostructures of oxide nanomaterials. This approach can also be utilized for various light-emitting, energy conversion, and photodetector applications. When Ge wafers are etched in aqueous solutions containing different concentrations of HF and H2O2 at different temperatures, various shapes/morphologies and ordered arrays of column-like nanostructures are obtained.13 Figures 1a−f and S1 show scanning electron microscopy (SEM) images of Ge wafers etched in HF/H2O2/H2O solution under different conditions, that is, the molar ratio, temperature, and duration. All of these Ge wafers showed rough surface morphologies that depended on the etching temperature. Tuning the etching conditions resulted in various morphologies, such as sharp- or flat-tipped pyramids, nanoparticles, nanorods, and porous films, as shown in Figures 1 and S1b−h. Table S1 presents the detailed experimental parameters for tuning the nanostructres. For example, the arrays of nanopyramids were obtained by treating Ge wafers in H2O2 at 90 °C for 10 min and then dipping them in HF[7 M]/H 2 O 2 [0.1 M]/H 2 O = 10.0:0.25:14.75 at 50 °C for 3 h. These nanostructures are confirm to be very stable and unchanged for several months in air. The typical height and diameter of the pyramids (see Figure 1c) were ∼7 and ∼3 μm, respectively. To determine the chemical compositions of the products, the as-prepared samples were investigated by Raman scattering and X-ray photoelectron spectroscopy (XPS). Figure S1i shows the Raman spectra of various nanostructures. The Ge wafer showed a strong Raman peak at 300 cm−1, which is the characteristic peak of the Ge bulk phonon.15,16,30,31 The Raman spectra from 1723
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various morphologies also showed a strong peak at 300 cm−1 and confirmed the presence of Ge in the nanostructures. For further study, we chose the pyramid-shaped samples because they had highly ordered cone structure, which effectively confined/trap light compared to the other nanostructures. To examine the stability and change in composition depending on the annealing temperature, we annealed pyramid-shaped etched Ge nanostructures (Figures 1a−c) at different temperatures for 2 h under oxygen-deficient conditions and confirmed that the pyramid-shaped morphology was retained up to 400 °C (Figures 1d−f). Figures 1g,h show Raman spectra taken at the tip and side walls of the pyramid before and after annealing. At the side walls, the Ge−Ge phonon peak at 300 cm−1 was sharp both before and after annealing. This strongly suggests the presence of pristine crystalline Ge, and this along with the weak peak at 167 cm−1, which indicates Ge−O−Ge stretching of trigonal GeO2, suggests low amounts of GeO2 at the walls.15,16,30,31 At the tip of the pyramids, the Ge peak was broader and weaker. The peak became further broader and weaker after annealing in oxygen-deficient conditions, while Ge−O−Ge peak becomes relatively stronger. This broadening and significant tailing to the lower energy side is a typical feature originating from momentum relaxation due to the surface/interface roughness and is associated with the etching or oxidation process.30−33 The pyramids center should be etched or oxidized more heavily than at the side walls of the pyramids.30−33 The annealed samples had more disorders caused by oxidation and strain at the Ge interfaces, which may lead to the formation of partially oxidized Ge or GeO2 with a high defect density. The Raman measurements confirmed the presence of a heavily roughened GeO2/Ge interface at the tip and a smoother GeO2/Ge interface at the side walls of the pyramids. Thus, a higher density of oxygen defects would naturally be expected at the tip. Figures 2a,b show the XPS analysis of the Ge structures. Figure 2a shows the XPS spectrum of Ge 3d core-level electrons for a typical Ge wafer, etched Ge, and etched Ge annealed at 400 °C for 2 h. For all samples, the peaks are broad in shape and can be deconvoluted into three components of the germanium oxidation states. The first was assigned to the elemental Ge around 29.9 eV, and the other two peaks exhibit binding energy peak shifts with respect to elemental Ge. The other peaks at 31.5−32.6 eV were assigned to the Ge monoxide (GeO) state and dioxide state (GeO2).34,35 Both pristine Ge and etched Ge exhibited the components associated with the GeO and GeO2 phases. The elemental Ge peak become stronger after etching, since surface oxide layer is removed by etching. After annealing, however, the GeO component reduces and a new peak emerged at a higher binding energy (33.6 eV) that was associated with GeOx (x ≈ 2) and a high defect density. Figure 2b shows the O 1s core-level spectra. The O 1s spectrum for the Ge wafer exhibited major peaks around 530.6 and 531.8 eV corresponding to suboxides of Ge (GeO and GeOx (x ≈ 1)) and low intensity around 533.1 eV for GeO2. When the Ge wafer was etched, a strong peak was observed at 531.5 eV (GeO), and a weak peak was observed at 532.6 eV (GeO2). Due to etching, a suboxide phase at 530.1 eV (GeOx (x ≈ 1)) and the GeO2 phase formed.34,35 After annealing, the majority of the oxide was converted into GeO2 yielding a strong peak at 532.4 eV. The low binding energy component at 531.1 eV and high binding energy component at 533.4 eV were assigned to GeO and GeOx (x ≈ 2) with a high defect
Figure 2. XPS spectra of etched Ge wafers annealed at 400 °C for 2 h: (a) Ge 3d and (b) O 1s. (c) PL spectra taken at tip and side walls of pyramid-shaped samples before and after annealing. The excitation wavelength was 470 nm.
density.34,35 The broadening in the peak was due to structural disorders and defects in the oxides. Overall, the XPS analysis confirmed the presence of suboxide GeOx (x ≈ 1 or 2) phases rich with defects at the roughened and strained Ge surfaces. Among the various shape/morphologies of fabricated Ge nanostructures, the pyramid-shaped samples showed weak PL, while the nanoparticle- and columnar-shaped samples were nonluminescent (see Figure S1j). In the case of pyramids, the tip region contained a rough GeO2/Ge interface. The GeO2/ Ge interfacial defects and oxygen vacancies defects in GeO2 should be responsible for the luminescence from the etched Ge.14,29,36−40 In the case of the other structures, most of the surface region contained mainly pristine Ge or GeO2 based on the Raman study (Figure S1i); these are nonluminescent in the visible region. The PL study confirmed that only the pyramid shapes were luminescent in the visible region. Figure 2c shows the PL spectra taken at the tip and side walls of the pyramid of GeO2/Ge samples before and after annealing. The prepared asetched sample showed very weak visible light emission with negligible position dependence. After annealing, the tips of the pyramids showed strong visible light emissions, and the side walls showed a blue shift in the PL peak position. The different defect levels due to nonstoichiometry and oxygen vacancies in GeO2 were responsible for the variation in PL peaks at different positions. The combined effects of the roughened GeO2/Ge interface and high defect density of oxygen vacancies made the tips of the pyramids more luminescent. To associate the defect-related luminescence to the spatial information, we performed PL imaging of the etched Ge samples. Figure 3a−c show optical micrograph, SEM, and PL mapping images, respectively, for the annealed sample. The 1724
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Figure 3. Optical, SEM, and 3D PL mapping image of the annealed samples (a−c) before and (d−f) after the addition of CdTe QDs. All of the images were taken from exactly the same place. (The scale bars for c and f are normalized) (g) PL spectra of an as-etched sample and an annealed sample before and after the addition of CdTe QDs. The excitation wavelength was 470 nm. (h) Excitation-dependent PL of CdTe/GeO2/Ge. The excitation wavelengths were 470, 570, and 670 nm. The dotted curve shows the PL of CdTe QDs suspension before loading on the GeO2/Ge nanopyramids. The PL curves for pristine GeO2/Ge at an excitation wavelength of 470 and 670 nm are also shown (black and green lines at the bottom).
QDs on a quartz glass substrate (see Figure S2a in Supporting Information). The CdTe QDs had a lifetime of 7.14 ns on a quartz glass substrate. However, the lifetime decreased significantly to 1.94 ns when they were placed on the GeO2/ Ge nanopyramids, indicative of strong interaction with the substrate. According to the energy band diagram in Figure S2b there will be no energy transfer between CdTe and quartz glass (i.e., SiO2) due to the absence of a compatible interband transition between the two. Because GeO2/Ge nanopyramids support various interband transitions between their energy levels (Figure S2c), CdTe QDs can nonradiatively transfer their excited energy to the interband excitation in GeO2/Ge. For the better understanding of the role of CdTe on PL enhancement, we examined the excitation-dependent PL of a CdTe/GeO2/ Ge sample. When CdTe/GeO2/Ge was excited by light with energy above the bandgap of CdTe, PL enhancement was observed, as shown in Figure 3h. When the excitation energy was below the bandgap of CdTe (576 nm), no PL enhancement was observed, which confirmed that photoexcitation of CdTe helped to enhance the PL of the Ge samples. When the excitation energy was below the bandgap of CdTe, there will be no excited e−−h+ pair available to transfer its energy from CdTe to GeO2/Ge, which resulted in weak luminescence originated only from GeO2/Ge. Now, we discuss the possible mechanism for the resonance energy transfer from CdTe that enhances the luminescence intensity of GeO2/Ge nanopyramids. Because trigonal GeO2 is an indirect bandgap material, the luminescence observed in the current study is believed to be due to the roughened GeO2/Ge
three images were taken at exactly the same position on each sample. As expected, only the tip (center) part of pyramidshaped region was bright (yellow color), that is, luminescent. The luminescence decreased away from the pyramid’s tip, which suggests a variation in the defect density along the side wall. The oxidation increased the defect density, which resulted in a substantial increase in the PL intensity especially at the tip of the pyramid. Surprisingly, when the Ge nanostructures were coated with CdTe QDs (Figures 3d−f), the PL intensity was further enhanced remarkably. The average size of the QDs is 3.5 nm and the band gap is 2.1 eV.41 When excited with 470 nm, the QDs exhibit an orange luminesce centered around 630 nm that corresponds to the band edge transition.41 In the PL mapping image (Figure 3c and f), the tip (center) region and side wall were more luminescent compared to the peripheral flat surface. This trend is similar to that of the bare GeO2/Ge surfaces. As shown in Figure 3g, the PL of the etched Ge surface was very weak but increased 2-fold with annealing and further increased 15-fold with the addition of the CdTe QDs. The absolute quantum yield at an excitation wavelength of 470 nm for bare, annealed, and CdTe QDs coated GeO2/Ge pyramids are 0.14, 0.32 and 3.44%, respectively. The variation in the PL intensity at among the tip/side walls was within 5%. Surprisingly, the luminescence signature for the CdTe QDs themselves was not observed in the enhanced luminescence. This strongly suggests that CdTe QDs nonradiatively transfer most of their energy to GeO2/Ge. This was further confirmed by the PL lifetime measurement, which exhibited a shortened emission lifetime for the QDs on GeO2/Ge compared to the emission lifetime of 1725
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Figure 4. (a) Energy band diagram of CdTe QDs/GeO2/Ge for enhanced luminescence. The bandgap of GeO2 is 5.7 eV. (b) Schematic of a CdTe/ GeO2/Ge nanostructure (nanocones or nanopyramids) modeled in the simulations. (c) Experimental (solid line) and simulated (dash line) absorptance spectra of GeO2/Ge nanopyramids with different cone diameters. Inset of (c) shows the SEM image of pyramids array. (d) SEM image of individual pyramid with diameter of ∼3 μm. (e) Simulated absorption cross section of a nanopyramids excited at 570 nm.
trapped in these defect levels, then relaxes into lower levels, that results in weak visible luminescence from GeO2/Ge. Based on our experimental results, Figure 4a presents the proposed energy level diagram for GeO2/Ge pyramids and CdTe. When excited at 470 nm blue light, CdTe QDs absorb significant amounts of excitation energy and nonradiatively transfer it through Förster resonance energy transfer (FRET)42 to the defect levels in GeO2/Ge as confirmed from lifetime measurement (Figure S2a). The transfer energy generate significant amount of excited carriers which partly occupy some of the nonradiative trap states present in GeO2, which
interface and various emission tips in GeO2, such as oxygen vacancies (VO), germanium−oxygen vacancy centers (VGe−O), defects in oxides and interfacial defects between GeO2 and Ge (VX−VDef), and oxygen deficiency centers (VX−VODC).36−40 Thus, the very weak visible light emission of the GeO2/Ge interface is attributed to the momentum relaxation and localized quantized levels at the roughened Ge interface as well as defects in GeO2 and interface states between Ge and GeO 2 .14,25,36−40 Because of these various defects and roughened interface, trap levels exist at the GeO2/Ge interface. When excited with high-energy light, the excited electrons are 1726
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defect- and roughness-controlled nanopyramid morphology of GeO 2/Ge was wet-chemically synthesized and showed moderate visible light luminescence from the GeO2/Ge interface as well as perfect broadband absorption from the ultraviolet to near-infrared region. The 3D PL imaging confirmed that the strongest luminescence was from the tip of the pyramids, which had the highest roughness and defect density. Furthermore, bright visible light was obtained when the pyramids were coated with CdTe QDs. The spectral features and lifetime of the emissions imply that QDs on nanopyramids effectively act as a strong light absorber and transfer their excitation energy to GeO2/Ge to dramatically enhance the visible light emission. We anticipate that the results demonstrated here will pave a way to convert nonluminescent WBG materials into strongly luminescent ones by engineering the interface, roughness, and defects, as well as tailoring the nanostructures of semiconductor/oxide heterointerfaces to realize perfect absorptivity and light trapping.
increases the radiative rate of excited electron relaxation in GeO2/Ge defect levels. As a consequence, the luminescence of GeO2 increases and that of QDs decreases. In order to probe whether CdTe QDs, when loaded on etched Ge bring more defects or not, we measured the Raman spectra before and after loading CdTe QDs on etched Ge. The Raman spectra (Figures 1h and S3) show negligible variation with and without CdTe QDs coated etched Ge which confirms that CdTe QDs has minimal effect on defect generation in etched Ge. However, CdTe QDs help to absorb more light and transfer the energy to the GeO2/Ge and enhances the luminescence intensity. Overall, the experimental results can be explained by the energy transfer mechanism, where CdTe acts as the donor, GeO2 acts as the acceptor, and the enhancement of the luminescence is dependent on the GeO2/Ge defect density. Together with the interfacial energy transfer mechanism, the surface morphology of the Ge pyramids also plays an important role in the enhanced luminescence. In order to investigate the morphology effect of the etched samples, the experimental absorptance spectra for the Ge wafers and etched samples for different sizes of GeO2/Ge nanopyramids (for schematic see Figure 4b) and a thin film of GeO2/Ge obtained by simulation were compared, as shown in Figure 4c. For the Ge wafer and flat thin film, the absorptance was only around 60%. It further increased above 94% for the etched and nanopyramid films. As the diameter of the pyramids decreased, the absorptance increased as shown in Figure 4c. The nanopyramids effectively work as graded index interface to reduce the large refractive index contrast between air and the substrate, resulting in enhanced light-trapping and absorptance at the interface.27,28 The experimentally measured absorptance of the etched Ge matched the simulated spectra of 2−3 μm GeO2 /Ge nanopyramids. The small difference in adsorption was due to the nonuniform rough surface morphology (Figure 4d) and variation in the dielectric constant of GeO2/Ge. This result suggests that the pyramid-shaped GeO2/Ge structures act as an antireflection layer or perfect absorber for visible light harvesting. In addition, the simulated absorption mapping for the 2 μm diameter pyramids obtained at an excitation wavelength of 570 nm (Figure 4e) suggests that the tips of the pyramids absorb more than the side walls. This also support the scenario that the higher absorption at the tip of pyramids leads to stronger excitation and, thus, stronger PL intensity at the tip, as observed in the PL mapping (Figure 3c). Overall, due to the high antireflection property, high absorption efficiency, and high defect density of the Ge pyramids, the PL intensity of the annealed samples showed highest enhancement when they were sensitized with CdTe QDs. One of the major implications of this study is the application of defects containing luminescent WBG nanostructured materials as light harvester and bright white-light-emitting phosphors. With the presence of light-absorbing QDs and antireflecting nanomorphology, this type of WBG material can efficiently convert ultraviolet/blue light into visible light. It may be adopted as a spectral converter for solar cells and photocatalysis, infrared detection, and so on. The wavelength conversion efficiency can be further enhanced by tuning the morphology, defect level, and QD bandgap. In summary, we demonstrated extraordinary light harvesting property and quantum dots mediated enhanced visible light emission from nanopyramids of Ge, based on interface engineering and perfect absorption by controlling the nanomorphology of the semiconductor/oxide heterointerface. The
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METHODS Synthesis. Single-crystalline Ge (100) wafers with dimensions of 3 mm × 3 mm were washed ultrasonically with methanol and then acetone before being rinsed with deionized water. The etching process was carried out in a 50 mL Teflon beaker. The mixed aqueous solution of HF and H2O2 was used as the etchant, and the concentrations of HF and H2O2 were varied from 4.8 to 7 and 0.1 to 0.2 M, respectively. The Teflon beaker were filled with 30 mL of etchant and kept at 25−50 °C. Then, the cleaned Ge wafers were placed into the etchant. After certain etching duration times, the wafers were taken out, rinsed with distilled water, and dried under ambient and room temperature conditions. Finally, the etched Ge wafers were annealed at 400 °C in an oxygen-deficient environment for 2 h. The CdTe QDs were prepared by the hydrothermal method. The detailed experimental conditions have been reported elsewhere.41 The 5 μL of CdTe QDs was drop-casted on 3 mm × 3 mm etched Ge wafers and dried on a hot plate at 90 °C for 30 min. Characterizations. The morphology and phase of the Ge microstructures were studied by SEM (Hitachi FE-SEM SU8000). Raman spectroscopy and PL imaging were performed with a WITec system 300 alpha. The light source for Raman and PL are 532 nm Nd:YAG laser and high power supercontinuum laser (SuperK EXTREME, NKT Photonics) combined with a wavelength selector (SuperK SELECT, NKT Photonics), respectively. The power density for the excitation wavelengths of 470, 570, and 670 nm are 1.5, 50, and 70 kW/ cm2, respectively. The lifetime was measured with the identical WITec system attached to a PDM series time-correlated single photon counting detector with a resolution of 30 ps. XPS spectra were recorded by Micro-XPS (Quantera SXM). The absorptance spectra were simulated by using DiffractMOD package (Synopsys’s RSoft), which is based on rigorous coupled-wave analysis. The absorption mapping was simulated using finite-difference time-domain method (FullWAVE, Synopsys’s Rsoft). The dielectric constant of GeO2 was 2.8, and the dielectric function of Ge was taken from literature.43
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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/acsphotonics.7b00214. 1727
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Supporting figures and table (PDF).
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AUTHOR INFORMATION
Corresponding Authors
*E-mail:
[email protected]. *E-mail:
[email protected]. ORCID
Satish Laxman Shinde: 0000-0002-0353-3705 Satoshi Ishii: 0000-0003-0731-8428 Karuna Kar Nanda: 0000-0001-9496-1408 Tadaaki Nagao: 0000-0002-6746-2686 Notes
The authors declare no competing financial interest.
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ACKNOWLEDGMENTS This work was partly supported by JSPS Invitation Fellowship Program S14057 and JSPS-KAKENHI (16K17496, 16F16315, 16H06364, 15K17447, 17H04801) and CREST “Phase Interface Science for Highly Efficient Energy Utilization” (JPMJCR13C3), Japan Science and Technology Agency. The authors thank the Materials Analysis Station, Sengen, NIMS, Tsukuba, for providing the XPS facility.
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