On the Theoretical and Experimental Control of Defect Chemistry

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On the Theoretical and Experimental Control of Defect Chemistry, Electrical and Photoelectrochemical Properties of Hematite Nanostructures Jian Wang, Nicola H Perry, Liejin Guo, Lionel Vayssieres, and Harry L. Tuller ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b16911 • Publication Date (Web): 21 Dec 2018 Downloaded from http://pubs.acs.org on December 21, 2018

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ACS Applied Materials & Interfaces

On the Theoretical and Experimental Control of Defect Chemistry, Electrical and Photoelectrochemical Properties of Hematite Nanostructures Jian Wang1,2, Nicola H. Perry2,3,†, Liejin Guo1, Lionel Vayssieres1,*, Harry L. Tuller2,3,* 1

International Research Center for Renewable Energy, State Key Laboratory of Multiphase Flow in Power

Engineering, Xi’an Jiaotong University, Xi’an, 710049, China 2

Department of Materials Science & Engineering, Massachusetts Institute of Technology, Cambridge, MA

02139, U.S.A. 3

International Institute for Carbon Neutral Energy Research (I2CNER), Kyushu University, Fukuoka 819-0395,

Japan † current address: Department of Materials Science & Engineering and Materials Research Laboratory, University of Illinois at Urbana-Champaign, Urbana IL 61801, U.S.A.

KEYWORDS α-Fe2O3, dilatometry, electrical conductivity, defect chemistry, photoelectrochemistry

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ABSTRACT Hematite (α-Fe2O3) is regarded as one of the most promising cost-effective and stable anode materials in photoelectrochemical applications, and its performance, like other transition metal oxides, depends strongly on its electrical and defect properties. In this work, the electrical and thermo-mechanical properties of undoped and Sn-doped α-Fe2O3 nanoscale powders were characterized in-situ under controlled temperatures (T = 250 to 400oC) and atmospheres (pO2 = 10-4 to 1 atm O2) to investigate their transport and defect properties. Frequencydependent complex impedance spectra show that interfacial resistance between particles is negligible in comparison with particle resistance. Detailed defect models predicting the dependence of electron, hole, iron and oxygen vacancy concentrations on temperature and oxygen partial pressures for undoped and doped αFe2O3 were derived. Using these defect equilibria models, the operative defect regimes were established and the bandgap energy of undoped α-Fe2O3 and oxidation enthalpy of Sn-doped α-Fe2O3 were obtained from the analysis of the temperature and pO2 dependence of the electrical conductivity. Based on these results, we are able to explain the surprisingly weak impact of donor doping on the electrical conductivity of α-Fe2O3. Furthermore, experimental means based on the results of this study are given for successfully tuning hematite to enhance its photocatalytic activity for the water oxidation reaction.


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INTRODUCTION Nanostructured materials often exhibit distinct characteristics of high relevance for efficient energy storage and conversion applications. In particular, hematite (α-Fe2O3), has been studied with great interest as a low-cost, earth-abundant and non-toxic photoanode for photo-assisted hydrogen generation by water splitting.1-3 However, its photoelectrochemical (PEC) performance is limited by its low electron mobility (on the order of 10-2 cm2 V-1 s-1), very short hole diffusion length (2-4 nm) and slow water oxidation kinetics.4-7 Although many studies have been carried out to improve its performance, including morphology control, doping, and cocatalyst coupling,8-12 fundamental properties affecting performance, such as the impact of nanostructuring on defect chemistry, conduction, and stability, requires further attention.13 It is well known that zero-dimensional, or point defects, play a particularly important role in the control and optimization of semiconductor physicochemical properties.14-16 Despite the strong interest in α-Fe2O3 as a PEC photoanode candidate, its simple binary character and its abundance in the earth’s crust, there have been surprisingly limited amounts of reports on its underlying defect structure and corresponding bulk transport properties.17-20 There are even fewer studies regarding the defect properties of nanosized α-Fe2O3, which show particular promise for energy conversion applications. There is clear evidence that such nanosized particles can exhibit substantially different defect chemistries than their bulk counterparts.21-23 In this study, we are presenting our approach for achieving a detailed understanding of the defect regimes and electrical properties of α-Fe2O3 nanostructures (doped and undoped) and comment on their significance in relation to its use as a photoanode with enhanced PEC performance. Hematite is usually observed to be an n-type semiconductor as a result of loss of oxygen upon heating and the consequent formation of oxygen vacancies in the bulk.24 Doping is expected to modify the electrical conductivity by impacting the relative concentration of electrons and holes as well as the carrier mobility. However, the actual effect of different dopants and doping concentrations on the PEC properties of α-Fe2O3 photoanodes remains unclear.25-26 For example, tetravalent tin (Sn) is one of the most popular donor dopants for hematite, but nevertheless, exhibits a surprisingly weak dopant effect. Ling et al found less than a three-fold increase in electron density (1.89 × 1019 vs 5.38 × 1019 cm-3) following the diffusion of Sn into hematite (9.9% ACS Paragon Plus Environment

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Sn) as an inadvertent dopant from fluorine-doped tin oxide (FTO) conductive substrates upon high-temperature (≥ 750oC) annealing.27 Goncalves and Leite showed an even weaker effect on electron density (1.3 × 1018 to 1.7 × 1018 cm-3) upon Sn doping (8% Sn).28 Carvalho and Souza also reported that the electron density remains nearly unchanged when comparing pure and Sn-doped α-Fe2O3.29 In a later study, they found that α-Fe2O3 photoelectrodes could only be activated by the synergetic effect of Sn doping and oxygen deficient atmosphere.30 Although most of these studies reveal considerable enhancements in photocurrents obtained with Sn-doped α-Fe2O3 photoelectrodes, the exact role of Sn incorporation remains unclear and of various origins, somewhat controversial.30-31 We have investigated the defect chemistry of undoped and Sn-doped hematite nanostructures to better understand the nature of the defects that are formed to compensate the net charge of the four valent donors. This follows an earlier study on Ti-doped α-Fe2O3 thin films in which there was evidence suggesting that negatively charged Fe vacancies compensate for the Ti donor instead of electrons under oxidizing conditions.13 In this present study, the electrical and thermo-mechanical properties of undoped and Sn-doped α-Fe2O3 nanopowders were investigated by combined in-situ dilatometry and electrochemical impedance spectroscopy (EIS) as means of exploring the impact of temperature, oxygen partial pressure (pO2), and nanostructuring on the defect and transport properties. This was achieved using a modified dilatometer and specialized sample holder as previously described.21 Dilatometry was carried out to track the thermal and chemical expansion/contraction of the α-Fe2O3 nanoparticles and their correlation with defect generation/annihilation. The impedance spectra were generally found to form nearly perfect semicircles represented by a resistor in parallel with a constant phase capacitive element. Such a result, together with the direct-current (DC) voltage independence of conductivity, points to negligible interfacial resistance between nanoparticles. Defect equilibria models were adopted to understand which defect regimes are operative under the different experimental conditions. We demonstrate that the conductivity results can be successfully analyzed by our proposed defect model with key thermodynamic and transport parameters derived.


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EXPERIMENTAL SECTION Sample Preparation The preparation method used here is the typical surfactant-free, low temperature aqueous chemical growth technique to synthesize hematite nanorods and nanorod-arrays.32-33 Typically, a 50 mL aqueous solution of 0.15 M FeCl3 (Alfa, Aesar, iron (III) chloride, 97.0-102% purity) and 1 M NaNO3 (Alfa Aesar, sodium nitrate, ≥99.0% purity) at pH 1.5 (adjusted by HCl) was poured in a 125 mL glass bottle and maintained at 100oC for 4 h in a regular oven. After the reaction ended and cooled down naturally, the precipitated powders were collected and washed five times with deionized water, assisted with a centrifuge. The obtained yellow (βFeOOH, akaganeite) powders were then dried at 60oC for 8 h, and subsequently heated at 400oC for 2 h in air to turn into α-Fe2O3. The Sn-doped α-Fe2O3 powders were prepared using the same method but with an additional selected amount of SnCl4·xH2O (Alfa Aesar, tin (IV) chloride hydrate, 98% purity) into the solution. The exact atomic percent of Sn in α-Fe2O3 was estimated to be 1.7 ± 0.7% from energy-dispersive spectroscopy (EDS) analysis. The method for preparing α-Fe2O3 photoelectrodes was the same as that used for preparing α-Fe2O3 powders except that one FTO substrate was put inside the bottle with its conductive side facing to the wall. After 4 h of reaction at 100oC, a uniform akaganeite film was formed on the FTO substrate. The film was thoroughly rinsed with DI water, dried on a hot plat, and then heated in a muffle furnace in air at 500oC for 3 h (to form α-Fe2O3). After that, the films were sintered in a rapid temperature processing (RTP) tube furnace (OTF-1200X, Hefei Ke Jing Materials Technology Co., Ltd.) at 800oC for 5 min in pure oxygen (identified as O2/Sn-Fe2O3), air atmosphere (identified as Air/Sn-Fe2O3) and 1000 ppm H2 in Ar bubbled through roomtemperature water with a pO2 of ~10-13 atm (identified as (H2/H2O/Ar)/Sn-Fe2O3), respectively. 800oC was chosen because Sn dopant from the FTO substrates could diffuse into α-Fe2O3 at this temperature.27, 34 Physical Characterization The sample morphologies were investigated by field-emission scanning electron microscopy (FESEM, JEOL 7800F, Japan) and transmission electron microscopy (TEM, FEI G2F30, USA). The crystal structure of the samples was characterized by x-ray diffraction (XRD, PANalytical X’pert PRO MPD). The elemental


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compositions and chemical states were analyzed with x-ray photoelectron spectrometry (XPS, Shimadzu/Kratos Analytical AXIS Ultra DLD, UK). Electrical and Thermo-mechanical Measurements In-situ electrochemical impedance spectroscopy and thermal-mechanical expansion were measured simultaneously in a modified dilatometer (L75 platinum series, Linseis) as reported previously.21 The initial sample heights in the sample holder were 0.6 and 1.2 mm for undoped and Sn-doped α-Fe2O3, respectively with change in height upon cooling and heating measured by the dilatometer. The thermal expansion of the empty sample holder with Pt electrodes could be neglected because it only accounts to approximately 1% of the thermal expansion of the sample itself. The electrical properties of the samples were measured at various temperatures and oxygen partial pressures by impedance spectroscopy (ModuLab XM MTS) over the frequency range of 0.01 Hz to 1 MHz with an AC amplitude of 300 mV. The sample measurements were limited to a maximum of 400oC. Measurement procedures followed the following protocol: heating up to 400 oC at a rate of 5oC min-1 and hold for 24 h at a given pO2. Subsequently, decrease temperature at 5oC min-1 in 25 oC increments. At each increment, the temperature was held for 3 h and impedance data measured every ten minutes. The conductivity was calculated from the fitting resistance of the equivalent circuit model using the equation: =

l RA

where l is the sample thickness adjusted according to the dilatometry measurement, R is the resistance obtained by the fit to the impedance spectra, and A is the area determined by the size of the electrodes in contact with the powder. PEC Measurements Photocurrent densities of Sn-doped α-Fe2O3 electrodes were measured in a typical three-electrode cell with a Ag/AgCl electrode as reference electrode, a Pt counter electrode and α-Fe2O3 thin film as working electrode. The simulated solar illumination was obtained using a 500 W Xe arc lamp equipped with an AM 1.5 G filter (light intensity: 100 mW cm-2). The electrolyte was 1 M NaOH aqueous solution (pH 13.6) and a CHI 760D potentiostat was used during the measurements. The contact area between electrode and electrolyte was 0.785 ACS Paragon Plus Environment

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cm2. The applied potential vs. Ag/AgCl was converted to reversible hydrogen electrode (RHE) using the equation: ERHE = EAg/AgCl + 0.059 × pH + 0.1976 V. Mott-Schottky measurements were carried out in the dark at frequency of 1 kHz with AC amplitude of 10 mV. The donor density, ND, and flat band potential, VFB, can be estimated by the following equation:

1 C 2   2 e 0 N D V  VFB  kT e  in which C is the specific capacitance, e the electron charge, ε the relative permittivity of semiconductor (80 for hematite35), ε0 the permittivity of vacuum, k the Boltzmann constant and T the absolute temperature. THEORETICAL BASIS Defect Chemistry of α-Fe2O3 Our previous study presented electrical measurements performed on Ti-doped α-Fe2O3 thin films, and a preliminary defect equilibrium model was developed.13 To begin, we consider intrinsic ionic and electronic disorder in α-Fe2O3. Intrinsic ionic disorder in α-Fe2O3 is reported to occur via the formation of Schottky defects as described in the following defect reaction written in Kröger-Vink notation:17   3VO nil  2VFe

where

VFe ,

and

VO

(1) are triply negatively and doubly positively charged vacancies on Fe and O sites, respectively.

The corresponding mass action relation is given as: 2 3  E  VFe  VO   K S T   K S0 exp   S   kT 

(2)

where square brackets denote the concentration of the charge carrier species, KS is the equilibrium constant, ES the Schottky formation energy,

KS0

the pre-exponential factor, k the Boltzmann constant, and T the absolute

temperature. The well-known intrinsic electron-hole generation reaction where the carriers are excited across the bandgap together with the corresponding mass action reaction are shown below: nil  e  h


(3) 7


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(4)

 E  np  K i T  NC NV exp   g   kT 



where n and p are the electron and hole concentrations, NC and NV the effective densities of states in the conduction and valence band, respectively, and

Eg

the bandgap energy.

Under sufficiently reducing conditions, oxygen is released from the crystalline structure and accommodated by the formation of O vacancies and electrons as described by the following defect reaction: 1 OO  2e  VO  O2  g  2

where

OO

(5)

is an oxygen ion at an oxygen site. The corresponding mass action relation for the reduction reaction

can be written as:  E  n 2 VO  pO21/2  K R T  K R0 exp   R   kT 

 

(6)

in which KR is the equilibrium constant, ER the reduction enthalpy,

K R0

the pre-exponential term. Under

sufficiently oxidizing conditions, oxygen is incorporated into the lattice and accommodated by the formation of Fe vacancies and holes as follows: 3 O2  3OO  2VFe  6h 2

(7)

The corresponding mass action reaction for the oxidation reaction can be written as: 2  E  0   pO2-3 2  KOx T   KOx p 6 VFe exp   Ox   kT 

(8)

in which EOx is the oxidation enthalpy with equilibrium constant KOx and pre-exponential

0 . KOx

The equilibrium

constants of the afore-mentioned reactions are not independent of one another, and can be combined to:

KS 

K R 3KOx Ki 6

(9)

In undoped α-Fe2O3, the charge carriers consist of electrons, holes, iron vacancies, and oxygen vacancies. The overall charge neutrality condition of undoped α-Fe2O3 can then be expressed by: ACS Paragon Plus Environment

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n  3 VFe   p  2 VO 

(10)

In Sn-doped α-Fe2O3, Sn incorporation by substitution for Fe, can be described by the following equation: Fe2O3  2SnO2  2SnFe  3OOx  2e 

1 O2 2

(11)

Thus, the charge neutrality of Sn-doped α-Fe2O3 is given instead by:    2 VO  n  3 VFe   p  SnFe

(12)

To solve for the concentrations of various defects in terms of the equilibrium constants, we apply the Brouwer approximation, in which one assumes that only one term on either side of the electroneutrality equation predominates under given operating conditions of temperature and pO2. The regimes and predicted solutions to the defect concentrations of undoped and Sn-doped α-Fe2O3 are summarized in Table 1 and 2, respectively, as well as in Figure 1. RESULTS Morphology and Crystal Structure The sample morphologies were obtained by field-emission scanning electron microscopy (FE-SEM) and transmission electron microscopy (TEM), as shown in Figure 2a-d. It reveals that the undoped α-Fe2O3 sample consists of anisotropic (oblong) nanoparticles with size of 438 ± 158 nm (Figure 2a-b). In contrast, the Sndoped α-Fe2O3 sample is composed of smaller spheroidal nanoparticles with size of 67 ± 24 nm (Figure 2c-d). High resolution TEM (insets in Figure 2b and d) show distinct crystal lattice spacing of 0.22 and 0.37 nm in accordance with the (113) and (012) crystal planes of hematite, respectively. The Sn/(Sn+Fe) atomic ratio (Sn%) of Sn-doped α-Fe2O3 deduced from TEM energy-dispersive spectroscopy (EDS) analysis is 1.7 ± 0.7%. The EDS elemental mapping shows that Sn is successfully incorporated in Sn-doped α-Fe2O3 (bottom-left of Figure 2d). X-ray powder diffraction (XRD) patterns (shown in Figure 2e) reveal that all diffraction peaks match well with the rhombohedral structured hematite (JCPDS No. 24-0072). No impurity phase was detected in both undoped and Sn-doped α-Fe2O3, indicating no other phase, such as SnO2, was formed. The relatively lower and broader peaks of Sn-doped α-Fe2O3 suggest much smaller crystallite sizes. The elemental chemical states and composition of the samples were analyzed by X-ray photoelectron spectroscopy (XPS). As shown in ACS Paragon Plus Environment

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Figure 2f, while Fe and O were detected in both undoped and Sn-doped samples, only Sn was present in the latter. The very small Cl 2p signal originates from the precursor residual. Fe 2p XPS spectra (Figure 2g) display two distinct peaks of Fe 2p1/2 at 724.3 eV and Fe 2p3/2 at 710.7 eV with a satellite peak at 719.1 eV, which are typical values observed for Fe3+ in α-Fe2O3.36 One may note the difference in the spectra located at ~ 716 eV which is attributed to the extra satellite peak at 716.7 eV arising from the Fe2+ species in α-Fe2O3 that is expected with Sn doping.30, 37 The binding energy of O 1s main line (529.8 eV for the undoped and 530.0 eV for the Sn-doped sample) is consistent with the reported value for α-Fe2O3, as shown in Figure 2h.38 Two shoulders at higher binding energy for both samples could be assigned to -OH and H2O, respectively.38 Sn 3d XPS spectrum (Figure 2i) for the Sn-doped sample exhibits two distinct peaks at 495.0 and 486.6 eV, corresponding to the Sn 3d3/2 and Sn 3d5/2 peaks, respectively.10 These results confirm the successful doping of Sn, with creation of Fe2+, in α-Fe2O3. Thermal Expansion and Contraction The nanopowders were subsequently used as starting materials for the thermo-mechanical and electrical characterization. As discussed in our previous study, the dilatometry and electrical measurements of the nanopowders were performed simultaneously in a modified dilatometer setup (shown in Figure 3a and b).21 This specially designed sample holder consists of two smaller quartz tubes inside a larger one, two quartz rods, with one serving as a pushrod, and two platinum foils as electrodes attached to platinum wires connected to an impedance analyzer (Figure 3c). The powders exhibited a general thermal expansion upon the temperature ramp, taken in 50oC steps, from room temperature to 800oC in simulated air (21% O2 in N2), as displayed in Figure 4a. At each temperature step, an expansion can be observed. During the initial annealing treatment below 200oC, a small shrinkage was observed which is attributed to the desorption of water from the nanoparticles surface. Above 500oC, shrinkage is observed following the initial expansion and with significant shrinkage observed at temperatures above 600oC. The observed shrinkage very likely originates from the densification of the nanoparticles due to sintering into a denser structure. To maintain stable nanopowder structure, subsequent measurements on as-prepared nanopowders were only conducted at temperatures below 400oC. As shown in Figure 4b, the sample was first heated to 400oC at 5oC min-1 and held for 24 h at a given ACS Paragon Plus Environment

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pO2. Subsequently, the temperature was decreased at 5 oC min-1 in 25 oC increments. At each increment, the temperature was held for 3 h and electrochemical impedance spectra (EIS) were measured every 10 minutes. The shrinkage (~ 1 m) at 400oC after the 24 h anneal is negligible in comparison with initial sample thickness (~600 m). In each temperature step change, the temperature reaches a steady state value in less than 40 minutes and the conductivity reaches a steady state value in about 1 h. Even at 250oC, the system could reach thermodynamic equilibrium in about 2 h (see inset of Figure 4b). Electrochemical Characterization Electrochemical impedance spectroscopy (EIS) measurements were performed as a function of temperature and oxygen partial pressure. Figure 5a and e show typical Nyquist plots of undoped and Sn-doped α-Fe2O3 obtained at 400oC as a function of pO2 respectively. Under all experimental conditions, the measured impedance spectra were nearly ideal semicircles displayed from the origin on the real axis. This is different from our earlier study of TiO2 quantum dots whose impedance spectra showed a somewhat depressed semicircle at higher frequencies and a second semicircle at low frequencies representing bulk impedance, and particle-to-particle impedance, respectively.21 Although the two (undoped and doped) samples have very different morphologies, their impedance spectra shapes are similar. The spectra of α-Fe2O3 are all represented by one resistor in parallel with a constant phase element (CPE). CPEs are used to take into account any inhomogeneities resulting in “depressed” arcs not well represented by ideal capacitors.39-40 The impedance of a CPE is given by:

Z

1 Q  i 

(13)

nq

where ω, i, and nq are angular frequency,

1

, and the factor related to the deviation from ideal capacitance,

respectively. The equivalent capacitance is derived from Q using the following equation:41



C R

1 nq

Q



1 nq

(14)

The equivalent circuit fits match the data very well. Typical nq values are near unity (ca. 0.992-0.997) for Q demonstrating nearly ideal capacitances. The intercept of the spectra with the real axis at low frequency, which ACS Paragon Plus Environment

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represents the resistance, increases with increasing pO2 indicating that undoped and Sn-doped α-Fe2O3 are both n-type semiconductors. To establish the nature of the resistance derived from the IS measurements, i.e. bulk vs interfacial, the impedance was measured as a function of DC bias. The bulk ohmic resistance should show no bias dependence, while interfacial resistances is expected to show a bias dependence if they are caused by space-charge barriers, since concentration profiles of charged species in space charge regions are modified under bias.42-43 Figures 5b and f show that DC biases of up to 5 V (electric field intensity of ~5000 V m-1 on the sample) show no effect on the effective conductivities at 400oC derived from the IS, thus supporting the interpretation that the measured resistance and capacitance arise from the bulk (i.e. the nanoparticles themselves) rather than from the interfaces between the nanoparticles. This result differs form that reported by Warren et al. in their study of α-Fe2O3 cauliflower electrodes where electron transport was interpreted as being limited by potential barriers formed at high-angle grain boundaries.5 This discrepancy may arise from the different surface chemistries of their nanoparticles (different preparation method) as well as the much lower temperatures at which their measurements were performed. Figures 5c and g show the Arrhenius plots of undoped and Sn-doped α-Fe2O3 in terms of log T vs 1000/T for isobaric conditions, respectively. The activation energy, Ea, is obtained from the slopes of the linear fits according to:  Ea    kT 

   0 exp  

(15)

where σ is conductivity, σ0 a pre-exponential constant, k Boltzmann’s constant and T the absolute temperature. It shows that the activation energies of undoped α-Fe2O3 increase with decreasing the pO2, obtaining values of 1.05 ± 0.02 eV at pO2 = 1 atm and 1.20 ± 0.01 eV at pO2 = 0.0001 atm, respectively. This is due to the transition of defect regime from II to I (see Figure 1a), which is discussed thereafter. Because most of α-Fe2O3 photoelectrodes for PEC water splitting in the literature were heated in air to improve their performance,10, 27, 34 in this study, we focus on the activation energy at pO2 = 0.21 atm, which is 1.06 ± 0.01 eV. The activation energies of Sn-doped α-Fe2O3 remain the same with varying the pO2, which are 1.05 ± 0.01 eV at all pO2s. ACS Paragon Plus Environment

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The electrical conductivity of undoped and Sn-doped α-Fe2O3 powders as a function of pO2 are plotted respectively in Figure 5d and h at isotherms from 250oC to 400oC. The conductivity of undoped α-Fe2O3 exhibits insensitivity to pO2 at the lower temperatures, but gradually shows a weakly increasing conductivity with decreasing pO2 at higher temperatures. It results in the increasing activation energy of undoped α-Fe2O3 with decreasing pO2 observed in Figure 5c. The conductivity of Sn-doped α-Fe2O3, on the other hand, exhibits a strong increase in conductivity with decreasing pO2 at all temperatures, displaying a power law dependence with a slope of approximately -1/4. In contrast with Ti-doped α-Fe2O3 whose conductivity reportedly becomes independent of pO2 below 0.001 atm O2 at as low as 100oC,13 the conductivity of Sn-doped α-Fe2O3 is not fixed at the donor concentration, but varies with the pO2 between 10-4 and 1 atm O2. These results are further discussed below in terms of the defect chemistry of hematite. DISCUSSION The electronic conductivity of undoped α-Fe2O3 exhibits nearly pO2 independence, more so at the lower temperatures. Based on the defect model, it points to the intrinsic electronic defect control (see regime II in Figure 1a). With n=p (regime II) and electron and hole mobilities much higher than ion mobilities, we obtain an expression for the total conductivity of undoped α-Fe2O3 given by: 

Eg   2  kT 

  nen  pe p  e  n   p  NC NV exp  

(16)

Assuming that electron and hole mobilities and effective density of states in the conduction and valence band are only weakly temperature dependent, the activation energy is given by Ea = Eg/2. Thus, the thermal bandgap energy of undoped α-Fe2O3 is calculated to be 2.12 ± 0.02 eV (at pO2 = 0.21 atm), comparable to results previously published in the literature.44 With increasing temperature, the driving force to form oxygen vacancies increases, ultimately resulting in a transition to defect regime I. This transition is not observed over the relatively narrow temperature range studied here but is suggested by the beginning of pO2 sensitivity under the more reducing conditions at temperatures above 325oC. Next, we consider Sn-doped α-Fe2O3 data. Given the -1/4 power law dependence of conductivity on pO2, this points to defect regime III in Figure 1b for which

  3 VFe   SnFe

as the experimentally operative regime. In this


13


Page 14 of 34

regime, the Sn4+ donors are compensated largely by iron vacancies rather than electrons. In the literature on Sndoped α-Fe2O3, the electron concentration is usually ascribed to donor compensation or due to the sum of concentrations of dopant donors and oxygen vacancies.45 Obviously, this fails to recognize the importance of Fe vacancies as a critical factor in compensating for the Sn donor. Our previous study on Ti-doped hematite indicated that the conductivity could vary with changes in the annealing pO2 at elevated temperatures and stated the importance of iron vacancies as likely compensating defects for Ti donors.13 However, since the defect regime of Ti-doped α-Fe2O3 did not fall within regime III, the compensation of Ti by iron vacancies was not further investigated. The -1/4 power law dependence of conductivity on pO2 indicates predominant n-type conduction in Sn-doped α-Fe2O3 (e.g. on the left side of region III of Figure 1b); the hole contribution can thus be neglected. The conductivity can thus be expressed as (see Table 2):    en 

NC NV 0 16 KOx

13

  SnFe    1   exp    3   kT  

EOx   Eg  6 

 1 4   pO2 

(17)

Assuming that the electron mobility is temperature independent, and disregarding the weak temperature dependencies of NC, NV and Ea  Eg  EOx 6

0 KOx

, we expect the activation energy of Sn-doped α-Fe2O3 to be quantified as:

. Interestingly, while the defect regimes in which the undoped and Sn-doped α-Fe2O3 operate are

different, the measured activation energies are very similar. We ascribe this to a coincidence, since the two materials exhibit very different pO2 dependencies. In studies that report bandgap values of α-Fe2O3, Sn doping does not result in measureable changes in bandgap energy.27, 46 Therefore, assuming the same Eg for Sn-doped α-Fe2O3 as for our undoped materials, the oxidation enthalpy of Sn-doped α-Fe2O3 is calculated to be 6.42 ± 0.06 eV. To the best of our knowledge, this is the first report of the oxidation enthalpy for α-Fe2O3. One may further question whether the conductivity of Sn-doped α-Fe2O3 will fall within regime II at lower pO2. To answer this, the conductivity was measured at lower pO2 with streaming H2/H2O/Ar gas mixtures. As shown in Figure 6, the conductivity does show pO2 independence at low pO2s, consistent with expectations for regime II within which the n-type conductivity of α-Fe2O3 is predicted to be independent of pO2, and satisfying the equality

  . n  SnFe

Here, the majority electrons fully compensate the charged donors and given the

temperature independent conductivity in this regime (see Figure 6), this further implies full ionization of the Sn ACS Paragon Plus Environment

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donors. The electron mobility can be calculated accordingly from the donor concentration of 1.7 ± 0.7% Sn (6.7 ± 2.7 × 1020 cm-3) to be 8 ± 3 × 10-4 cm2 V-1 s-1. This value is in agreement with some of the reported values of electron mobility for hematite,47-48 but lower than that in Ti- and Si-doped hematite films (~0.01 cm2 V-1 s-1).13, 49

It should be noted that the conductivity of Sn-doped α-Fe2O3 at low pO2 is nearly five orders of magnitude higher than that at high pO2 at 300oC. If one extrapolates the conductivity in air to room temperature according to the activation energy (see Figure 5g) and compare it with the value at low pO2, this difference can be as large as fourteen orders of magnitude given that the conductivity at low pO2 is temperature independent. This result suggests that the conductivity of Sn-doped α-Fe2O3 could be boosted at extremely low pO2. Accordingly, the PEC performance of Sn-doped α-Fe2O3 electrode could also be enhanced by reducing atmosphere treatment. To test this hypothesis, three α-Fe2O3 electrodes on fluorine-doped tin oxide (FTO) substrates were prepared and treated in pure oxygen (named as O2/Sn-Fe2O3), air (named as Air/Sn-Fe2O3) and reducing atmosphere (named as (H2/H2O/Ar)/Sn-Fe2O3) at 800oC, respectively. The annealing temperature was chosen to be 800oC to induce Sn diffusion from the substrate into α-Fe2O3 to cause Sn doping as widely reported in the literature.27, 31, 34 The preparation details are given in the experimental section. While the SEM images (Figure 7b-c) and XRD patterns (Figure 7e) of α-Fe2O3 electrodes annealing in various atmosphere display no significant difference, the measured photocurrents (Figure 7f) show that annealing in reducing atmosphere does have a positive effect on the photocurrent of Sn-doped α-Fe2O3 although obviously not proportional to the expected increase in the majority electron density. This is not entirely surprising since the overall rate limiting step for oxygen generation at the photoanode is photo-induced minority holes. Mott-Schottky plots shown in Figure 7g exhibit that the flat band potential of α-Fe2O3 electrodes shift negatively with decreasing the oxygen partial pressure during annealing. The donor densities were calculated to be 2.3×1020, 3.6×1020 and 8.3×1020 cm-3 for O2/SnFe2O3, Air/Sn-Fe2O3 and (H2/H2O/Ar)/Sn-Fe2O3 electrodes, respectively, demonstrating the increasing donor density with decreasing the oxygen partial pressure. It should be noted that the Mott-Schottky equation is derived from a planar electrode so these values are only for comparison purposes.50-51


15


Page 16 of 34

The long-term stability test, as shown in Figure 7h, indicates (H2/H2O/Ar)/Sn-Fe2O3 could retain its high photoactivity after a 24 h continuous operation without significant attenuation at 1.23 V vs. RHE. The crystal structure, determined by XRD, shows no changes after 24 h PEC measurements (Figure 7e). These results demonstrate that the defects induced by annealing in reducing atmosphere could be well maintained even in harsh conditions, e.g. strong alkaline solutions. Looking back at Figure 6, one finds that the transition pO2s from defect regime III to II are around 10-17 atm at 400oC and 10-23 atm at 300oC. However, in the case of Ti-doped α-Fe2O3 thin film with similar donor concentration (4 × 1020 cm-3), the transition pO2 is 10-4 atm in a similar temperature range.13 This implies that the oxidation enthalpy of our Sn-doped α-Fe2O3 nanopowders is lower than that of Ti-doped α-Fe2O3 thin film, given their greater readiness to oxidize and form iron vacancies. This could be understood by the fact that the thin films are chemo-mechanically constrained by their substrate while the nanopowders are free to expand/contract with loss/gain of oxygen.52 The donor dominated conductivity of Sn-doped α-Fe2O3 also appears to be lower than that of Ti-doped α-Fe2O3. This is very likely due to the very different sample morphologies (loosely packed nanopowders vs. dense thin film). The observed distinctions between Sn- and Tidoped α-Fe2O3 may explain, in part, the difference in PEC performance obtained in the literature for Sn- and Tidoped α-Fe2O3,53-54 although it should be noted that we do not presently have similar defect and electrical data for Ti-doped α-Fe2O3 nanopowders. To test this hypothesis, further studies are needed to fully understand the defect chemistry of α-Fe2O3, in its various morphologies. A detailed study of the role and impact of defects and their transport properties on photoelectrochemical performance is under way. CONCLUSIONS In conclusion, undoped and Sn-doped α-Fe2O3 nanopowders were thermo-mechanically and electrically characterized in-situ using a specially modified dilatometer. Specifically: 

At temperatures above 600oC, the loosely packed nanopowders undergo significant shrinkage and so all electrical characterization was limited to a maximum temperature of 400oC in order to maintain the initial nanoscale powder morphology.


16

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


The single, nearly ideal, semicircles obtained in the impedance spectra and the DC voltage independence of the measured resistance (ohmic behavior) indicated that the interfacial resistance between α-Fe2O3 particles was negligible in comparison to the particle resistance.



The electrical properties of undoped α-Fe2O3 were found to be largely pO2-independent and governed by intrinsic electron-hole generation at the elevated pO2s of this study. The bandgap energy was determined to be 2.12 ± 0.02 eV from the thermally activated electrical measurements performed at pO2 of 0.21 atm.



In contrast, the conductivity of Sn-doped α-Fe2O3, showed a power law dependence on pO2 with a -1/4 slope at pO2s between 10-4 and 1 atm. Based on the derived defect model, this points to donors being largely compensated by iron vacancies, rather than electrons. From the thermally activated electrical measurements in this regime, and assuming no change in bandgap energy upon Sn doping, the oxidation enthalpy was calculated to be 6.42 ± 0.06 eV. At considerably lower pO2s, the conductivity falls within the donor dopant controlled defect regime where Sn donors are compensated by electrons as evidenced by the pO2-independent conductivity. The effective electron mobility was derived to be 8 ± 3 × 10-4 cm2 V-1 s-1 in this regime.



Better fundamental understanding and experimental control of the distinct defect regimes and electrical properties of Sn-doped α-Fe2O3 yield to improved efficiency for photoelectrochemical water splitting55,56. Other applications such as sensors and energy storage for example could also benefits from the findings of this study on the theoretical and experimental control of the defect regime and electrical properties of hematite to improve devices performance.


17


Page 18 of 34

Table 1. Predicted defect densities of undoped α-Fe2O3 for each possible charge neutrality regime.

Regime

I

II

Defect

n  2 VO 

np

2K 

13

n

R

Ki

p

V   Fe

V   O

2K R 1 3

pO2

1 6

pO2

16

12

2

K Ox K R Ki

3

pO2

14

13

 KR     4 

pO2

1 6

 K Ox  K 3  i

KR Ki

III

Ki

12

Ki

12

 p  3VFe

Ki

9K 

18

pO2

3 16

Ox

9K Ox 1 8 pO2 3 16

12

   



18

pO2

pO2

3 4

1 2

 K Ox     729 

9K Ox 1 4

pO2

KR Ki

2

3 16

pO2

1 8


18

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Table 2. Predicted defect densities of Sn-doped α-Fe2O3 for each possible charge neutrality regime. Regime

I

II

III

Defect

n  2 VO 

  n  SnFe

  3 VFe   SnFe

2K 

13

n

R

Ki

p

V   Fe

V   O

pO2

pO2

2 K R 

13

K Ox K R Ki

3

pO2

14

Ki

pO2

1 6

12 3

Sn  pO  3 Fe

KR

Sn 

 2 Fe

pO2

   pO 1 4  SnFe  2   3 

13

Sn 

KOx

16

3  1 6  pO21 4 KOx     SnFe  

 Fe

13

 KR     4 

KOx

Ki

16

12

2

 Fe

p  3 VFe

13

Ki

Sn 

1 6

IV

Sn 

1 2

Ki

2

13

 3     Sn     Fe 

pO2

3 16

9K 

18

Ox

pO2

3 16

 K Ox  3 16   pO2  729 

3

K R KOx

18

Ox

18

 Fe

3 4

2

Ki

9K 



2 3

9K 

14

Ox

KR Ki

2

pO2

1 8


19


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Figure 1. Defect diagrams of α-Fe2O3 showing the pO2 dependent defect concentrations: (a) undoped α-Fe2O3 and (b) Sn-doped α-Fe2O3.


20

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Figure 2. Morphological and structural characterization of undoped and Sn-doped α-Fe2O3 nanopowder samples. (a and c) scanning electron microscopy (SEM) images with size distribution histograms insets (bottom left); (b and d) transmission electron microscopy (TEM) images of undoped α-Fe2O3 and Sn-doped α-Fe2O3, respectively; (e) X-ray powder diffraction (XRD) patterns; (f-i) X-ray photoelectron spectroscopy (XPS): (f) survey scan; (g) Fe 2p; (h) O 1s and (i) Sn 3d spectra. The top-right insets in (b and d) are the corresponding high resolution TEM images extracted from the regions within the white squares and the bottom-left inset in (d) is the EDS Sn mapping of Sn-doped α-Fe2O3.


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Figure 3. Modified dilatometer and sample holder used for nanopowder measurement. (a) Top view still image of the modified dilatometer. (b) Image of the home-made quartz sample holder in connection with the dilatometer with α-Fe2O3 powders inside. (c) Schematic diagram of the sample holder.


22

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Figure 4. Dilatometry and conductivity of α-Fe2O3 nanopowders as a function of temperature. (a) Thermal expansion (blue line) of α-Fe2O3 nanopowders from room temperature to 800oC in simulated air with the light and dark grey regions showing the negligible and significant shrinkage, respectively. (b) Simultaneous measurements of expansion and electrical properties in low temperature regime: expansion (blue line) and conductivity (black hollow squares) of as prepared undoped α-Fe2O3 nanopowder as a function of temperature in simulated air. The inset of (b) is the magnified conductivity and thermal expansion of α-Fe2O3 at 250 oC.


23


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Figure 5. Electrochemical characterization of (a-d) undoped and (e-h) Sn-doped α-Fe2O3 nanopowders. (a and e) Typical impedance spectra at 400oC at various pO2s with peak (arc top) frequencies indicated; (b and f) DC voltage dependence of the conductivity at 400oC at various pO2s; (c and g) Arrhenius isobaric plots in terms of log T vs 1000/T with linear fitting data and representative activation energy obtained at pO2 of 0.21 atm; (d and h) pO2 dependence of conductivity for several isotherms. In (a) and (e), the open symbols are the experimental data while the solid data points are the equivalent circuit fit, and the insets show the equivalent circuits used to fit the data.


24

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Figure 6. Conductivity of Sn-doped α-Fe2O3 nanopowder at low and high pO2s. Dashed lines represent fits from the defect model.


25


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Figure 7. Morphological, Structural and photoelectrochemical (PEC) characterization of α-Fe2O3 electrodes. Scanning electron microscopy (SEM) images: (a) as-prepared α-Fe2O3 electrode heated at 500oC for 3 h; (b, c and d) α-Fe2O3 electrodes sintered at 800oC for 5 min to induce Sn-doping in pure O2 (identified as O2/SnFe2O3), air atmosphere (identified as Air/Sn-Fe2O3) and 1000 ppm H2 in Ar bubbled through room-temperature water with a pO2 of ~10-13 atm (identified as (H2/H2O/Ar)/Sn-Fe2O3), respectively. (e) X-ray diffraction (XRD) patterns of Sn-Fe2O3 annealed in O2, air and H2/H2O/Ar treatment before and after 24 h PEC measurements along with SnO2 and Fe2O3 JCPDS standards. (f) Photocurrent-potential (J-V) curves obtained with α-Fe2O3 nanorod photoelectrodes. (g) Mott-Schottky plots collected in the dark. (h) Long-term stability test of (H2/H2O/Ar)/Sn-Fe2O3 electrode at 1.23 V vs. RHE. The electrodes were heated at 800 oC to induce Sn diffusion from the FTO substrate into α-Fe2O3.


26

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AUTHOR INFORMATION Corresponding Authors *Prof. Lionel Vayssieres (ORCID: 0000-0001-5085-5806) and Prof. Harry L. Tuller (ORCID: 0000-0001-83393222) Emails: [email protected] (L.V.) and [email protected] (H.L.T.) Author Contributions L.V. and H.L.T. proposed the project. J.W., N.H.P., L.V. and H.L.T. designed the experiment details and J.W. conducted the experiments. N.H.P. helped with the electrochemical measurements. L.G. helped with the physical characterization. All authors discussed and comments on the results. J.W. wrote the original manuscript. N.H.P., L.V. and H.L.T. revised the manuscript. Notes The authors declare no competing interests. ACKNOWLEDGEMENT This work was supported by the National Science Foundation under award number DMR-1507047 and National Natural Science Foundation of China under number 51236007 and 51323011. J. W. thanks the China Scholarship Council Postgraduate Scholarship Program provided by the Ministry of Education, China for his one year stay at MIT under the supervision of H.T. N. H. P. acknowledges support from WPI-I2CNER (MEXT, Japan). The authors also acknowledge Chang Sub Kim, Department of Materials Science and Engineering, Massachusetts Institute of Technology, and Penghui Guo, School of Energy and Power Engineering, Xi’an Jiaotong University for assistance with dilatometer and XPS measurements, respectively.


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Table of Contents Graphic


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On the Theoretical and Experimental Control of Defect Chemistry

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