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Exploring the mechanism of a pure and amorphous black-blue TiO2:H thin film as a photoanode in water splitting
Author’s Accepted Manuscript Exploring the Mechanism of a Pure and Amorphous Black-Blue TiO 2:H Thin Film as a Photoanode in Water Splitting Junhui Liang, Ning Wang, Qixing Zhang, Bofei Liu, Xiangbin Kong, Changchun Wei, Dekun Zhang, Baojie Yan, Ying Zhao, Xiaodan Zhang www.elsevier.com/locate/nanoenergy
PII: DOI: Reference:
S2211-2855(17)30667-5 https://doi.org/10.1016/j.nanoen.2017.10.062 NANOEN2296
To appear in: Nano Energy Received date: 15 August 2017 Revised date: 30 September 2017 Accepted date: 27 October 2017 Cite this article as: Junhui Liang, Ning Wang, Qixing Zhang, Bofei Liu, Xiangbin Kong, Changchun Wei, Dekun Zhang, Baojie Yan, Ying Zhao and Xiaodan Zhang, Exploring the Mechanism of a Pure and Amorphous Black-Blue TiO 2:H Thin Film as a Photoanode in Water Splitting, Nano Energy, https://doi.org/10.1016/j.nanoen.2017.10.062 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Exploring the Mechanism of a Pure and Amorphous Black-Blue TiO2:H Thin Film as a Photoanode in Water Splitting Junhui Lianga,b,c,d, Ning Wanga,b,c,d, Qixing Zhanga,b,c,d, Bofei Liua,b,c,d, Xiangbin Konga,b,c,d, Changchun Weia,b, Dekun Zhanga,b, Baojie Yana,b, Ying Zhaoa,b,c,d, Xiaodan Zhanga,b,c,d,* a
Institute of Photoelectronic Thin Film Devices and Technology of Nankai University, Tianjin 300071, P. R.
China b
Key Laboratory of Photoelectronic Thin Film Devices and Technology of Tianjin, Tianjin 300071, P. R. China
c
Key Laboratory of Optical Information Science and Technology of Ministry of Education, Tianjin 300071, P. R.
China d
Collaborative Innovation Center of Chemical Science and Engineering (Tianjin), Tianjin 300071, P. R. China
*
Corresponding author: Tel.: +86-22-23499304; fax: +86 22-23499304; E-mail address:
[email protected]
Abstract The use of amorphous disordered surface layers on TiO2 nanocrystals prepared by hydrogenated post-treatment has been proven to be an effective approach for enhancing the light absorption and photocatalytic activity of TiO2 photoanodes. However, the mechanism underlying the enhanced optical-electrical performance caused by the hydrogenated amorphous shell has still not been fully unveiled. Herein, to explore the mechanism without the effect of the crystalline structure, a pure, amorphous hydrogen-doped TiO2 (a-TiO2:H) film was prepared using a magnetron sputtering technique under reactive hydrogen plasma. We propose that the cooperative effects of the extended mid-gap states and valence band tail are responsible for the enhanced visible and near-infrared optical absorption. In addition, the doped H acts as a shallow donor to provide carriers and shift the Fermi level to the conduction band, ultimately accelerating charge transport and transfer at the semiconductor/electrolyte interface. Finally, the photoconversion efficiency of a-TiO2:H was improved one order of magnitude compared to undoped a-TiO2:H. Through the investigation of a-TiO2:H, we gained 1
further insight into black TiO2. In addition, we believe our a-TiO2:H film grown at room temperature opens new opportunities for a broad range of applications, including PEC water splitting, supercapacitors, dye-sensitized solar cells, and perovskite solar cells.
KEYWORDS: Amorphous TiO2:H; Photoanode; Optical absorption theoretical model; Extended mid-gap states; Band tail states
1. Introduction Solar-driven water splitting into H2 and O2 is considered as a promising approach to overcome the energy crisis and greenhouse effect caused by the gradually exhausted fossil fuels. Since Fujishima and Honda discovered photocatalytic water splitting on a titanium dioxide (TiO2) electrode in the 1970s, TiO2 has been extensively investigated as a photoanode for photoelectrochemical (PEC) water splitting owing to its favorable band-edge positions, superior chemical stability and low cost.[1-4] However, the solar conversion efficiency is seriously limited on TiO2 because of its large band gap ranging from 3.2 to 3.7 eV,[5-7] only making TiO2 effective under ultraviolet (UV) light, which accounts for less than 5% of total solar irradiation. For practical applications, two distinct approaches can be employed to improve the photocatalytic activity: (1) improving the conductivity and suppressing the recombination centers to prolong the lifetime of photogenerated carriers and (2) band gap engineering to enhance light absorption. Very recently, an amorphous shell/crystalline core structure of black TiO2 prepared by hydrogenation post-treatment of TiO2 nanocrystals demonstrated high visible and nearinfrared (VIS-NIR) optical absorption and excellent conductivity and thus triggered a surge of interest.[2, 3, 8-19] Theoretical models, including extended valence band tail states induced by a disordered surface amorphous shell,[9, 15] mid-gap electronic states caused by oxygen vacancies (Vo)[12, 14, 20-22] and H-related deficiencies,[11-13, 23-25] and localized surface plasmon resonance (LSPR) resulting from a high carrier concentration in the amorphous shell,[11, 26-28] have been employed to explore the reasons for the black color. However, the above models and corresponding experiments have not entirely uncovered the reason for the color change. Notably, the hydrogenated amorphous shell plays an important role in the color 2
change and enhancement of solar-driven photocatalytic activity. Unfortunately, a deep understanding of the mechanism of the enhanced optical absorption and conductivity of the hydrogenated amorphous shell without the effects of crystalline structures is still lacking. Thus, more careful investigations on the optical-electrical properties of pure and amorphous hydrogenated TiO2 are necessary to further understand this mechanism and open new applications for amorphous TiO2.[29] To gain insight into the effects of hydrogen doping (H-doping) on the optical-electrical properties of the amorphous shell, we employ magnetron sputtering technology to fabricate the pure and amorphous black-blue H-doped TiO2 (a-TiO2:H) under Ar and H2 atmosphere without any post-treatment process. We believe that both the valence band tail and mid-gap states caused by H-induced deficiencies lead to enhanced light harvesting. In addition, magnetron sputtering is a versatile technique, allowing us to change the H-doping flow rate more freely and precisely. As a result, we can control the passivating and etching effect of hydrogen by accurately controlling the H-doping flow rate, thus reducing recombination centers. We propose a novel theoretical model to explain the enhanced VIS-NIR light absorption, providing further insight into the effect of H-doping on the electrical properties of a-TiO2:H.
1. Experimental Section 2.1 Preparation of undoped a-TiO2 and a-TiO2:H All the samples were prepared on fluorine-doped tin oxide (FTO) glass substrates by pulsed direct current magnetron sputtering from an intrinsic TiO2 target (99.999%). The undoped a-TiO2 thin films were fabricated with a constant Ar gas flow (55 Sccm). The aTiO2:H thin films were prepared by introducing H2 (10 Sccm) and Ar (45 Sccm) gas mixture into the chamber during the preparation process. In addition, for all the samples, the distance between the target and substrate, sputtering power and working pressure were kept at 50 mm, 380 W and 4.0 mTorr, respectively. The thickness for all samples was kept at 280 nm. 2.2 Photoelectrochemical measurements The photoelectrochemical properties for all the samples were analyzed in a threeelectrode configuration immersed in a 1 M NaOH aqueous solution (pH = 13.6) under 1-sun 3
conditions (AM 1.5G, 100 mWcm-2) using a dual-beam solar simulator (Wacom WXs-156s12). The three-electrode configuration consisted of a Pt counter electrode, Ag/AgCl reference electrode and the working electrode with a well-defined area of 0.283 cm2. The potential of the working electrode was controlled by a Princeton potentiostat (PARSTAT 4000).
2. Results and discussions
Figure 1. Properties of undoped a-TiO2 and a-TiO2:H. a) Absorption spectra and the inset sample photos. b) XPS survey spectra. c) TEM images. d) Spectral dependence of the logarithm of the absorption coefficient on the photon energy.
Fig. 1a displays the UV-visible spectra of the undoped a-TiO2 and a-TiO2:H films fabricated by pulsed direct current magnetron sputtering. Evidently, the color of the sample changes from white (undoped a-TiO2) to black-blue (a-TiO2:H). The X-ray photoelectron spectroscopy (XPS) survey spectra (Fig. 1b) display no characteristic peaks for Ti or Sn metals. Thus, the enhanced light absorption is mainly ascribed to H-doping. Moreover, both the undoped a-TiO2 and a-TiO2:H films possess large absorption at wavelength shorter than 350 nm, which can be attributed to the intrinsic band gap absorption. The optical band gaps of 4
undoped a-TiO2 and a-TiO2:H can be calculated using the Tauc model (3.64 eV and 3.78 eV, respectively, as shown in Fig. S1), which are higher than the values of crystalline anatase and rutile TiO2 (3.2 and 3.0 eV, respectively).[30] These higher-energy band gaps are induced by the amorphous effect and quantum size effect.[31, 32] The amorphous structure is confirmed by the X-ray diffraction (XRD) patterns (Fig. S2) and transmission electron microscopy (TEM) images (Fig. 1c). A greatly enhanced light absorption peak of the a-TiO2:H film from approximately 600 nm to 930 nm is seen when compared with undoped a-TiO2, which is different from the previous research of the enhanced light harvesting for core-shell structure. For instance, the extended band tail model proposed by Chen et al. presents a gradual decrease of the optical harvesting with the wavelength increase due to the approach exponential distribution of band tail density.[9] The LSPR absorption model proposed by Wang et al. illustrates the optical harvesting continuously enhances with the wavelength increase.[11] As our result present a light absorption peak, the existing theory models couldn’t explain this phenomenon accurately. In order to investigate the cause of the enhanced light absorption, the spectral dependence of the logarithm of the absorption coefficient on the photon energy are presented (Fig. 1d). For amorphous materials, the absorption spectrum is divided into three regions: a strong absorption region (SAR), exponential absorption region (EAR) and weak absorption region (WAR). Among these regions, the SAR is related to the band-to-band transition absorption. The EAR or Urbach absorption edge, is associated with the transition absorption from the extended valence band tail to localized conduction band tail states. The WAR originates from transition absorption of mid-gap states. Compared with undoped a-TiO2, the absorption coefficient (α) in the WAR improves over one order of magnitude after H-doping. The high absorption coefficient of a-TiO2:H in the WAR indicates the existence of extended mid-gap states. In contrast, the absorption coefficient and light absorption of undoped a-TiO2 in the VIS-NIR region caused by oxygen vacancies and unintentional H-doping are much lower than those of a-TiO2:H, thus the mid-gap states caused by unintentional H-doping in undoped a-TiO2 are negligible.[33] We believe that the enhanced light absorption in the VISNIR region (blue, dashed region in Fig. 1a) is attributed to the extended mid-gap states absorption. In addition, the slope of the exponential absorption decreases after H-doping, indicating an enhancement of the localized state density of the band tail. Therefore, the enhanced light absorption in the short wavelength region (black, dashed region in Fig. 1a) for a-TiO2:H is mainly caused by the overlap of enhanced band tail state and mid-gap states. Our 5
mid-gap absorption model resulting from H-related deficiency differs with the oxygen vacancy absorption model proposed by Alberto Naldoni and Wang et al.[12, 20]. Distinguishing from the hydrogenated treatment TiO2, the doped H in a-TiO2:H prefers to passivate the oxygen vacancies and form Ti-H bonds rather than introducing oxygen vacancies, which will be discussed later. Thus, the enhanced optical harvesting in VIS-NIR region of a-TiO2:H is mainly attributed to the mid-gap caused by the H-related deficiency, and the oxygen vacancy only plays a weak role.
Figure 2. Properties of undoped a-TiO2 and a-TiO2:H. a) Ti 2p and O 1s XPS spectra. b) The experimental and fitting results of a-TiO2:H film. Ti 2p XPS spectra (upper). Normalized O 1s XPS spectra (lower). c) PL emission spectra. d) Proposed electronic energy band diagrams.
Ti 2p and O 1s XPS spectra (shown in Fig. 2a) were employed to study the origins of the extended band tail states and mid-gap states for a-TiO2:H, as the introduced H-bonding has a great influence on the electronic energy band structure.[3, 23] For both samples, the Ti 2p and O 1s XPS spectra are very similar with Ti 2p3/2, 2p1/2 and O 1s peaks centered at binding energies of 458.9 eV, 464.7 eV and 530.4 eV, respectively, which are typical for the Ti4+-O bonds in TiO2.[34] Notably, an additional broad peak was found at 457.4 eV and fitted from the Ti 2p spectra for a-TiO2:H (Fig. 2b upper), corresponding to the Ti-H bonds (H substitutional oxygen, Ho, or passivated dangling bonds).[35] Compared with undoped a-TiO2, 6
the a-TiO2:H film exhibits a slightly broader O 1s peak, which is deconvoluted into two peaks centered at 530.4 and 531.6 eV (Fig. 2b lower). The 531.6 eV peak is attributed to Ti-H-O bonds (interstitial hydrogen, Hi), in which the doped H is located at the lattice interstitial site and bound with oxygen.[34, 36] The ratios of Ho and Hi calculated from the peak areas in the Ti 2p and O 1s spectra for undoped a-TiO2 and a-TiO2:H are shown in Table S1. Ho is the key species to form mid-gap states and responsible for the enhanced VIS-NIR light absorption (blue, dashed region in Fig. 1a).[11, 12] The introduced Hi can enhance the lattice disorder and localized state density, and the energy levels of Hi have been reported to be located at the valence band.[37] Thus, Hi plays an important role in enhancing the light absorption in the black, dashed region in Fig. 1a. The photoluminescence (PL) emission spectra of undoped a-TiO2 and a-TiO2:H in the wavelength range of 360-600 nm with an excitation of 310 nm are shown in Fig. 2c to explore the mid-gap states in the electronic energy band gaps. The emission peaks predominantly at 396 nm (3.13 eV) and 391 nm (3.17 eV) for undoped a-TiO2 and a-TiO2:H, respectively, are attributed to the band gap transitions, which are smaller than the optical band gaps calculated by the Tauc model. The lower-energy transitions are attributed to the band tail states merging with the valence band and the conduction band caused by lattice distortion of the amorphous structure.[23] The emission peak at approximately 390 nm originates from the band gap transition from the valence band tail to the conduction band tail. The conduction band tail is neither sensitive to local lattice distortions nor sensitive to changes induced by bonding with H. Thus, the localized states of the conduction band tail are located at 0.12 eV below the conduction band maximum for both undoped a-TiO2 and a-TiO2:H.[23] The valence band tails are located at 0.39 eV and 0.49 eV above the valence band maximum for undoped a-TiO2 and a-TiO2:H, respectively. The larger valence band tail of a-TiO2:H than undoped a-TiO2 is mainly due to Hi, which results in a smaller slope for the exponential absorption and enhanced band tail state density. A weak emission peak located at 466 nm (2.66 eV) was found for a-TiO2:H, resulting from the emission of transitions from the mid-gap states to the conduction band. Thus, the mid-gap state energy level is located at 1.12 eV higher than the valence band maximum, which is well consistent with calculations (1.14 eV) based on density functional theory (DFT) 7
using the Perdew-Burke-Ernzerhof (PBE) function.[23] In addition, the onset of the optical absorption of a-TiO2:H is lowered to approximately 0.48 eV (2600 nm, shown in Fig. 1a), and an abrupt change in absorption at approximately 932 nm (1.33 eV) suggests a narrowed band gap resulting from intraband transitions.[9] Clearly, H-bonding (Hi and Ho) can not only extend the valence band tail, leaving the conduction band unchanged, but also induce extended mid-gap states. Based on the discussion above, the electronic energy band diagrams of undoped a-TiO2 and a-TiO2:H are schematically shown in Fig. 2d. The enhanced light harvesting ability of a-TiO2:H in VIS-NIR region is attributed to transitions from the valence band tail to the mid-gap states or from the mid-gap states to the conduction band. Table 1. Electrical properties of the undoped a-TiO2 and a-TiO2:H photoanodes. The carrier concentrations and flat-band potentials were calculated using the Mott-Schottky plots (Figure 3a). Donor density
Flat-band potential
(1019/cm3)
(V vs NHE)
Undoped a-TiO2
5.82
-0.68
a-TiO2:H
873.29
-1.09
Samples
Figure 3. PEC performance of undoped a-TiO2 and a-TiO2:H deposited on FTO glass 8
substrates. a) Mott-Schottky plots measured at 5k Hz. b) Band energetics of the semiconductor/electrolyte interface after equilibration for undoped a-TiO2 and a-TiO2:H. c) Nyquist plots from EIS measured at 1.23 V vs. RHE. The inset shows the equivalent circuit model used to fit the Nyquist plots.
As discussed above, doped H mainly exists in the forms of Ho and Hi.[38, 39] Both Ho and Hi can act as shallow donors to improve the electrical conductivity. Mott-Schottky and electrochemical impedance spectroscopy (EIS) measurements are employed to characterize the electrical properties. Both samples show positive slopes in the Mott-Schottky plots (Fig. 3a), as expected for n-type semiconductors. In addition, a-TiO2:H shows a substantially smaller slope than undoped a-TiO2, suggesting an increased donor density. To more intuitively analyze the donor density, we calculate the donor density (Nd) and flat-band potential (Vfb) for both the samples from the slopes of the Mott-Schottky plots using the following equations: 2 2 d 1/ C 1 2 kT and Nd (V Vfb ) 2 2 C 0 A eN d e e0 0 dV 1
(1)
where e0 is the electron charge, 𝜀 the dielectric constant of TiO2 ( 𝜀 =73),[40] 𝜀 0 is the permittivity of vacuum, and V is the applied bias at the electrode. C and A are the interfacial capacitance and area, respectively, k is Boltzmann’s constant, T is the absolute temperature, and e is the electronic charge. Therefore, a plot of 1/C2 against V should yield a straight line from which Vfb can be determined from the intercept on the V axis. The calculated values of the donor density and Vfb are shown in Table 1. Undoped aTiO2 demonstrates the properties of an n-type semiconductor, whose electrons originate from intrinsic defects, such as oxygen vacancies and unintentional H-doping. An improvement of two orders of magnitude in the donor density is achieved after H-doping. H-induced defects in a-TiO2:H mainly consist of two forms, Hi and Ho.
2TiO2 3H 2 2Ti H O( Hi ) 2H 2O 2TiO2 5H 2 2Ti H ( HO ) 4H 2O
(2) (3)
Among of these forms, Hi and Ho are located at the shallow donor energy levels, favorable to provide donors.[38] Thus, a-TiO2:H shows a higher donor density than undoped a-TiO2. 9
Additional potential advantages of these H-bonding deficiencies involve the deficiencies passivating the recombination centers (Vo and dangling bonds), providing trap sites for photogenerated carriers and preventing these trap sites from rapid recombination, thus promoting electron transfer and photocatalytic reactions. In addition, the Fermi energy level is a function of the carrier density, and the relationship between the Fermi energy level and Vfb has been discussed in our previous report.[41] A more negative flat-band potential of a-TiO2:H indicates larger band bending, which will accelerate charge separation and transfer at the semiconductor/electrolyte interface (as shown in Fig. 3b). In addition, a more negative flat-band potential can also block electrons and suppress charge recombination at the interface. The potential vs. Ag/AgCl can be converted into the potential vs. the normal hydrogen electrode (NHE) as follows: 0 ENHE E Ag / AgCl EAg / AgCl ( 0.197 V)
(4)
Table 2. Electrical properties and catalytic activities of the undoped a-TiO2 and a-TiO2:H photoanodes. The fitting results calculated by the Nyquist plots (Figure 3c). Rs
Rtransport
Rtransfer
(𝛀 cm2)
(𝛀 cm2)
(𝛀 cm2)
Undoped a-TiO2
26.2
1590
39590
a-TiO2:H
26.3
473.6
3526
Samples
The electrical performance for charge transport and transfer can be calculated by fitting the Nyquist plots of undoped a-TiO2 and a-TiO2:H (Fig. 3c), which are deduced from the EIS measurements under illumination. A smaller diameter is found for the a-TiO2:H film compared with undoped a-TiO2, suggesting better PEC performance. To more intuitively explain this behavior, Table 2 shows the external circuit resistance (Rs), charge-transport resistance (Rtransport) and charge-transfer resistance (Rtransfer) deduced from the equivalent circuit fitting (the inset image in Fig. 3c). Rs represents the charge-transport capacity in the external circuit and electrolyte. Rtransport corresponds to the charge-transport capacity in the bulk of TiO2. Rtransfer represents the charge-transfer resistance at the semiconductor/electrolyte interface. From Table 2, Rs is approximately the same for the two samples, suggesting a 10
stable testing environment. Both Rtransport and Rtransfer of a-TiO2:H decreases greatly, indicating faster charge transport and transfer. We believe that this behavior is attributed to the improved donor density and larger band bending.
Figure 4. PEC performances of undoped a-TiO2 and a-TiO2:H deposited on FTO glass substrates. a) Three-electrode voltammograms curves. b) Photoconversion efficiencies at different applied voltages.
Fig. 4a demonstrates the three-electrode voltammograms of undoped a-TiO2 and aTiO2:H. Compared with undoped a-TiO2, a considerable improvement in the saturation current density is obtained for a-TiO2:H. The saturation photocurrent density increases from 0.13 mA/cm2 to 0.67 mA/cm2 at 1.23 V vs. the reversible hydrogen electrode (RHE) for aTiO2:H compared to undoped a-TiO2. The potential vs. Ag/AgCl can be converted into the potential vs. RHE as follows: 0 ERHE E Ag / AgCl EAg / AgCl ( 0.197 V) 0.0591 pH
(5)
The higher saturation current density of a-TiO2:H is mainly attributed to three issues: higher light absorption, smaller charge-transport and charge-transfer resistance, and less recombination centers. Meanwhile, a-TiO2:H can achieve saturation current at a more negative applied voltage, indicating better catalytic performance. The photoconversion efficiency can be calculated as follows:[42]
J mp 1.23 Vmp Pin
(6)
where Jmp and Vmp are the current density and voltage at the maximum power point (MPP), respectively, and Pin is the power of the incoming illumination. The photoconversion 11
efficiencies with different applied voltages of undoped a-TiO2 and a-TiO2:H are presented in Fig. 4b. An order of magnitude improvement in the photoconversion efficiency (from 0.02% to 0.26%) for a-TiO2:H is achieved compared with undoped a-TiO2. Compared with most hydrogenated crystalline TiO2 materials (Table S2), our a-TiO2:H film obtains a high photocurrent density of 674 μA/cm2 at an applied voltage of 1.23 V vs. RHE. This high photocurrent density is attributed to the high light absorption, reduced recombination centers, and more shallow donors (Hi and Ho) to provide a higher carrier concentration and accelerate charge transport and transfer.
4. Conclusions In summary, we successfully prepared a-TiO2:H at room temperature with simple magnetron sputtering technology. We reported an efficient theoretical model to explain the cause of the color change of a-TiO2:H. We believed that the introduced extended mid-gap states and band tail merging with the valence band causing by H-bonding was responsible for the enhanced VIS-NIR light absorption. In addition, a suitable H-doping flow rate passivated the recombination centers to reduce carrier recombination and introduced shallow donors (Ho and Hi) to provide free carriers, thus improving the carrier transport and PEC performance. Finally, based on a planar structure, we obtained a photocurrent density of 674 𝜇A/cm2 for aTiO2:H at an applied voltage of 1.23 V vs. RHE, which is comparable with most crystalline TiO2 materials, even with nanostructures. We gained insight into the mechanism to understand the cause of the color change for hydrogenated TiO2. We believe the ability to fabricate highly photoactive a-TiO2:H films will open new opportunities in various areas.
Acknowledgements The authors gratefully acknowledge the financial supports from the International cooperation project of the Ministry of Science and Technology (2014DFE60170), National Natural Science Foundation of China (Grant No. 61474065 and 61674084), Tianjin Research Key Program of Application Foundation and Advanced Technology (15JCZDJC31300), Key Project in the Science & Technology Pillar Program of Jiangsu Province (BE2014147-3), and the 111 Project (B16027). 12
Appendix A. Supplementary material Supplementary data associated with this article can be found in the online version at
References [1] A. Fujishima, K. Honda, Nature 238 (1972) 37-38. [2] Y.H. Hu, Angew. Chem., Int. Ed. 51 (2012) 12410-12412. [3] X. Chen, L. Liu, F. Huang, Chem. Soc. Rev. 44 (2015) 1861-1885. [4] X. Chen, S.S. Mao, Chem. Rev. 107 (2007) 2891-2959. [5] Z. Wang, U. Helmersson, P. O. Käll, Thin Solid Films 405 (2002) 50-54. [6] A. Di Paola, M. Bellardita, L. Palmisano, Brookite, Catalysts 3 (2013) 36-73. [7] A. Kruth, S. Peglow, A. Quade, M. M. Pohl, R.d. Foest, V. Brüser, K. D. Weltmann, J. Phys. Chem. C 118 (2014) 25234-25244. [8] L. Mascaretti, S. Ferrulli, P. Mazzolini, C.S. Casari, V. Russo, R. Matarrese, I. Nova, G. Terraneo, N. Liu, P. Schmuki, A. Li Bassi, Sol. Energy Mater. Sol. Cells 169 (2017) 1927. [9] X.B. Chen, L. Liu, P.Y. Yu, S.S. Mao, Science 331 (2011) 746-750. [10] C. Fan, C. Chen, J. Wang, X. Fu, Z. Ren, G. Qian, Z. Wang, J. Mater. Chem. A 2 (2014) 16242-16249. [11] Z. Wang, C. Yang, T. Lin, H. Yin, P. Chen, D. Wan, F. Xu, F. Huang, J. Lin, X. Xie, M. Jiang, Adv. Funct. Mater. 23 (2013) 5444-5450. [12] G.M. Wang, H.Y. Wang, Y.C. Ling, Y.C. Tang, X.Y. Yang, R.C. Fitzmorris, C.C. Wang, J.Z. Zhang, Y. Li, Nano Lett. 11 (2011) 3026-3033. [13] M. Mehta, N. Kodan, S. Kumar, A. Kaushal, L. Mayrhofer, M. Walter, M. Moseler, A. Dey, S. Krishnamurthy, S. Basu, J. Mater. Chem. A 4 (2016) 2670-2681. [14] T. Leshuk, R. Parviz, P. Everett, H. Krishnakumar, R.A. Varin, F. Gu, ACS Appl. Mater. Interfaces 5 (2013) 1892-1895. [15] X. Chen, L. Liu, Z. Liu, M.A. Marcus, W.-C. Wang, N.A. Oyler, M.E. Grass, B. Mao, P.A. Glans, Y.Y. Peter, Sci. Rep. 3 (2013) 1510. [16] W. Zhou, W. Li, J. Q. Wang, Y. Qu, Y. Yang, Y. Xie, K. Zhang, L. Wang, H. Fu, D. Zhao, 13
J. Am. Chem. Soc. 136 (2014) 9280-9283. [17] K. Zhang, J.H. Park, J. Phys. Chem. Lett. 8 (2017) 199-207. [18] C.Y. Yang, Z. Wang, T.Q. Lin, H. Yin, X.J. Lu, D.Y. Wan, T. Xu, C. Zheng, J.H. Lin, F.Q. Huang, X.M. Xie, M.H. Jiangl, J. Am. Chem. Soc. 135 (2013) 17831-17838. [19] X. Lu, G. Wang, T. Zhai, M. Yu, J. Gan, Y. Tong, Y. Li, Nano Lett. 12 (2012) 1690-1696. [20] A. Naldoni, M. Allieta, S. Santangelo, M. Marelli, F. Fabbri, S. Cappelli, C.L. Bianchi, R. Psaro, V. Dal Santo, J. Am. Chem. Soc. 134 (2012) 7600-7603. [21] M. Choi, J.H. Lee, Y.J. Jang, D. Kim, J.S. Lee, H.M. Jang, K. Yong, Sci. Rep. 6 (2016) 36099. [22] S. Wendt, P.T. Sprunger, E. Lira, G.K. Madsen, Z. Li, J.Ø. Hansen, J. Matthiesen, A. Blekinge-Rasmussen, E. Lægsgaard, B. Hammer, Science 320 (2008) 1755-1759. [23] L. Liu, P.Y. Yu, X. Chen, S.S. Mao, D.Z. Shen, Phys. Rev. Lett. 111 (2013) 065505. [24] H.H. Pham, L. W. Wang, Phys. Chem. Chem. Phys. 17 (2015) 11908-11913. [25] S. Pan, X. Liu, M. Guo, S.f. Yu, H. Huang, H. Fan, G. Li, J. Mater. Chem. A 3 (2015) 11437-11443. [26] F. Zhang, Y. Zheng, Y. Cao, C. Chen, Y. Zhan, X. Lin, Q. Zheng, K. Wei, J. Zhu, J. Mater. Chem. 19 (2009) 2771-2777. [27] M. Hari, S.A. Joseph, S. Mathew, P. Radhakrishnan, V. Nampoori, J. Appl. Phys. 112 (2012) 074307. [28] Z.W. Seh, S. Liu, M. Low, S.Y. Zhang, Z. Liu, A. Mlayah, M.Y. Han, Adv. Mater. 24 (2012) 2310-2314. [29] C. Fan, C. Chen, J. Wang, X. Fu, Z. Ren, G. Qian, Z. Wang, Sci. Rep. 5 (2015) 11712. [30] D.O. Scanlon, C.W. Dunnill, J. Buckeridge, S.A. Shevlin, A.J. Logsdail, S.M. Woodley, C.R.A. Catlow, M.J. Powell, R.G. Palgrave, I.P. Parkin, G.W. Watson, T.W. Keal, P. Sherwood, A. Walsh, A.A. Sokol, Nat. Mater. 12 (2013) 798-801. [31] K. Eufinger, D. Poelman, H. Poelman, R. De Gryse, G.B. Marin, Appl. Surf. Sci. 254 (2007) 148-152. [32] D.M. King, X. Du, A.S. Cavanagh, A.W. Weimer, Nanotechnology 19 (2008) 445401. [33] D.M. Hofmann, A. Hofstaetter, F. Leiter, H. Zhou, F. Henecker, B.K. Meyer, S.B. Orlinskii, J. Schmidt, P.G. Baranov, Phys. Rev. Lett. 88 (2002) 045504. 14
[34] M. Lazarus, T. Sham, Chem. Phys. Lett. 92 (1982) 670-674. [35] Z. Zheng, B. Huang, J. Lu, Z. Wang, X. Qin, X. Zhang, Y. Dai, M. H. Whangbo, Chem. Commun. 48 (2012) 5733-5735. [36] E. McCafferty, J. Wightman, Surf. Interface Anal. 26 (1998) 549-564. [37] A. Imanishi, T. Okamura, N. Ohashi, R. Nakamura, Y. Nakato, J. Am. Chem. Soc. 129 (2007) 11569-11578. [38] A. Janotti, C.G. Van de Walle, Nat. Mater. 6 (2007) 44-47. [39] C. Sun, Y. Jia, X. H. Yang, H. G. Yang, X. Yao, G.Q. Lu, A. Selloni, S.C. Smith, J. Phys. Chem. C 115 (2011) 25590-25594. [40] M.D. Stamate, Appl. Surf. Sci. 218 (2003) 318-323. [41] J. Liang, H. Tan, M. Liu, B. Liu, N. Wang, Q. Zhang, Y. Zhao, A.H.M. Smets, M. Zeman, X. Zhang, J. Mater. Chem. A 4 (2016) 16841-16848. [42] M.G. Walter, E.L. Warren, J.R. McKone, S.W. Boettcher, Q. Mi, E.A. Santori, N.S. Lewis, Chem. Rev. 110 (2010) 6446-6473.
Graphical Abstract
A theoretical model, the cooperative effects of the extended mid-gap states and valence band tail caused by H-induced deficiencies, is proposed to explore the underlying cause for enhanced light harvesting of H-doped black-blue amorphous TiO2.
Highlights Magnetron sputtering technique has been employed to fabricate the pure amorphous black-blue TiO2:H. 15
Suitable H-doping flow rate can passivate the recombination centers to reduce carrier recombination and improve conductivity. A theoretical model is proposed to explore the underlying cause for enhanced light harvesting of H-doping black-blue TiO2. Based on a planar structure, we obtain a photocurrent density of 674 𝜇A/cm2 for a-TiO2:H at an applied voltage of 1.23 V vs. RHE
16
PII: DOI: Reference:
S2211-2855(17)30667-5 https://doi.org/10.1016/j.nanoen.2017.10.062 NANOEN2296
To appear in: Nano Energy Received date: 15 August 2017 Revised date: 30 September 2017 Accepted date: 27 October 2017 Cite this article as: Junhui Liang, Ning Wang, Qixing Zhang, Bofei Liu, Xiangbin Kong, Changchun Wei, Dekun Zhang, Baojie Yan, Ying Zhao and Xiaodan Zhang, Exploring the Mechanism of a Pure and Amorphous Black-Blue TiO 2:H Thin Film as a Photoanode in Water Splitting, Nano Energy, https://doi.org/10.1016/j.nanoen.2017.10.062 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Exploring the Mechanism of a Pure and Amorphous Black-Blue TiO2:H Thin Film as a Photoanode in Water Splitting Junhui Lianga,b,c,d, Ning Wanga,b,c,d, Qixing Zhanga,b,c,d, Bofei Liua,b,c,d, Xiangbin Konga,b,c,d, Changchun Weia,b, Dekun Zhanga,b, Baojie Yana,b, Ying Zhaoa,b,c,d, Xiaodan Zhanga,b,c,d,* a
Institute of Photoelectronic Thin Film Devices and Technology of Nankai University, Tianjin 300071, P. R.
China b
Key Laboratory of Photoelectronic Thin Film Devices and Technology of Tianjin, Tianjin 300071, P. R. China
c
Key Laboratory of Optical Information Science and Technology of Ministry of Education, Tianjin 300071, P. R.
China d
Collaborative Innovation Center of Chemical Science and Engineering (Tianjin), Tianjin 300071, P. R. China
*
Corresponding author: Tel.: +86-22-23499304; fax: +86 22-23499304; E-mail address:
[email protected]
Abstract The use of amorphous disordered surface layers on TiO2 nanocrystals prepared by hydrogenated post-treatment has been proven to be an effective approach for enhancing the light absorption and photocatalytic activity of TiO2 photoanodes. However, the mechanism underlying the enhanced optical-electrical performance caused by the hydrogenated amorphous shell has still not been fully unveiled. Herein, to explore the mechanism without the effect of the crystalline structure, a pure, amorphous hydrogen-doped TiO2 (a-TiO2:H) film was prepared using a magnetron sputtering technique under reactive hydrogen plasma. We propose that the cooperative effects of the extended mid-gap states and valence band tail are responsible for the enhanced visible and near-infrared optical absorption. In addition, the doped H acts as a shallow donor to provide carriers and shift the Fermi level to the conduction band, ultimately accelerating charge transport and transfer at the semiconductor/electrolyte interface. Finally, the photoconversion efficiency of a-TiO2:H was improved one order of magnitude compared to undoped a-TiO2:H. Through the investigation of a-TiO2:H, we gained 1
further insight into black TiO2. In addition, we believe our a-TiO2:H film grown at room temperature opens new opportunities for a broad range of applications, including PEC water splitting, supercapacitors, dye-sensitized solar cells, and perovskite solar cells.
KEYWORDS: Amorphous TiO2:H; Photoanode; Optical absorption theoretical model; Extended mid-gap states; Band tail states
1. Introduction Solar-driven water splitting into H2 and O2 is considered as a promising approach to overcome the energy crisis and greenhouse effect caused by the gradually exhausted fossil fuels. Since Fujishima and Honda discovered photocatalytic water splitting on a titanium dioxide (TiO2) electrode in the 1970s, TiO2 has been extensively investigated as a photoanode for photoelectrochemical (PEC) water splitting owing to its favorable band-edge positions, superior chemical stability and low cost.[1-4] However, the solar conversion efficiency is seriously limited on TiO2 because of its large band gap ranging from 3.2 to 3.7 eV,[5-7] only making TiO2 effective under ultraviolet (UV) light, which accounts for less than 5% of total solar irradiation. For practical applications, two distinct approaches can be employed to improve the photocatalytic activity: (1) improving the conductivity and suppressing the recombination centers to prolong the lifetime of photogenerated carriers and (2) band gap engineering to enhance light absorption. Very recently, an amorphous shell/crystalline core structure of black TiO2 prepared by hydrogenation post-treatment of TiO2 nanocrystals demonstrated high visible and nearinfrared (VIS-NIR) optical absorption and excellent conductivity and thus triggered a surge of interest.[2, 3, 8-19] Theoretical models, including extended valence band tail states induced by a disordered surface amorphous shell,[9, 15] mid-gap electronic states caused by oxygen vacancies (Vo)[12, 14, 20-22] and H-related deficiencies,[11-13, 23-25] and localized surface plasmon resonance (LSPR) resulting from a high carrier concentration in the amorphous shell,[11, 26-28] have been employed to explore the reasons for the black color. However, the above models and corresponding experiments have not entirely uncovered the reason for the color change. Notably, the hydrogenated amorphous shell plays an important role in the color 2
change and enhancement of solar-driven photocatalytic activity. Unfortunately, a deep understanding of the mechanism of the enhanced optical absorption and conductivity of the hydrogenated amorphous shell without the effects of crystalline structures is still lacking. Thus, more careful investigations on the optical-electrical properties of pure and amorphous hydrogenated TiO2 are necessary to further understand this mechanism and open new applications for amorphous TiO2.[29] To gain insight into the effects of hydrogen doping (H-doping) on the optical-electrical properties of the amorphous shell, we employ magnetron sputtering technology to fabricate the pure and amorphous black-blue H-doped TiO2 (a-TiO2:H) under Ar and H2 atmosphere without any post-treatment process. We believe that both the valence band tail and mid-gap states caused by H-induced deficiencies lead to enhanced light harvesting. In addition, magnetron sputtering is a versatile technique, allowing us to change the H-doping flow rate more freely and precisely. As a result, we can control the passivating and etching effect of hydrogen by accurately controlling the H-doping flow rate, thus reducing recombination centers. We propose a novel theoretical model to explain the enhanced VIS-NIR light absorption, providing further insight into the effect of H-doping on the electrical properties of a-TiO2:H.
1. Experimental Section 2.1 Preparation of undoped a-TiO2 and a-TiO2:H All the samples were prepared on fluorine-doped tin oxide (FTO) glass substrates by pulsed direct current magnetron sputtering from an intrinsic TiO2 target (99.999%). The undoped a-TiO2 thin films were fabricated with a constant Ar gas flow (55 Sccm). The aTiO2:H thin films were prepared by introducing H2 (10 Sccm) and Ar (45 Sccm) gas mixture into the chamber during the preparation process. In addition, for all the samples, the distance between the target and substrate, sputtering power and working pressure were kept at 50 mm, 380 W and 4.0 mTorr, respectively. The thickness for all samples was kept at 280 nm. 2.2 Photoelectrochemical measurements The photoelectrochemical properties for all the samples were analyzed in a threeelectrode configuration immersed in a 1 M NaOH aqueous solution (pH = 13.6) under 1-sun 3
conditions (AM 1.5G, 100 mWcm-2) using a dual-beam solar simulator (Wacom WXs-156s12). The three-electrode configuration consisted of a Pt counter electrode, Ag/AgCl reference electrode and the working electrode with a well-defined area of 0.283 cm2. The potential of the working electrode was controlled by a Princeton potentiostat (PARSTAT 4000).
2. Results and discussions
Figure 1. Properties of undoped a-TiO2 and a-TiO2:H. a) Absorption spectra and the inset sample photos. b) XPS survey spectra. c) TEM images. d) Spectral dependence of the logarithm of the absorption coefficient on the photon energy.
Fig. 1a displays the UV-visible spectra of the undoped a-TiO2 and a-TiO2:H films fabricated by pulsed direct current magnetron sputtering. Evidently, the color of the sample changes from white (undoped a-TiO2) to black-blue (a-TiO2:H). The X-ray photoelectron spectroscopy (XPS) survey spectra (Fig. 1b) display no characteristic peaks for Ti or Sn metals. Thus, the enhanced light absorption is mainly ascribed to H-doping. Moreover, both the undoped a-TiO2 and a-TiO2:H films possess large absorption at wavelength shorter than 350 nm, which can be attributed to the intrinsic band gap absorption. The optical band gaps of 4
undoped a-TiO2 and a-TiO2:H can be calculated using the Tauc model (3.64 eV and 3.78 eV, respectively, as shown in Fig. S1), which are higher than the values of crystalline anatase and rutile TiO2 (3.2 and 3.0 eV, respectively).[30] These higher-energy band gaps are induced by the amorphous effect and quantum size effect.[31, 32] The amorphous structure is confirmed by the X-ray diffraction (XRD) patterns (Fig. S2) and transmission electron microscopy (TEM) images (Fig. 1c). A greatly enhanced light absorption peak of the a-TiO2:H film from approximately 600 nm to 930 nm is seen when compared with undoped a-TiO2, which is different from the previous research of the enhanced light harvesting for core-shell structure. For instance, the extended band tail model proposed by Chen et al. presents a gradual decrease of the optical harvesting with the wavelength increase due to the approach exponential distribution of band tail density.[9] The LSPR absorption model proposed by Wang et al. illustrates the optical harvesting continuously enhances with the wavelength increase.[11] As our result present a light absorption peak, the existing theory models couldn’t explain this phenomenon accurately. In order to investigate the cause of the enhanced light absorption, the spectral dependence of the logarithm of the absorption coefficient on the photon energy are presented (Fig. 1d). For amorphous materials, the absorption spectrum is divided into three regions: a strong absorption region (SAR), exponential absorption region (EAR) and weak absorption region (WAR). Among these regions, the SAR is related to the band-to-band transition absorption. The EAR or Urbach absorption edge, is associated with the transition absorption from the extended valence band tail to localized conduction band tail states. The WAR originates from transition absorption of mid-gap states. Compared with undoped a-TiO2, the absorption coefficient (α) in the WAR improves over one order of magnitude after H-doping. The high absorption coefficient of a-TiO2:H in the WAR indicates the existence of extended mid-gap states. In contrast, the absorption coefficient and light absorption of undoped a-TiO2 in the VIS-NIR region caused by oxygen vacancies and unintentional H-doping are much lower than those of a-TiO2:H, thus the mid-gap states caused by unintentional H-doping in undoped a-TiO2 are negligible.[33] We believe that the enhanced light absorption in the VISNIR region (blue, dashed region in Fig. 1a) is attributed to the extended mid-gap states absorption. In addition, the slope of the exponential absorption decreases after H-doping, indicating an enhancement of the localized state density of the band tail. Therefore, the enhanced light absorption in the short wavelength region (black, dashed region in Fig. 1a) for a-TiO2:H is mainly caused by the overlap of enhanced band tail state and mid-gap states. Our 5
mid-gap absorption model resulting from H-related deficiency differs with the oxygen vacancy absorption model proposed by Alberto Naldoni and Wang et al.[12, 20]. Distinguishing from the hydrogenated treatment TiO2, the doped H in a-TiO2:H prefers to passivate the oxygen vacancies and form Ti-H bonds rather than introducing oxygen vacancies, which will be discussed later. Thus, the enhanced optical harvesting in VIS-NIR region of a-TiO2:H is mainly attributed to the mid-gap caused by the H-related deficiency, and the oxygen vacancy only plays a weak role.
Figure 2. Properties of undoped a-TiO2 and a-TiO2:H. a) Ti 2p and O 1s XPS spectra. b) The experimental and fitting results of a-TiO2:H film. Ti 2p XPS spectra (upper). Normalized O 1s XPS spectra (lower). c) PL emission spectra. d) Proposed electronic energy band diagrams.
Ti 2p and O 1s XPS spectra (shown in Fig. 2a) were employed to study the origins of the extended band tail states and mid-gap states for a-TiO2:H, as the introduced H-bonding has a great influence on the electronic energy band structure.[3, 23] For both samples, the Ti 2p and O 1s XPS spectra are very similar with Ti 2p3/2, 2p1/2 and O 1s peaks centered at binding energies of 458.9 eV, 464.7 eV and 530.4 eV, respectively, which are typical for the Ti4+-O bonds in TiO2.[34] Notably, an additional broad peak was found at 457.4 eV and fitted from the Ti 2p spectra for a-TiO2:H (Fig. 2b upper), corresponding to the Ti-H bonds (H substitutional oxygen, Ho, or passivated dangling bonds).[35] Compared with undoped a-TiO2, 6
the a-TiO2:H film exhibits a slightly broader O 1s peak, which is deconvoluted into two peaks centered at 530.4 and 531.6 eV (Fig. 2b lower). The 531.6 eV peak is attributed to Ti-H-O bonds (interstitial hydrogen, Hi), in which the doped H is located at the lattice interstitial site and bound with oxygen.[34, 36] The ratios of Ho and Hi calculated from the peak areas in the Ti 2p and O 1s spectra for undoped a-TiO2 and a-TiO2:H are shown in Table S1. Ho is the key species to form mid-gap states and responsible for the enhanced VIS-NIR light absorption (blue, dashed region in Fig. 1a).[11, 12] The introduced Hi can enhance the lattice disorder and localized state density, and the energy levels of Hi have been reported to be located at the valence band.[37] Thus, Hi plays an important role in enhancing the light absorption in the black, dashed region in Fig. 1a. The photoluminescence (PL) emission spectra of undoped a-TiO2 and a-TiO2:H in the wavelength range of 360-600 nm with an excitation of 310 nm are shown in Fig. 2c to explore the mid-gap states in the electronic energy band gaps. The emission peaks predominantly at 396 nm (3.13 eV) and 391 nm (3.17 eV) for undoped a-TiO2 and a-TiO2:H, respectively, are attributed to the band gap transitions, which are smaller than the optical band gaps calculated by the Tauc model. The lower-energy transitions are attributed to the band tail states merging with the valence band and the conduction band caused by lattice distortion of the amorphous structure.[23] The emission peak at approximately 390 nm originates from the band gap transition from the valence band tail to the conduction band tail. The conduction band tail is neither sensitive to local lattice distortions nor sensitive to changes induced by bonding with H. Thus, the localized states of the conduction band tail are located at 0.12 eV below the conduction band maximum for both undoped a-TiO2 and a-TiO2:H.[23] The valence band tails are located at 0.39 eV and 0.49 eV above the valence band maximum for undoped a-TiO2 and a-TiO2:H, respectively. The larger valence band tail of a-TiO2:H than undoped a-TiO2 is mainly due to Hi, which results in a smaller slope for the exponential absorption and enhanced band tail state density. A weak emission peak located at 466 nm (2.66 eV) was found for a-TiO2:H, resulting from the emission of transitions from the mid-gap states to the conduction band. Thus, the mid-gap state energy level is located at 1.12 eV higher than the valence band maximum, which is well consistent with calculations (1.14 eV) based on density functional theory (DFT) 7
using the Perdew-Burke-Ernzerhof (PBE) function.[23] In addition, the onset of the optical absorption of a-TiO2:H is lowered to approximately 0.48 eV (2600 nm, shown in Fig. 1a), and an abrupt change in absorption at approximately 932 nm (1.33 eV) suggests a narrowed band gap resulting from intraband transitions.[9] Clearly, H-bonding (Hi and Ho) can not only extend the valence band tail, leaving the conduction band unchanged, but also induce extended mid-gap states. Based on the discussion above, the electronic energy band diagrams of undoped a-TiO2 and a-TiO2:H are schematically shown in Fig. 2d. The enhanced light harvesting ability of a-TiO2:H in VIS-NIR region is attributed to transitions from the valence band tail to the mid-gap states or from the mid-gap states to the conduction band. Table 1. Electrical properties of the undoped a-TiO2 and a-TiO2:H photoanodes. The carrier concentrations and flat-band potentials were calculated using the Mott-Schottky plots (Figure 3a). Donor density
Flat-band potential
(1019/cm3)
(V vs NHE)
Undoped a-TiO2
5.82
-0.68
a-TiO2:H
873.29
-1.09
Samples
Figure 3. PEC performance of undoped a-TiO2 and a-TiO2:H deposited on FTO glass 8
substrates. a) Mott-Schottky plots measured at 5k Hz. b) Band energetics of the semiconductor/electrolyte interface after equilibration for undoped a-TiO2 and a-TiO2:H. c) Nyquist plots from EIS measured at 1.23 V vs. RHE. The inset shows the equivalent circuit model used to fit the Nyquist plots.
As discussed above, doped H mainly exists in the forms of Ho and Hi.[38, 39] Both Ho and Hi can act as shallow donors to improve the electrical conductivity. Mott-Schottky and electrochemical impedance spectroscopy (EIS) measurements are employed to characterize the electrical properties. Both samples show positive slopes in the Mott-Schottky plots (Fig. 3a), as expected for n-type semiconductors. In addition, a-TiO2:H shows a substantially smaller slope than undoped a-TiO2, suggesting an increased donor density. To more intuitively analyze the donor density, we calculate the donor density (Nd) and flat-band potential (Vfb) for both the samples from the slopes of the Mott-Schottky plots using the following equations: 2 2 d 1/ C 1 2 kT and Nd (V Vfb ) 2 2 C 0 A eN d e e0 0 dV 1
(1)
where e0 is the electron charge, 𝜀 the dielectric constant of TiO2 ( 𝜀 =73),[40] 𝜀 0 is the permittivity of vacuum, and V is the applied bias at the electrode. C and A are the interfacial capacitance and area, respectively, k is Boltzmann’s constant, T is the absolute temperature, and e is the electronic charge. Therefore, a plot of 1/C2 against V should yield a straight line from which Vfb can be determined from the intercept on the V axis. The calculated values of the donor density and Vfb are shown in Table 1. Undoped aTiO2 demonstrates the properties of an n-type semiconductor, whose electrons originate from intrinsic defects, such as oxygen vacancies and unintentional H-doping. An improvement of two orders of magnitude in the donor density is achieved after H-doping. H-induced defects in a-TiO2:H mainly consist of two forms, Hi and Ho.
2TiO2 3H 2 2Ti H O( Hi ) 2H 2O 2TiO2 5H 2 2Ti H ( HO ) 4H 2O
(2) (3)
Among of these forms, Hi and Ho are located at the shallow donor energy levels, favorable to provide donors.[38] Thus, a-TiO2:H shows a higher donor density than undoped a-TiO2. 9
Additional potential advantages of these H-bonding deficiencies involve the deficiencies passivating the recombination centers (Vo and dangling bonds), providing trap sites for photogenerated carriers and preventing these trap sites from rapid recombination, thus promoting electron transfer and photocatalytic reactions. In addition, the Fermi energy level is a function of the carrier density, and the relationship between the Fermi energy level and Vfb has been discussed in our previous report.[41] A more negative flat-band potential of a-TiO2:H indicates larger band bending, which will accelerate charge separation and transfer at the semiconductor/electrolyte interface (as shown in Fig. 3b). In addition, a more negative flat-band potential can also block electrons and suppress charge recombination at the interface. The potential vs. Ag/AgCl can be converted into the potential vs. the normal hydrogen electrode (NHE) as follows: 0 ENHE E Ag / AgCl EAg / AgCl ( 0.197 V)
(4)
Table 2. Electrical properties and catalytic activities of the undoped a-TiO2 and a-TiO2:H photoanodes. The fitting results calculated by the Nyquist plots (Figure 3c). Rs
Rtransport
Rtransfer
(𝛀 cm2)
(𝛀 cm2)
(𝛀 cm2)
Undoped a-TiO2
26.2
1590
39590
a-TiO2:H
26.3
473.6
3526
Samples
The electrical performance for charge transport and transfer can be calculated by fitting the Nyquist plots of undoped a-TiO2 and a-TiO2:H (Fig. 3c), which are deduced from the EIS measurements under illumination. A smaller diameter is found for the a-TiO2:H film compared with undoped a-TiO2, suggesting better PEC performance. To more intuitively explain this behavior, Table 2 shows the external circuit resistance (Rs), charge-transport resistance (Rtransport) and charge-transfer resistance (Rtransfer) deduced from the equivalent circuit fitting (the inset image in Fig. 3c). Rs represents the charge-transport capacity in the external circuit and electrolyte. Rtransport corresponds to the charge-transport capacity in the bulk of TiO2. Rtransfer represents the charge-transfer resistance at the semiconductor/electrolyte interface. From Table 2, Rs is approximately the same for the two samples, suggesting a 10
stable testing environment. Both Rtransport and Rtransfer of a-TiO2:H decreases greatly, indicating faster charge transport and transfer. We believe that this behavior is attributed to the improved donor density and larger band bending.
Figure 4. PEC performances of undoped a-TiO2 and a-TiO2:H deposited on FTO glass substrates. a) Three-electrode voltammograms curves. b) Photoconversion efficiencies at different applied voltages.
Fig. 4a demonstrates the three-electrode voltammograms of undoped a-TiO2 and aTiO2:H. Compared with undoped a-TiO2, a considerable improvement in the saturation current density is obtained for a-TiO2:H. The saturation photocurrent density increases from 0.13 mA/cm2 to 0.67 mA/cm2 at 1.23 V vs. the reversible hydrogen electrode (RHE) for aTiO2:H compared to undoped a-TiO2. The potential vs. Ag/AgCl can be converted into the potential vs. RHE as follows: 0 ERHE E Ag / AgCl EAg / AgCl ( 0.197 V) 0.0591 pH
(5)
The higher saturation current density of a-TiO2:H is mainly attributed to three issues: higher light absorption, smaller charge-transport and charge-transfer resistance, and less recombination centers. Meanwhile, a-TiO2:H can achieve saturation current at a more negative applied voltage, indicating better catalytic performance. The photoconversion efficiency can be calculated as follows:[42]
J mp 1.23 Vmp Pin
(6)
where Jmp and Vmp are the current density and voltage at the maximum power point (MPP), respectively, and Pin is the power of the incoming illumination. The photoconversion 11
efficiencies with different applied voltages of undoped a-TiO2 and a-TiO2:H are presented in Fig. 4b. An order of magnitude improvement in the photoconversion efficiency (from 0.02% to 0.26%) for a-TiO2:H is achieved compared with undoped a-TiO2. Compared with most hydrogenated crystalline TiO2 materials (Table S2), our a-TiO2:H film obtains a high photocurrent density of 674 μA/cm2 at an applied voltage of 1.23 V vs. RHE. This high photocurrent density is attributed to the high light absorption, reduced recombination centers, and more shallow donors (Hi and Ho) to provide a higher carrier concentration and accelerate charge transport and transfer.
4. Conclusions In summary, we successfully prepared a-TiO2:H at room temperature with simple magnetron sputtering technology. We reported an efficient theoretical model to explain the cause of the color change of a-TiO2:H. We believed that the introduced extended mid-gap states and band tail merging with the valence band causing by H-bonding was responsible for the enhanced VIS-NIR light absorption. In addition, a suitable H-doping flow rate passivated the recombination centers to reduce carrier recombination and introduced shallow donors (Ho and Hi) to provide free carriers, thus improving the carrier transport and PEC performance. Finally, based on a planar structure, we obtained a photocurrent density of 674 𝜇A/cm2 for aTiO2:H at an applied voltage of 1.23 V vs. RHE, which is comparable with most crystalline TiO2 materials, even with nanostructures. We gained insight into the mechanism to understand the cause of the color change for hydrogenated TiO2. We believe the ability to fabricate highly photoactive a-TiO2:H films will open new opportunities in various areas.
Acknowledgements The authors gratefully acknowledge the financial supports from the International cooperation project of the Ministry of Science and Technology (2014DFE60170), National Natural Science Foundation of China (Grant No. 61474065 and 61674084), Tianjin Research Key Program of Application Foundation and Advanced Technology (15JCZDJC31300), Key Project in the Science & Technology Pillar Program of Jiangsu Province (BE2014147-3), and the 111 Project (B16027). 12
Appendix A. Supplementary material Supplementary data associated with this article can be found in the online version at
References [1] A. Fujishima, K. Honda, Nature 238 (1972) 37-38. [2] Y.H. Hu, Angew. Chem., Int. Ed. 51 (2012) 12410-12412. [3] X. Chen, L. Liu, F. Huang, Chem. Soc. Rev. 44 (2015) 1861-1885. [4] X. Chen, S.S. Mao, Chem. Rev. 107 (2007) 2891-2959. [5] Z. Wang, U. Helmersson, P. O. Käll, Thin Solid Films 405 (2002) 50-54. [6] A. Di Paola, M. Bellardita, L. Palmisano, Brookite, Catalysts 3 (2013) 36-73. [7] A. Kruth, S. Peglow, A. Quade, M. M. Pohl, R.d. Foest, V. Brüser, K. D. Weltmann, J. Phys. Chem. C 118 (2014) 25234-25244. [8] L. Mascaretti, S. Ferrulli, P. Mazzolini, C.S. Casari, V. Russo, R. Matarrese, I. Nova, G. Terraneo, N. Liu, P. Schmuki, A. Li Bassi, Sol. Energy Mater. Sol. Cells 169 (2017) 1927. [9] X.B. Chen, L. Liu, P.Y. Yu, S.S. Mao, Science 331 (2011) 746-750. [10] C. Fan, C. Chen, J. Wang, X. Fu, Z. Ren, G. Qian, Z. Wang, J. Mater. Chem. A 2 (2014) 16242-16249. [11] Z. Wang, C. Yang, T. Lin, H. Yin, P. Chen, D. Wan, F. Xu, F. Huang, J. Lin, X. Xie, M. Jiang, Adv. Funct. Mater. 23 (2013) 5444-5450. [12] G.M. Wang, H.Y. Wang, Y.C. Ling, Y.C. Tang, X.Y. Yang, R.C. Fitzmorris, C.C. Wang, J.Z. Zhang, Y. Li, Nano Lett. 11 (2011) 3026-3033. [13] M. Mehta, N. Kodan, S. Kumar, A. Kaushal, L. Mayrhofer, M. Walter, M. Moseler, A. Dey, S. Krishnamurthy, S. Basu, J. Mater. Chem. A 4 (2016) 2670-2681. [14] T. Leshuk, R. Parviz, P. Everett, H. Krishnakumar, R.A. Varin, F. Gu, ACS Appl. Mater. Interfaces 5 (2013) 1892-1895. [15] X. Chen, L. Liu, Z. Liu, M.A. Marcus, W.-C. Wang, N.A. Oyler, M.E. Grass, B. Mao, P.A. Glans, Y.Y. Peter, Sci. Rep. 3 (2013) 1510. [16] W. Zhou, W. Li, J. Q. Wang, Y. Qu, Y. Yang, Y. Xie, K. Zhang, L. Wang, H. Fu, D. Zhao, 13
J. Am. Chem. Soc. 136 (2014) 9280-9283. [17] K. Zhang, J.H. Park, J. Phys. Chem. Lett. 8 (2017) 199-207. [18] C.Y. Yang, Z. Wang, T.Q. Lin, H. Yin, X.J. Lu, D.Y. Wan, T. Xu, C. Zheng, J.H. Lin, F.Q. Huang, X.M. Xie, M.H. Jiangl, J. Am. Chem. Soc. 135 (2013) 17831-17838. [19] X. Lu, G. Wang, T. Zhai, M. Yu, J. Gan, Y. Tong, Y. Li, Nano Lett. 12 (2012) 1690-1696. [20] A. Naldoni, M. Allieta, S. Santangelo, M. Marelli, F. Fabbri, S. Cappelli, C.L. Bianchi, R. Psaro, V. Dal Santo, J. Am. Chem. Soc. 134 (2012) 7600-7603. [21] M. Choi, J.H. Lee, Y.J. Jang, D. Kim, J.S. Lee, H.M. Jang, K. Yong, Sci. Rep. 6 (2016) 36099. [22] S. Wendt, P.T. Sprunger, E. Lira, G.K. Madsen, Z. Li, J.Ø. Hansen, J. Matthiesen, A. Blekinge-Rasmussen, E. Lægsgaard, B. Hammer, Science 320 (2008) 1755-1759. [23] L. Liu, P.Y. Yu, X. Chen, S.S. Mao, D.Z. Shen, Phys. Rev. Lett. 111 (2013) 065505. [24] H.H. Pham, L. W. Wang, Phys. Chem. Chem. Phys. 17 (2015) 11908-11913. [25] S. Pan, X. Liu, M. Guo, S.f. Yu, H. Huang, H. Fan, G. Li, J. Mater. Chem. A 3 (2015) 11437-11443. [26] F. Zhang, Y. Zheng, Y. Cao, C. Chen, Y. Zhan, X. Lin, Q. Zheng, K. Wei, J. Zhu, J. Mater. Chem. 19 (2009) 2771-2777. [27] M. Hari, S.A. Joseph, S. Mathew, P. Radhakrishnan, V. Nampoori, J. Appl. Phys. 112 (2012) 074307. [28] Z.W. Seh, S. Liu, M. Low, S.Y. Zhang, Z. Liu, A. Mlayah, M.Y. Han, Adv. Mater. 24 (2012) 2310-2314. [29] C. Fan, C. Chen, J. Wang, X. Fu, Z. Ren, G. Qian, Z. Wang, Sci. Rep. 5 (2015) 11712. [30] D.O. Scanlon, C.W. Dunnill, J. Buckeridge, S.A. Shevlin, A.J. Logsdail, S.M. Woodley, C.R.A. Catlow, M.J. Powell, R.G. Palgrave, I.P. Parkin, G.W. Watson, T.W. Keal, P. Sherwood, A. Walsh, A.A. Sokol, Nat. Mater. 12 (2013) 798-801. [31] K. Eufinger, D. Poelman, H. Poelman, R. De Gryse, G.B. Marin, Appl. Surf. Sci. 254 (2007) 148-152. [32] D.M. King, X. Du, A.S. Cavanagh, A.W. Weimer, Nanotechnology 19 (2008) 445401. [33] D.M. Hofmann, A. Hofstaetter, F. Leiter, H. Zhou, F. Henecker, B.K. Meyer, S.B. Orlinskii, J. Schmidt, P.G. Baranov, Phys. Rev. Lett. 88 (2002) 045504. 14
[34] M. Lazarus, T. Sham, Chem. Phys. Lett. 92 (1982) 670-674. [35] Z. Zheng, B. Huang, J. Lu, Z. Wang, X. Qin, X. Zhang, Y. Dai, M. H. Whangbo, Chem. Commun. 48 (2012) 5733-5735. [36] E. McCafferty, J. Wightman, Surf. Interface Anal. 26 (1998) 549-564. [37] A. Imanishi, T. Okamura, N. Ohashi, R. Nakamura, Y. Nakato, J. Am. Chem. Soc. 129 (2007) 11569-11578. [38] A. Janotti, C.G. Van de Walle, Nat. Mater. 6 (2007) 44-47. [39] C. Sun, Y. Jia, X. H. Yang, H. G. Yang, X. Yao, G.Q. Lu, A. Selloni, S.C. Smith, J. Phys. Chem. C 115 (2011) 25590-25594. [40] M.D. Stamate, Appl. Surf. Sci. 218 (2003) 318-323. [41] J. Liang, H. Tan, M. Liu, B. Liu, N. Wang, Q. Zhang, Y. Zhao, A.H.M. Smets, M. Zeman, X. Zhang, J. Mater. Chem. A 4 (2016) 16841-16848. [42] M.G. Walter, E.L. Warren, J.R. McKone, S.W. Boettcher, Q. Mi, E.A. Santori, N.S. Lewis, Chem. Rev. 110 (2010) 6446-6473.
Graphical Abstract
A theoretical model, the cooperative effects of the extended mid-gap states and valence band tail caused by H-induced deficiencies, is proposed to explore the underlying cause for enhanced light harvesting of H-doped black-blue amorphous TiO2.
Highlights Magnetron sputtering technique has been employed to fabricate the pure amorphous black-blue TiO2:H. 15
Suitable H-doping flow rate can passivate the recombination centers to reduce carrier recombination and improve conductivity. A theoretical model is proposed to explore the underlying cause for enhanced light harvesting of H-doping black-blue TiO2. Based on a planar structure, we obtain a photocurrent density of 674 𝜇A/cm2 for a-TiO2:H at an applied voltage of 1.23 V vs. RHE
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