Shaped‐controlled silicon‐doped hematite nanostructures for enhanced PEC water splitting

Accepted Manuscript Title: Shaped controlled silicon doped hematite nanostructures for enhanced PEC water splitting Authors: Mattia Allieta, Marcello Marelli, Francesco Malara, Claudia L. Bianchi, Saveria Santangelo, Claudia Triolo, ˇ ep´an Kment, Alessandro Salvatore Patane, Anna M. Ferretti, Stˇ Ponti, Alberto Naldoni PII: DOI: Reference:

S0920-5861(18)31102-7 https://doi.org/10.1016/j.cattod.2018.10.010 CATTOD 11673

To appear in:

Catalysis Today

Received date: Revised date: Accepted date:

14-7-2018 5-10-2018 8-10-2018

Please cite this article as: Allieta M, Marelli M, Malara F, Bianchi CL, Santangelo S, ˇ Ponti A, Naldoni A, Shaped controlled Triolo C, Patane S, Ferretti AM, Kment S, silicon doped hematite nanostructures for enhanced PEC water splitting, Catalysis Today (2018), https://doi.org/10.1016/j.cattod.2018.10.010 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 proof before it is published in its final 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.

Shaped controlled silicon doped hematite

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nanostructures for enhanced PEC water splitting

Mattia Allietaa, Marcello Marellia, Francesco Malaraa, Claudia L. Bianchib, Saveria

Santangeloc, Claudia Triolod, Salvatore Patane'd, Anna M. Ferrettia, Štěpán Kmente,*,

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Dipartimento di Chimica, Università degli Studi di Milano, Via Golgi 19, 20133 Milan, Italy

Dipartimento di Ingegneria Civile, dell’Energia, dell’Ambiente e dei Materiali (DICEAM),

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CNR-Istituto di Scienze e Tecnologie Molecolari, Via Golgi 19, 20133 Milan, Italy

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Alessandro Pontia,Alberto Naldonie,*

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Università“Mediterranea” di Reggio Calabria, Loc. Feo di Vito, 89122 Reggio Calabria, Italy Dipartimento di Scienze Matematiche e Informatiche, Scienze Fisiche e Scienze della Terra

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(MIFT), Università di Messina, Viale F. Stagno d’Alcontres, 31 - 98166 Messina, Italy e

Regional Centre of Advanced Technologies and Materials, Faculty of Science,

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PalackýUniversity Olomouc, 17. listopadu 1192/12, 771 46 Olomouc, Czech Republic

Graphical abstract

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Shape-controlled hematite (a-Fe2O3) nanostructures were fabricated using hydrothermal

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Highlights:

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method.

Different concentration of Si (1÷20 at.%) induced the formation of different crystal shape.

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Crystal preferential growth is directed by silicon.

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Silicon-doped hematite showed tunable electronic and magnetic properties.

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Photoelectrochemical activity of Silicon doped hematite nanocrystals is strongly dependent

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

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:

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from their shape.

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Shape-controlled hematite (a-Fe2O3) nanostructures were fabricated using hydrothermal method.

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Different concentration of Si (1÷20 at.%) induced the formation of different crystal shape.

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Crystal preferential growth is directed by silicon.

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Silicon-doped hematite showed tunable electronic and magnetic properties.

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Photoelectrochemical activity of Silicon doped hematite nanocrystals is strongly dependent

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from their shape.

Hematite (-Fe2O3) is one of the most promising photoanode materials for photoelectrochemical

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(PEC) water splitting although great challenges hinder high performance. Silicon-doped -Fe2O3

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shows improved PEC activity but the relationship among dopant content and enhanced

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conductivity, structure, and particle morphology is only poorly understood. Here, we present a

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systematic study on hydrothermally grown -Fe2O3 nanocrystals by using XRD, Raman, UV-vis

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spectroscopy, TEM, XPS, SQUID magnetometry, electrochemical impedance spectroscopy, and

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photocurrent measurements. We find that the Si content controls the morphology of -Fe2O3 already at Si 5 at.% inducing a transition from nanostructures with ellipsoidal shape to

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nanowires. Si doping is effective in improving PEC activity in the case of Si1% at. sample,

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which shows a 20% photocurrent enhancement in comparison with pure -Fe2O3. On the contrary, -Fe2O3 containing Si content higher than 5 at.% presents lower PEC activity. Results

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are rationalized in the view of the interplay of morphological, structural, magnetic, and electronic properties in doped -Fe2O3 thus providing general guidelines for the design of efficient photoelectrodes for solar water splitting.

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KEYWORDS: water splitting, substitutional doping, hematite, semiconductor, hydrothermal synthesis

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1. Introduction Photoelectrochemical water splitting is one the most prominent processes able to convert solar energy into sustainable energy vector such as hydrogen[1–7].The application of PEC water

splitting on a large scale requires cheap, abundant, stable, and nontoxic photoelectrode materials. Hematite fulfils all these requirements and recently has attracted a great interest as photoanode

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material due to its appropriate bandgap (2.0–2.2 eV) enabling the absorption of 40% of the solar

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spectrum[1,8].However, the low conductivity, the short hole diffusion length (typically 2–4 nm),

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and the fast electron-hole recombination limit the PEC performance of pure -Fe2O3[9–12].

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Doping with aliovalent elements (e.g., Si, Ti, Ge) and nanostructuring have been shown to

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dramatically affect -Fe2O3activity by enhancing its conductivity and hole diffusion length,

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respectively[1,4,7,13–18]. Among all the doped materials, Si- and Ti-doped -Fe2O3 are the most studied photoanodes. Approaches that systematically correlate the effect of doping on

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conductivity, electronic structure, morphology, and PEC activity rely on titanium doping because

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the Fe2-xTixO3 solid solution offers complete solubility of Ti in the -Fe2O3 structure, with a stable end-member FeTiO3 (ilmenite). For instance, Ti-doped -Fe2O3 have shown increased

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conductivity and improved PEC performance up to Ti 20–30 at.%, while FeTiO3 has showed lower conductivity than that of other composition. In contrast to titanium, silicon is poorly soluble (i.e.,Si 1–2 at.%) in iron oxides at room temperature and ambient pressure showing PEC performance strictly dependent on sample preparation routes[9,16,19].

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In particular, Si-doped -Fe2O3 made by atmospheric pressure chemical vapor deposition (APCVD) currently represents the best performing -Fe2O3 films[9,16,19]. With a doping level of 1–2 at.%, these films showed photocurrents of 2.3 mA/cm2 at 1.23 V vs. the reversible

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hydrogen electrode (RHE) under 1 sun illumination[19]. Others authors prepared Si-doped Fe2O3by spin-coating deposition and showed that the highest photocurrent was obtained for a

0.5% Si content, while increasing Si up to 2% led to a lower activity against all expectations[20]. In addition to the contribution of Si n-type doping, the conductivity of -Fe2O3 photoelectrodes is also strongly influenced by the orientation and by the morphology of nanocrystals[16,21–

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23].The generally accepted conduction mechanism involves small polaron hopping through

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Fe2+/Fe3+, giving rise to an anisotropic conductivity four times more efficient within the (001)

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planes. -Fe2O3 nanostructures preferentially grown along the [110] crystallographic direction

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have shown improved PEC activity[23,24]. Si doping in -Fe2O3 films obtained through

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chemical vapor deposition have been reported to direct the growth along the [110] direction to a

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limited extent[3,16], while physical vapor deposition (PVD) techniques such as sputtering and heteroepitaxial pulsed-laser deposition showed an exquisite control over the crystal

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orientation[24,25].

However, the film conformation imposes a physical constraint on the growth of -Fe2O3

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nanocrystals[16,23,25] which does not allow a reliable investigation of the intimate relationships

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between the morphology/structural changes, and the PEC performance upon doping. Nanoparticles deposition assembly is a powerful approach to study fundamental properties[26] of nanomaterials and their structure–activity relationships, combining the enhanced control over morphology withthe tunability of physical properties of nanocrystals [27–31].

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Here, we present a systematic study on structural, electronic, and magnetic properties, as well as on PEC activity, of Si-doped -Fe2O3 nanostructures prepared by using a hydrothermal method. We investigate the effects of a high level of silicon doping on the structure and physical

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properties of the new Si-doped -Fe2O3 compounds. We elucidate the actual effect of Si incorporation in the-Fe2O3 structure showing that the sole Si content controls the morphology of the-Fe2O3 nanostructures already at Si 1 at.%. Above this doping level we observe that Si induces the transition from nanostructures with ellipsoidal shape to nanowires growing

preferentially along the [110] direction that self-assemble in layered superstructures, similar to

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those formed in many silicates. We present a study on the significant change in morphological,

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structural, magnetic, and electronic properties of Si-doped -Fe2O3 nanostructures and relate

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them to the PEC activity.

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2. Material and methods

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2.1 Synthesis of -Fe2O3 nanostructures and preparation of photoanodes -Fe2O3 nanostructures were prepared as reported elsewhere [32].Briefly, 45 mLof H2O, 45

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mLof ethanol, and 1.05 g of Fe(CH3COO)3 were mixed and heated at 150°C for 15 h in

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autoclave. The doped -Fe2O3 samples were prepared by adding tetramethoxysilane (TMOS) to obtain 5 different Si/Fe at.% concentrations (1%, 5%, 10%, 15%, and 20%). The powders were

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subsequently sintered at 550 °C for 1 h. Hematite electrodes were prepared by electrophoretic deposition of -Fe2O3 nanoparticles on fluorine-doped tin oxide (FTO) coated glass[33]. Typically, 5 mL of the prepared -Fe2O3 solution was mixed with 0.02 g of iodine followed by the addition of 45 mL of acetone under ultrasonication. A 10 V bias was applied for 1 minute. Finally, the obtained films were annealed in air at 500 °C for 1 h.

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2.2 Morphological and physicochemical characterization X-ray powder diffraction (XRPD) patterns were recorded employing the Cu-K radiation at

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room temperature. Qualitative phase analyses of raw patterns were performed using X’Pert High Score Plus (version 2.2c) software (PANalytical BV, Almelo, The Netherlands). Rietveld analysis were performed using the GSAS software suite of programs (see Supplementary material for more details).

Raman scattering was measured in air at room temperature using an NTEGRA - Spectra SPM

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from NT-MDT. A thermo-cooled solid-state laser operating at 532 nm provided excitation. The

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use of a very low laser power (50 W at the sample surface) prevented local heating of the

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samples and annealing effects.

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Transmission electron microscopy (TEM) analysis and related selected area electron diffraction (SAED) were carried out by a ZEISS LIBRA200FE TEM. Energy-dispersive X-ray spectroscopy

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(EDS – Oxford INCA Energy TEM 200) and elemental mapping were collected along with

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HAADF-STEM (high angular annular dark field scanning transmission electron microscopy)

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images.

X-ray photoelectron spectroscopy (XPS) was performed with M-Probe - SSI instrument

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equipped with a monochromatic Al Ksource (1486.6 eV) with a spot size of 200 mm× 750 mm and pass energy of 25 eV, providing a resolution for 0.74 eV.

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UV-vis spectroscopy was performed with a Thermo Scientific Evolution 600 spectrophotometer, equipped with a diffuse reflectance accessory Praying–Mantis sampling kit (Harrick Scientific Products, USA).

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The thermal behavior of the saturation remanent magnetization of -Fe2O3was measured by a Quantum Design MPMS XL-5 SQUID magnetometer.

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2.3 Photoelectrochemical characterization The electrodes were electrochemically characterized in a three-electrode cell; the reference

electrode was the Ag/AgCl electrode, while a high surface area Pt mesh acted as the counter

electrode. The electrode potential (E) was converted to the RHE scaled by the Nernst equation.

Electrochemical impedance spectroscopy (EIS) and PEC measurements were performed with a

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PGSTAT204 Autolab potentiostat and under 1 Sun illumination (using a Lot Oriel Xenon lamp

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with applied an AM1.5G filter) in 1M NaOH. EIS data were collected using a 10 mV amplitude

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perturbation at frequency between 0.01 Hz and 1 MHz. At least three electrodes of each type

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were fabricated and tested. All electrodes showed similar characteristics and representative data

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are reported.

3. Results and discussion

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3.1 Structural and compositional properties

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Fig. 1a schematically summarizes the morphologies of the pure and Si-doped -Fe2O3 synthesized under the considered reaction conditions. After a hydrothermal step at 150°C all the

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samples displayed the typical diffraction pattern of goethite FeOOH (Fig.S1 of the Supplementary material). Following an annealing at 550°C for 1 h they underwent a structural phase transition to -Fe2O3 (Fig.1b). A single crystalline phase characterized the samples before and after the annealing process (Fig.S2, see an example of Rietveld refinement), with no evidence of Si-based spurious phases even at Si 20 at.%.

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Raman spectra confirmed the formation of pure -Fe2O3 for all samples after the calcination step (Figs. 1c and S3). Normal modes of -Fe2O3 at the Brillouin zone center include seven Ramanactive vibrations (2A1g + 5Eg), six IR-active vibrations (2A2u + 4Eu), and five optically silent

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vibrations (2A1u + 3A2g)[34]. Table S1 reports frequencies and assignment of Raman vibration modes commonly detected in FeOOH and -Fe2O3.

Pure and Si-doped -Fe2O3 samples showed the six peaks typical of crystalline -Fe2O3 at 227,

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246, 293, 412, 498, and 610 cm–1. They respectively correspond to A1g(1), Eg(1), unresolved

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Eg(2)-Eg(3) modes originating from the motion of iron cations, and toEg(4), A1g(2) andEg(5)

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modes associated to the motion of oxygen anions[34]. In addition, we detected the IR-active Eu

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mode at ~657 cm–1 in both the pure and all the doped samples. The presence of the Eu mode is consistent not only with the presence of lattice defects[35] and the small grain size[14] but also

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with the incorporation of dopants[13,14,36–39], all leading to the breakdown of symmetry

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selection rules. As the Si content was increased from 1 to 20 at.%, the disorder-activated Eu

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mode progressively intensified and broadened. This suggests an increased structural disorder owing to the incorporation of Si in the nanocrystal lattice. Furthermore, the substitutional nature

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of the doping is supported by the downshift of Eg(4) peak from 407 cm–1 to 403 cm–1 for the 1 and 20 at.% samples[15,21,40].

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Representative TEM micrographs highlighted that pure (Figs.2a, 2d, and S4) and Si1% - Fe2O3 (Figs.2b and 2e) exhibited a similar morphology. The goethite precursor showed an elongated solid shape, whereas after thermal phase transitionα-Fe2O3 hollow ellipsoids were formed for both samples. The fast heating rate (i.e., 30°C/min) used during calcination produced

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a quick dehydration of FeOOH thus favoring the formation of the hollow nanostructures. After the introduction of Si 1 at.%, the crystal dimension changed from a width of 30–40 nm and length of 50 nm for the un-doped -Fe2O3 to a width of 10–15 nm and length of 20–30 nm for

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1%Si - -Fe2O3. The former presented a nanocrystal wall 10 nm thick, whereas in the latter it shrank down to 2–4 nm, a length scale comparable to the hole diffusion length in Fe2O3[3,8,19].

With increasing Si content, a marked shape evolution was evident. At Si ≥5 at.%, -Fe2O3 (Fig.

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2, S4–S6) nanostructures mainly consisted in nanowires and fiber-like building blocks that self-

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assembled through electrostatic interaction into m long bundles, resembling the typical

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morphology of common silicates such as chrysotile. These samples showed a full solid structure,

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before and after the calcination, demonstrating that at dopant concentration higher than 1 at.%, the inclusion of silicon into the -Fe2O3 crystalline habit was the driving force directing the

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growth of nanowires and superstructures with higher hierarchical organization.

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The morphological transition was gradual. Si5% sample comprised a mixture of nanowires and

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ellipsoidal nanoparticles (Fig. S5), whereas Si20% - -Fe2O3 (Fig. S6) contained large aggregates of 10-nm-thick nanowires assembled along their longer lateral dimension. Figs. 2 and

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S4 show representative TEM and HRTEM micrographs of Si10% - -Fe2O3 nanostructures formed by adjacent nanowires having a width of 7–8 nm and some round nanocrystals of few nm

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in size.

Considering the Si-doping range (i.e., 1–20 at.%), one would have expected to observe some spurious phase, especially at concentration higher than 5 at.%. However, STEM-EDS elemental maps analysis (Fig. 2g) for Si20% - -Fe2O3 showed that both Fe and Si presented a

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homogeneous elemental distribution in the entire selected area. The measured Si at.% contents computed as Si/(Si+Fe) are in good agreement with the nominal ones and correspond to 0.9 Si at.%, 4.8 Si at.%, 10.1 Si at.%, and 19.8 at.% for Si1% -, Si5% -, Si10% -, Si20% - -Fe2O3,

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respectively. A throughout analysis of TEM micrographs highlighted the formation of an amorphous Si/SiO2 shell 0.5–1 nm thick around the Si-doped -Fe2O3 nanostructures as clearly shown in Figs. 2c and 2e.

To further investigate the elemental composition of hydrothermally grown -Fe2O3

nanostructures we performed XPS analysis. Survey analysis reported in Figs.S7a and S7b refers

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to pure and Si20% - -Fe2O3, the most representative samples. They revealed the following

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surface Si at.% concentration: 1.6% (Si1% - -Fe2O3), 7.6% (Si5% - -Fe2O3), 9.6% (Si10% -

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-Fe2O3), and 13.1% (Si20% - -Fe2O3). Results were different once taken into account the

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Si/(Si+Fe)at.% that was 23, 39, 47, and 60% for Si1% -, Si5% -, Si10% -, and Si20% - -Fe2O3,

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respectively. This increase highlighted a progressive Si surface enrichment of the nanostructures

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at both low and high Si doping at.% providing a complementary information to EDX analysis. High resolution XPS analysis of the Fe 3p region in pure -Fe2O3 nanoparticles showed the

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typical line shape associated to -Fe2O3[41,42], with the only presence of the Fe3+ component

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(Fig.S7c). In contrast, the spectra of all Si-doped samples suggested the presence of an additional component (see for example the line shape of Si20% sample reported in Fig.S7d), consistent

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with the presence of Fe2+[42]. The presence of Fe2+ is the direct signature of substitutional doping. According to n-type doping behavior expected for Si4+, the Fe2+ species must account for an increase of electron donor concentration upon doping. Table S2 reports the donor concentration (ND) extracted from Mott–Schottky measurements (Fig.S8) obtained by using EIS in dark conditions. ND values grew from 3.5 × 1019 for the bare

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-Fe2O3 to 9 × 1019 for the Si20%--Fe2O3 (Table S2) reaching a plateau at high doping concentration, which is consistent with the Si/SiO2 surface segregation shown by TEM and XPS analysis. The introduction of Si did not significantly change the flat band potential of the -

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Fe2O3 that resulted in values between 0.57 and 0.62 V (Table S3).Assuming a pure substitutional doping for Fe3+ in -Fe2O3, the charge introduced by Si4+ can be accommodated by the

generation of Fe2+ to achieve the electroneutrality. In this view the valence reduction is featured at the Fe3+ sites by negatively charged point defects, which fully account for the observed

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increase of electron donor concentration upon doping.

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To investigate the effect of Si doping into the growth of our -Fe2O3 nanostructures, we analyzed

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more deeply the XRPD patterns. Figure 3a shows the doping evolution of lattice parameters.

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Increasing Si concentration causes a drop of a axis value and a lengthening of c axis which drives

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a remarkable increase of hexagonal strain  = c/a saturating above 10 % Si (see inset to Fig.4b).

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This structural evolution upon doping can be fully consistent with a substitutional doping mechanism and it can be supported by ionic radii (ir) mismatch between Fe3+ (ir = 0.65 Å) and

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induced Fe2+ ( irFe(II) = 0.78Å) in 6-fold coordination featuring a sort of solubility limit at Si > 5%. In Fig.3b, selected portions of the XRPD patterns highlight the doping-dependent evolution of

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(104) and (110) peaks. Pure -Fe2O3 and Si-doped -Fe2O3 displayed a marked difference in the

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extracted position and full width at half maximum (FWHM) (see Fig.3c and 3d). We used them to estimate the average coherence length (CL/ nm) of crystallites parallel to the (104) and (110) crystallographic planes (Fig.3e). For the pure -Fe2O3, the CL related to both peaks almost coincides providing evidence of a homogenous distribution of spherical-shaped particles, in agreement with TEM images (Fig. 2a). Increasing the Si content up to 5 at.% the peak

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broadening becomes (hkl)-dependent and the (104) CL dramatically decreased, whereas the (110) CL remained nearly constant (Fig.3e). In Fig.3e we incorporated the width of the fiber-like shaped crystallites estimated through the

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analysis of TEM micrographs for 0%, 10%, and 20% Si-doped samples. The width of the nanostructures matched very well with the CL related to crystallites parallel to the (104) crystallographic planes, confirming the interplay between the anisotropic peak broadening and the morphological transition of crystallites[43]. Interestingly, the scattering related to the (110) and the (104) planes depends on the preferential growth of crystalline (001) planes perpendicular to

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the specimen plane[23]. The observed peak broadening of the (104) peak with increasing Si

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content can be viewed as a sort of preferential elongation of the crystal along the [110] crystallite

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direction. This corroborates with the narrower shape profile of (110) peaks, which display a weak

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doping dependence on CL parameter in the whole studied compositional range (see Fig.3d). In other words, the progressive increase of the Si content produced a morphological transition from

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hollow nanostructures to nanowires that self-assembled in -Fe2O3 bundles. The nanowires grew

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along the [110] crystallographic direction thus showing very developed (001) planes.

3.2 Light absorption

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UV-vis spectra of synthesized -Fe2O3 nanostructures (Fig.3e) show that the threshold absorption is located around 600 nm for all samples. -Fe2O3 absorption is dominated by a

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strong transition in the region 400–600 nm, which is assigned to the pair excitation processes 6

A1(6S) + 6A1(6S) to 4T1(4G) + 4T1(4G) at 485–550 nm, and is overlapped by the contributions of

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A1(6S) to 4E, 4A1(4G) ligand field transitions at 430 nm and the ligand to metal charge transfer

transition band tail[43]. The double exciton process 6A1(6S) + 6A1(6S) to 4T1(4G) + 4T1(4G) yields

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the strongest absorption band at around 535 nm and it is primarily responsible for the red color of -Fe2O3. Interestingly, increasing the Si content in our -Fe2O3 nanostructures produced a decrease in intensity of the absorption related to the double exciton process (Fig.3e). The most

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notable change in absorption was observed passing from pure -Fe2O3 to Si1% --Fe2O3, while at higher loading the absorption decrease is weaker with absorption spectra of Si10% and Si20% --Fe2O3 that show a similar behavior.

3.3 Magnetic properties

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Fig.S9 and Table S4 report the remanent magnetization (Mr) of our-Fe2O3 nanoparticles. The

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most characteristic magnetic property of bulk -Fe2O3 is the Morin transition between the strictly

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antiferromagnetic (AFM) low-temperature phase (with AFM axis parallel to the crystal c axis)

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and the high-temperature phase, which features a slight canting of the spins out of the AFM axis

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lying in the basal ab plane[44,45]. After application of a magnetic field at room temperature,

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bulk -Fe2O3 retains a weak Mr, which on cooling undergoes a rapid decrease at the Morin transition temperature TM (262 K in pure bulk -Fe2O3). On the other hand, the magnetic

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behavior of bulk-Fe2O3 displays two prominent features. First, upon re-heating through TM a

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partial recovery of Mr occurs (memory effect). Second, Mr does not vanish below TM. Some magnetization remains that is constant down to very low temperature which can be attributed to

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lattice defects (Mdefect)[46].Thus, Mr comprises two parts arising from lattice defects and canted AFM spins, Mr(T) = Mdefect + MCAFM(T). The thermal behavior of Mr of pure and Si-doped (1%, 5%, 10%, 20%) -Fe2O3 nanostructures naturally leads to the subdivision of the studied samples in two groups (Fig.S9), namely, Group I and II. Group I comprises pure and Si1%- -Fe2O3 and displays Mr(T) curves similar to that of

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bulk -Fe2O3, with the Morin transition and memory effect clearly discernible (Fig.S9a). Group II (Si5%, Si10%, and Si20%) displays much lower Mr and almost flat Mr(T) curves in comparison with Group I (Fig.S9b). A detailed discussion of magnetic measurements is reported

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in the Supplementary material. Briefly, pure and Si1% - -Fe2O3 nanostructures magnetically behave as nanostructured -Fe2O3 and the effect of Si doping seems to be limited to a decrease of the remanent magnetization related to the canted AFM. In contrast, the remanent

magnetization of highly Si-doped -Fe2O3 nanostructures is primarily the defect magnetization, which is scarcely affected in magnitude by Si doping. However, a small fraction of Mr still

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undergoes Morin transition. Such fraction decreases when the Si doping is increased. It could be

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attributed to a few particles with low doping level or low-doped regions within nanorods with

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higher doping level.

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3.4 Photoelectrochemical activity and electrical properties

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Fig.4a shows the photocurrent density (J) values extracted at 1.23 V vs RHE under 1 Sun illumination for all the -Fe2O3 samples (see Figs.S10 and S11). The pure -Fe2O3 produced 62

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A/cm2, while the most active photoanode, i.e., Si1% - -Fe2O3, showed a 20% photocurrent enhancement reaching 72.5 A/cm2. A further increase of Si content produced a monotonic

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decrease of the photocurrent that for Si ≥5 at.% resulting in a lower PEC activity than that of

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pure -Fe2O3. These photocurrent values are comparable with the PEC activity reported by Can Li’s group for similar nanoparticle-based studies[27–30]. Their works have shown how PEC activity of sub-micrometer particle photoanodes correlate well with microscopic physicochemical properties of the crystals[27].

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Photocurrent curves (Fig.S11) revealed a ~30 mV anticipation in onset potential for Si1%, while all the other samples showed similar values as reported for pure -Fe2O3. Interestingly, Si15% and 20% presented increased photocurrent with respect to pristine hematite in the potential range

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0.7–1.0 V, while at higher potential the performance were strongly damped. The increased performance at low potential might be tentatively assigned to oxidation peak related to Si-Fe species present in the samples at high Si content.

Fig. 4a shows the average CL of crystallites along with the photocurrent data. The latter was

depressed with the decrease of nanostructure CL, following the width contraction of fiber-like

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shaped Si-doped -Fe2O3. Since the electron conduction in the [110] is four orders of magnitude

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higher than that on the axial direction, one would expect that the growth of elongated

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nanostructures with well-developed (001) planes resulted in photoanodes with high PEC activity.

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To understand the structure–activity relationship in highly Si-doped -Fe2O3 we consider the

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role of the morphology and chemical composition of the nanostructures. Si1% - -Fe2O3, the

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most efficient sample, was the only one maintaining the hollow nanoparticle morphology and a similar crystallite orientation than that of pure -Fe2O3. The nanocrystal wall shrank down from

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10 nm in pure -Fe2O3 to 2–4 nm in Si1% - -Fe2O3, thus reaching a length-scale comparable to the hole diffusion length and, hence, justifying the higher PEC activity. On the other hand, TEM

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images evidenced that nanowires formed upon insertion of Si ≥5 at.% had lateral width of about

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8–10 nm (Fig. 2), a value still compatible with the diffusion length of photogenerated holes. However, the increase of the Si content produced the assembly of nanowires in fiber-like bundles containing tens of units aligned along their lateral dimensions, and stacked together to form large lamellar aggregates. The random orientation of these structural entities broke the preferential exposition of (001) planes accounting for the low PEC activity observed for the highly doped α-

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Fe2O3 nanostructures. The morphological transition from nanoparticles to assembled nanofibers and the consequent surface/volume ratio decrease as well as the increased Si/SiO2 surface enrichment detected by XPS are additional evidences depressing the PEC activity of highly

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doped α-Fe2O3. To confirm our hypothesis, we studied the electrical transport properties of the photoanodes

through EIS at 1.23 V under light conditions. A single arc depicted the impedance behavior of all the electrodes and thus an RC circuit was used as equivalent circuit to fit the data. In these

conditions, the arcs of the Nyquist plot represent the bulk properties of the electrodes: Rbulk (bulk

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transport resistance) and Cbulk (capacitance of the bulk). The dimensions of the arcs reflected the

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electrode properties: the bigger the arcs the higher the resistance. This resistance can be

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correlated both to the internal resistance of the -Fe2O3 nanoparticles or to the charge transfer

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resistance through -Fe2O3 nanoparticles. Fig.4b shows that, in comparison to bare -Fe2O3, a concentration of Si 1 at.% produced a smaller arc, whereas increasing the Si-concentration the

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arcs become larger. Only the presence of Si 1 at.% increased the conduction of the -Fe2O3,

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whereas higher Si concentration increased considerably the transport resistance. Mott–Schottky

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(Table S2 and Fig.S8) analysis shows a consistent increase of donor concentration along with Sidoping and then a better conduction of -Fe2O3. Therefore, the EIS results indirectly suggest that

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the photocurrent is mainly limited by the high charge transfer resistance through adjacent Fe2O3 nanoparticles. The observed trend can be ascribable to the different shape of the

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nanoparticles, pure and Si 1 at.% -Fe2O3 were formed by smaller nanoparticles thus allowing, by means of electrophoretic deposition, to obtain a more dense assembled electrode showing a better electrical conduction. Conversely, for Si concentration ≥5at.% the aggregation of nanowires did not allow optimal nanostructures packing into the electrode, thus reducing the

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electron mobility through the nanostructures, also influenced by the concomitant Si surface enrichment.

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4. Conclusions We showed that silicon directs the growth of -Fe2O3 nanostructures and it templates the

morphological transition from hollow nanoparticles to nanowires already at Si 1 at.%. Increasing Si above 5 at.%, these entities assembled into fiber-like bundles similar to the morphology found in common silicates. Importantly, we showed that even Si monotonically induced the increase of

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donor density accounting for the generation of Fe2+ species upon substitutional doping. The

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morphological transition was accompanied by a preferential nanocrystal growth in the [110]

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direction, so as the development of extended (001) crystallographic planes was observed.PEC

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water splitting measurements on pure and Si-doped -Fe2O3 photoanodes showed that the Si doping was effective in improving PEC activity only in the case of Si1% sample, which showed

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a 20% photocurrent enhancement in comparison with pure -Fe2O3. On the contrary, the -

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Fe2O3 containing Si content higher than 5 at.% presented lower PEC activity. The morphology

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transition and the increasing Si surface enrichment of the nanostructures at high Si content

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produced a lower electrode electrical conductivity thus depressing their performance.

ACKNOWLEDGMENTS

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The authors acknowledge the support by the project no. CZ.02.1.01/0.0/0.0/15_003/0000416 and the project 8E15B009 of the Ministry of Education, Youth and Sports of the Czech Republic.

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Fig. 1. (a) Scheme of the hydrothermal method used to synthesize pure and Si-doped -Fe2O3

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with Si content in the range 1–20 at.%. (b) XRPD patterns and (c) Raman spectra of pure and

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Si20% - -Fe2O3.

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Fig. 2.TEM and HRTEM micrographs for (a,d) pure, (b,e) Si 1 at.%, and (c,f) Si 10 at.% - -

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Fe2O3, respectively. The insets of (a), (b), and (c) represent low magnification TEM micrographs of the samples. The inset of (f) is the FFT pattern relative to the HRTEM image for Si10% - -

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Fe2O3 (g) STEM-EDS elemental maps for Fe and Si on the selected area for Si20% - -Fe2O3.

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Fig.3. (a) Doping evolution of the refined cell parameters. Squares and circles are hexagonal c and a axis, respectively of hematite. Inset shows the doping evolution of hexagonal strain order parameter (). (b) Selected XRPD pattern regions showing (104) and (110) reflections for pure Si-

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doped -Fe2O3powders. (c) Position and (d) full width at half maximum (FWHM) for (104) and (110) as a function of the Si content (at.%). (e) Coherence length (CL) of -Fe2O3 crystallites along (110) and (104) direction (blue and red line, respectively) as estimated from (c) and (d). The -Fe2O3 crystallite widths obtained from TEM micrographs analysis are also shown (green). (e)

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UV-vis spectra of pure and all Si-doped hematite samples.

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Fig.4. Electrochemical measurements on pure and Si-doped -Fe2O3 samples in 1M NaOH and under 1 Sun AM1.5G illumination. (a) Photocurrent density (J) (blue squares) and average

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coherence length (CL) of crystallite along the (104) and (110) direction (red circles) as a function of the Si content at 1.23 V vs. RHE. (b) Nyquist plot for all investigated photoanodes carried out at 1.23 V.

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