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Ultrathin Buffer Layers of SnO by Atomic Layer Deposition: Perfect Blocking Function and Thermal Stability Ladislav Kavan, Ludmilla Steier, and Michael Grätzel J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.6b09965 • Publication Date (Web): 13 Dec 2016 Downloaded from http://pubs.acs.org on December 14, 2016

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The Journal of Physical Chemistry

Ultrathin Buffer Layers of SnO2 by Atomic Layer Deposition: Perfect Blocking Function and Thermal Stability Ladislav Kavan1,2*‡, Ludmilla Steier2, ‡ and Michael Graetzel2 1

J. Heyrovský Institute of Physical Chemistry, v.v.i., Academy of Sciences of the Czech Re-

public, Dolejškova 3, CZ-18223 Prague 8, Czech Republic 2

Laboratory of Photonics and Interfaces, Institute of Chemical Sciences and Engineering,

Swiss Federal Institute of Technology, CH-1015 Lausanne, Switzerland ‡

LK and LS contributed equally to the work.

E-mail: [email protected]

ABSTRACT This study pinpoints the advantages of ultrathin electron selective layers (ESL) of SnO2 fabricated by atomic layer deposition (ALD). These layers recently caught attention in planar perovskite solar cells and appear as powerful alternatives to other oxides such as TiO2. Here, we carry out a thorough characterization of the nature of these ultrathin ALD SnO2 layers providing a novel physical insight for the design of various photoelectrodes in perovskite and dyesensitized solar cells and in photoelectrochemical water-splitting. We use a combination of cyclic voltammetry, electrochemical impedance spectroscopy, Hall measurements, X-ray photoelectron spectroscopy, AFM and electron microscopy to analyze the blocking behavior and energetics of as-deposited (low-temperature) and also calcined ALD SnO2 layers. First, we find that the low-temperature ALD-grown SnO2 layers are amorphous and perfectly pinholefree for thicknesses down to 2 nm. This exceptional blocking behavior of thin ALD SnO2 layers allows photoelectrode designs with even thinner electron selective layers thus potentially minimizing resistance losses. The compact nature and blocking function of thin SnO2 films is not perturbed by annealing at 450 oC, which is a significant benefit compared to other amorphous ALD oxides. Further on, we show that amorphous and crystalline ALD SnO2 films substantially differ in their flatband (and conduction band) positions – a finding to be taken into account when considering band alignment engineering in solar devices using these highquality blocking layers.

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1. INTRODUCTION The steep progress in the field of perovskite solar cells1,2 renewed the interest in thin layers of transparent oxide semiconductors, acting as electron-selective contacts. The electron collecting electrodes have been developed earlier as buffer layers for the recombinationblocking in a classical liquid-junction dye-sensitized solar cell (DSC)3 with iodine-based mediators4,5 and Co-based mediators.6 Here, the buffer layer prevents the back electron-transfer from the electron-collecting substrate (F-doped SnO2; FTO) underneath the TiO2 photoanode to the electrolyte (I3- or CoL33+ L=ligand, e.g. bipyridine) which is an undesired parasitic process in these solar cells.3 Nevertheless, the presence of blocking layers is not mandatory for the proper function of the liquid-junction DSCs4-6 since it only moderately improves their performance metrics, i.e. open-circuit voltage, fill factor and short circuit photocurrent. On the other hand, the buffer layer is pivotal for a solid-state DSC, where it inhibits the back electron-transfer from the FTO substrate to the hole-conductor, like spiro-OMeTAD.7 The same holds for perovskite solar cells, regardless whether they operate with or without the auxiliary hole-conductor.1,2,8 Up to now, these blocking layers were prepared almost exclusively from TiO2, with sporadic reports on other oxides, such as Al2O3.9,10 The buffer layer is requested to be transparent enough and pinhole-free to fully block charge carrier recombination between the electrode contact and the photon absorber. The most popular synthetic technique has been spray pyrolysis of titanium diisopropoxy bis(acetylacetonate).2,4,5,7,11 Other possible techniques are sol-gel12-14, spin-coating15, DC-magnetron sputtering16, electrochemical deposition17,18 and atomic layer deposition (ALD).6,18-27 The last method is attractive not only for solar cells,6,21,23 but also for photoelectrochemical water splitting, where the ALD-made titania layers prevent corrosion of certain electrode materials such as Cu2O or silicon.19,20,24-27 A comparative study

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of TiO2 buffer layers prepared by spray pyrolysis, ALD and electrochemical deposition on FTO substrates confirms that the latter two methods provide high-quality pinhole-free films, but the blocking function of titania is strongly impaired by calcination at 450-500oC by cracks and other defects.18 The calcination is a standard procedure for subsequent deposition of mesoporous oxide layers on top of the blocking layer to enhance the photo-conversion efficiency of solar cells.3,28 The application of other oxide coatings for buffer layers on FTO is less frequently encountered in the literature. Baena et al.29 demonstrated recently that thin ALD-SnO2 films work

surprisingly

well

in

planar

solar

cells

with

mixed

perovskite,

(CH3NH3PbBr3)0.15(NH2CH=NH2PbI3)0.85. The effect was attributed to a barrier-free alignment of the conduction band (CB) edge of amorphous ALD-SnO2 and that of the mixed perovskite. On the other hand, the interface of TiO2/mixed-perovskite shows a bandmisalignment of 0.3 eV. The studies of conduction-band energetics of single-crystal n-SnO2 are dating back to the works of Wrighton et al.30,31 The flatband potential (EFB) of SnO2 was found to be by ca. 0.3 V positive to that of TiO2 (rutile), which means that the conduction band of SnO2 is downshifted by roughly the same value (if we neglect the offset of positions of the Fermi level (EFB) from the CB edge and possible small deviations of the isoelectric points of TiO2/SnO2).30,32,33 A similar flatband potential was reported by Metikos-Hutkovic et al.34 for polycrystalline n-SnO2: EFB = -0.19 V vs. SCE (saturated calomel electrode) in pH=6, translating to 0.41 V vs. RHE (reversible hydrogen electrode). The advantage of using the RHE as a reference is that it avoids the Nernstian pH-dependence of the EFB values of oxide semiconductors.32 Comparable values for polycrystalline SnO2 were found by others: 0.37 VRHE35 and

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0.50 VRHE36, although sometimes also strongly deviating data are encountered in the literature.37 In Baena’s work, 29 ALD layers of TiO2 and SnO2 showed a conduction band offset of 0.25-0.27 eV. This downshift is equivalent to EFB between 0.41 to 0.43 VRHE, taking the anatase (101) as a reference.32 The matching of these values with the electrochemically determined EFB30,34,35 is actually surprising with respect to the recent work of Deák et al.32,38 showing that the CB edge of a semiconductor electrode exposed to vacuum or to electrochemical medium is different for fundamental reasons. Furthermore, EFB of anatase is known to be enhanced by 0.2-0.3 V for quasi-amorphous (ALD-made) titania against crystalline anatase.18 In addition to the use as electron selective layer in perovskite cells,29 SnO2 was reported to be a promising material for a counterelectrode in the traditional DSCs with an iodinemediator.39 Here, a greatly enhanced electrocatalytic activity for the I3- reduction was observed upon calcination of SnO2 in N2 atmosphere.39 Furthermore, nanocrystalline SnO2 was also tested as the photoanode in DSCs. Compared to TiO2, tin oxide is favored by larger electron mobility (≈10-100 cm2V-1s-1 vs. ≈0.1-1 cm2V-1s-1)40,41 and by a larger band gap (3.5-3.6 eV). The latter minimizes the formation of holes that would cause oxidative degradation of the DSC during long-term operation. In turn, the downshifted conduction band of SnO2 is also responsible for a smaller open-circuit voltage (VOC) in addition to smaller dye adsorption ability and faster electron recombination.35 However, no open circuit potential loss was observed in the planar SnO2-perovskite solar cells.29 Another area of application of ALD-made SnO2 thin films is in surface passivation of carbon-based electrocatalysts42 and graphene-supported anodes for Li-ion batteries.43 In spite of these numerous applications of SnO2 thin films, an electrochemical evaluation of the blocking function of SnO2 is still missing. Here we report a detailed analysis of the

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low-temperature ALD-grown SnO2 layers. We probe the quality of the blocking SnO2 layers of as-deposited as well as annealed (at 450 °C) ALD-SnO2 films of different thicknesses by cyclic voltammetry (CV) and X-ray photoelectron spectroscopy. We further elucidate major differences in band energetics by impedance spectroscopy and Hall measurements and propose a clearer band alignment picture based on the combination of our results and calculations. 2. RESULTS AND DISCUSSION The buffer layer in solar cells has to fully block charge carrier recombination between the electrode contact (FTO) and the photon absorber (in planar perovskite cells) or the electrolyte (in liquid-junction DSCs) or the hole-conductor (in solid-state DSCs). Hence, the analysis of pinhole-defects in these blocking layers is of great interest. A sensitive technique to detect pinholes is cyclic voltammetry, which even allows quantifying the pinhole concentration18 from the peak current according to equation (1): jp Au = A0 j o , FTO

(1)

Au is the effective pinhole area, A0 the total electrode area, jp the peak current density at the actual blocking electrode and jp,FTO the peak current density at the clean FTO electrode. This equation is valid assuming linear diffusion towards the FTO/electrolyte solution interface where the voltammetric peak current, Ip is given by the Randles-Sevcik equation: Ip = k·n3/2·A·c·D1/2·v1/2

(2)

with k being a constant (k = 2.69 ·105 C mol-1 V-1/2), n the number of electrons appearing in the half-reaction for the redox couple, A the electrode area, D the diffusion coefficient and ν the scan rate. (Specific diffusion fields might develop around pinholes in real electrodes, e.g.

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the microelectrode-like pinholes can show spherical diffusion and sigmoidal-shaped voltammogram,44,45 but these fields may also merge to form a single planar diffusion layer.) Our earlier work18 has confirmed the suitability of the [Fe(CN)6]3-/4- redox couple for probing the pinhole concentration in low-temperature ALD TiO2. Hence, we chose the same conditions to probe our ALD-SnO2, that has recently been shown to have an energetic offset of only 0.3 V with respect to the ALD-TiO2.29 Fig. 1 shows cyclic voltammograms of this redox couple at the low-temperature (118 °C) ALD SnO2 films of different thicknesses (1100 nm) on FTO. The low-temperature ALD-SnO2 is expected to form a rectifying contact with the electrolyte solution due to the formation of a depletion layer and corresponding barrier to electron transfer. Hence, we expect an absence of the oxidation peak in the cyclic voltammograms for pinhole-free ALD-SnO2 films. Fig. 1 shows cyclic voltammograms of this redox couple at the low-temperature (118 °C) ALD SnO2 films of different thicknesses (1100 nm) on FTO. The applicability of Eq. 1 and 2 is corroborated by (i) the scan-rate (v1/2) dependence of jp (not shown) and (ii) the Nernstian peaks-shaped voltammogram18,44 for our 1-nm ALD-SnO2 film (Fig. 1).

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Figure 1: Cyclic voltammograms at bare FTO electrode (dashed line) and that covered by SnO2 layers made by ALD. The layer thickness (in nm) is labeled in annotations. Scan rate: 50 mV/s. The electrolyte solution is 1 mM K4Fe(CN)6 + 1 mM K3Fe(CN)6 in aqueous 0.5 M KCl, pH 2.5. The voltammograms (except ‘FTO’ and ‘100’) are offset for clarity.

From 2 nm onwards our low-temperature ALD-SnO2 films fully suppress the oxidation of ferrocyanide even at potentials as high as 1.4 VRHE and thus, show perfect blocking behavior. Solely, the 1 nm thick SnO2 exhibits pinholes, which expose part of the FTO surface to the electrolyte solution giving Au/A0 = 59%. However, with respect to ALD-made TiO2 films (with Au/A0 of 79% or 87%, depending on the synthetic conditions),18 the blocking function of ALD SnO2 compares favorably. Indeed, the titania films actually required thicknesses of at least 3-5 nm for comparable blocking function.18 The excellent blocking of ultrathin SnO2 films on FTO might be due to the close structural similarity of SnO2 and FTO, in contrast to the structurally incompatible ALD-titania (anatase) and FTO (rutile).

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X-ray photoelectron spectra (XPS) of a 15 nm thick Au-supported ALD SnO2 film show no signal of the underlying Au (Fig. S1, Supporting Info) which confirms the compact (pinhole-free) nature of our ALD SnO2 films too. In fact, our XPS data suggest the ALD SnO2 films remain pinhole-free even after an annealing treatment at 450 °C. The annealed SnO2 films are of high interest as they could potentially replace the compact TiO2 blocking layer in mesoporous dye-sensitized and perovskite solar cells where the 450 °C calcination treatment is part of the standard protocol for the fabrication of the mesoporous layer. Corresponding CV measurements on these Au-supported SnO2 films (Fig. S2, Supporting Info) show excellent blocking of ferrocyanide oxidation for the as-deposited ALD SnO2 films, which is in good agreement with our results on FTO substrates (Fig. 1), indicating that ALD of SnO2 is not sensitive to the substrate surface (FTO or Au) since all films behave very similarly. Surprisingly, cyclic voltammograms of the [Fe(CN)6]3-/4- couple on the 10 and 100 nm thick annealed SnO2 films on FTO (Fig. 2) and on Au (Fig. S2, Supporting Info) strongly resemble those on pure FTO. From Eq. (1), the Au/A0 values would correspond to 93% or 88%, for the calcined 10 nm or 100 nm SnO2 films, respectively. However, it is very unlikely the [Fe(CN)6]3-/4- electrochemistry occurs on pinholes, bearing in mind the unrealistically large Au/A0 values and the absence of the signal of the supporting material in XPS (Fig. S1 Supporting Info).

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Figure 2: Cyclic voltammograms at bare FTO electrode (dashed line) and that covered by SnO2 layers made by ALD (black full line). Red line is for the same sample calcined at 450 oC in air. The layer thickness (in nm) is labeled in annotations. Scan rate: 50 mV/s. The electrolyte solution is 1 mM K4Fe(CN)6 + 1 mM K3Fe(CN)6 in aqueous 0.5 M KCl, pH 2.5.

We carried out top-view scanning electron microscopy (SEM) and cross-sectional transmission electron microscopy (TEM) measurements to investigate the morphology of the ALD SnO2 films (Fig. 3). The annealing induces a change from a smooth (Fig. 3a) to a more rough surface (Fig. 3b) that is caused by the crystallization of the dense amorphous SnO2 film (Fig. 3c) at 450 °C yielding a dense nanoparticulate film with particles on the order of 10 nm (Fig. 3d). From now on, we designate the pristine films as ‘amorphous’ (a-SnO2) and the calcined films as ‘crystalline’ (c-SnO2). Electron microscopy obviously excludes that ≈90% of the FTO electrode area is exposed to the electrolyte in c-SnO2 even though the film is composed from nanoparticles. There are two possible explanations for the observed differences in electrochemical behavior of a/c-SnO2 (Fig 1. vs. Fig. 2): (i) The calcination induces major energetic differences between the ALD-made a-SnO2 and c-SnO2 films, or (ii) the calcination increases the doping in c-SnO2 films such that they become degenerate. The latter is however, improbable for our layers which have quite low doping even after calcination and show rectifying behavior for redox probes with a more positive potential (cf. Table 1 and discussion be-

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low). We carried out electrochemical impedance spectroscopy and Hall-effect studies for clarification of the rectifying nature of the SnO2/[Fe(CN)6]3-/4- contact.

(a)

SnO2 As-deposited

(b)

200 nm (c)

SnO2 As-deposited

FTO

SnO2 Annealed

200 nm (d)

SnO2 Annealed

FTO 5 nm

15 nm

Figure 3 Top: Scanning electron micrographs in top view of a) as-deposited (at 118 °C) and b) annealed (at 450 °C) ALD SnO2 films on FTO substrates. Bottom: Transmission electron micrographs in cross-sectional view of c) as-deposited and d) annealed ALD SnO2 films on FTO. The flatband potential of our ALD SnO2 layers was measured using electrochemical impedance spectroscopy (EIS) in the dark. The impedance spectra were fitted to a Randlestype equivalent circuit, in which the space charge capacitance is modeled by a constant phase

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element (CPE) to account for non-ideal capacitive behavior (see details in Experimental Section).18 Fig. 4 shows the Mott-Schottky analysis of the EIS measurements, obeying the equation46:

   k T k T 1  2 1 2  E − EFB − B  + 2 =   E + EH − EFB − B  =  2 e  CH e  C  eε 0ε r N D   eε 0ε r N D 

(3)

where e is the electron charge, ε0 is the permittivity of free space, εr is the dielectric constant of the semiconductor (εr ≈ 10 for SnO2)34,36, ND the number of ionized donors per unit volume,

E the applied voltage, EH is the potential drop across the Helmholtz layer, EFB is the flat band potential, kB is Boltzmann’s constant, and T is the temperature. The capacitance values are referenced to the real area of the electrode. To this purpose, the projected area was corrected by a roughness factor (Rf) which was found experimentally from atomic force microscopy (AFM) images (Fig. S3, Supporting Info). The Rf values of our films were from 1.2 to 1.4; in good agreement a value reported earlier for the same FTO, but coated with ALD-hematite (8nm Rf= 1.26).47 The flatband potential is given by: E FB = E 0 + E H −

k BT e

(4)

where E0 is the potential at the intercept (for C-2 = 0). The double-layer correction further includes the contributions from diffuse (Gouy) layer, surface states and specifically adsorbed ions.33,48 This correction may further include a thin dielectric layer at the surface, which takes part of the voltage drop.46,48 However, in accord with a general practice, all these additional factors are simplified in terms of a constant Helmholtz capacitance (Eqs 3,4).48 Fig. 4 shows Mott-Schottky plots of a-SnO2 (left chart) and c-SnO2 (right chart) on FTO with different SnO2 layer thicknesses (3-15 nm).

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Figure 4: Mott-Schottky plot of a bare FTO electrode and that covered by SnO2 layers made by ALD. Left chart: as received samples; Right chart: the same samples calcined at 450 oC in air. The layer thickness (in nm) is labeled in annotations. The electrolyte solution is aqueous 0.5 M KCl, pH 2.5. Potentials were measured with Ag/AgCl reference electrode, but are recalculated against RHE for easier comparison.

In the case of a-SnO2, C-2 is almost independent of the film thickness. Assuming a Helmholtz capacitance of CH ≈ 50 µF/cm2 (as in Refs.34,36), an average flatband potential of 0.2 VRHE (in pH 2.5) could be calculated with Eq. 4 for the ALD a-SnO2 films. This value drastically shifts by 0.5 V for the annealed c-SnO2 thin films. Hence, we hypothesized that the [Fe(CN)6]3-/4couple was not suitable to probe pinholes in ALD c-SnO2 films. To guarantee a rectifying contact with the c-SnO2 for the pinhole analysis, we chose the [Ru(bpy)3]3+/2+ couple, that has a more positive redox potential of 1.38 VRHE. Indeed, Fig. 5 shows that the oxidation of [Ru(bpy)3]2+ is successfully suppressed on both, the a-SnO2 and the c-SnO2, thereby giving a direct evidence that all our ALD SnO2 films are pinhole-free. These findings disclose a significant benefit of ALD-SnO2 against ALD-TiO2 films: while titania is spontaneously cracking upon heat treatment18, the SnO2 films are perfectly stable.

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Figure 5: Cyclic voltammograms at bare FTO electrode (dashed line) and that covered by SnO2 layers made by ALD (black full line). Red line is for the same sample calcined at 450 oC in air. The layer thickness (in nm) is labeled in annotations. Scan rate: 50 mV/s. The electrolyte solution is 0.1 mM Ru(bpy)3Cl2 in aqueous 0.5 M KCl, pH 2.5.

Knowing that our ALD SnO2 layers are pinhole free, we selected the dimethylviologen, MV2+/+ redox couple with a redox potential of Eredox = -0.65 V vs. Ag/AgCl (sat’d KCl) that is independent of pH,18 to pinpoint the conduction band position of ALD a-SnO2. As the pH dependence of flatband potentials is Nernstian, we chose two pH values that set the EFB of a-SnO2 slightly more positive (at pH 10.6) or slightly more negative (at pH 11.9) versus the MV2+/+ redox potential. Fig. 6a compares the CVs of the a-SnO2/dimethylviologen system at pH 10.6 and 11.9. In both cases, MV+ is oxidized. In pH 10.6, the reduction peak position is coincident with the one on FTO, and both peak currents are symmetrical. This suggests, that here an ohmic contact is present. At pH 11.9 a weak rectifying behavior can be identified showing a smaller peak current of MV+ oxidation and a slight negative shift of the reduction peak of MV2+ that confirms nicely our determined flatband potentials of a-SnO2 from MottSchottky plots. This analysis indicates that the conduction band of a-SnO2 should be positioned at or lower than the MV2+/+ redox potential in pH 10.6 as shown in Fig 6b, which trans-

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lates to ECB = 0.17 VRHE and ECB-EFB = 0.02 V in a-SnO2 in agreement with the ohmic behavior seen at pH 10.6. a)

b)

Figure 6: a) Cyclic voltammograms at bare FTO electrode (dashed line) and that covered by 100 nm thick, amorphous SnO2 layer made by ALD. Scan rate: 50 mV/s. The electrolyte solution is 1 mM dimethylviologen (MV2+) in aqueous 0.5 M KCl with two different pH values (10.6 and 11.9, respectively). The expected positions of flatband potential are indicated by arrows for each pH value. The redox potential of dimethylviologen is E0 = -0.65 V vs. Ag/AgCl (sat’d KCl) independent of pH (dashed arrow). b) Estimation of conduction band position of a-SnO2 from the CVs in (a).

As the EIS measurements of flatband potentials are less perturbed by the depletion layer (which even disappears at EFB), the extracted flatband potentials herein can be considered correct for the c-SnO2 films and in agreement with literature.30,31,34,36 (To the best of our knowledge, there is no comparable earlier EIS measurement of EFB on amorphous ultrathin SnO2.) Nevertheless, our value for a-SnO2 is near the position assumed for the band alignment with perovskites29 as well as it is near the onset potential of ferricyanide reduction at this electrode (Fig. 1).

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Assuming a dielectric constant of our ALD SnO2 films of εr = 10,34,36 the MottSchottky analysis yields charge carrier densities that are on the order of the ones in FTO (aSnO2: ND = 1.5·1021 cm-3; c-SnO2: ND = 1.0·1021 cm-3 and FTO: ND = 1.4·1021 cm-3 obtained from Fig. 4 assuming εr = 9 for FTO49). These high carrier densities are typical for degenerate semiconductors such as FTO, but are unlikely for our ALD SnO2 films. According to Eq. 5, a charge carrier concentration of ND ≈ 2·1018 cm-3 is calculated for our amorphous ALD SnO2 films based on ECB-EFB ≈ 0.02 V as determined from Fig. 6a.

ECB = E FB −

k BT  N C   ln e N  D

(5)

NC is the effective density of conduction band states calculated to be 4.1⋅1018 cm-3 for an electron effective mass of me* = 0.3⋅m0 .50,51 This estimated charge carrier concentration is in very good agreement with measured charge carrier concentrations in ALD c-SnO2 but differs by 2 orders of magnitude from the ones measured by Hall effect for our a-SnO2 (Table 1) Table 1 Results of Hall measurements on 78 nm ALD-SnO2 films deposited on quartz substrate at 118°C (a-SnO2) and calcined at 450°C (crystalline, c-SnO2).

-3

ND / cm

Mobility / cm2(Vs)-1 -1

Conductivity / Scm

a-SnO2 (78 nm)

c-SnO2 (78 nm, calcined)

(1.3 ± 0.9)⋅10

(3.3 ± 1.3)⋅1018

16

5.4 ± 4.2 (7.2 ± 1.6)⋅10

3.7 ± 1.5 -3

1.7 ± 0.5

We assume that our Hall measurements on amorphous SnO2 were not reliable because of the poorly conductive a-SnO2. Hence, we consider the electrochemical determination of the conduction band position and the calculated charge carrier concentrations of ND ≈ 2·1018 cm-3 to be accurate. The conductivity of c-SnO2, however, is sufficiently high for a reliable Hall

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measurement and our results are in good agreement with the reported values for crystalline SnO2.52-54 The discrepancy between ND values measured by EIS and by Hall effect and electrochemical analysis is most likely due to the low thickness of our ALD SnO2 films: considering ND values on the order of ND ≈ 2·1018 cm-3 for a-SnO2 and ≈ 3·1018 cm-3 for c-SnO2, then the calculated Debye length: 1/ 2

1/ 2

 2ε ε   kT  LD =  0 r   E − EFB −  e   eN D  

(6)

is exceeding the film thickness (a-SnO2: LD = 16 nm at 0.6 VRHE for E-EFB = 0.45 V, which is the fitted region in the Mott-Schottky plot; c-SnO2: LD = 13 nm for E-EFB = 0.53 V). Moreover, amorphous layers very probably shrink upon calcination ending up with thicknesses smaller than the Debye length. Hence, the depletion layer extends over the whole SnO2 film, and penetrates into the supporting FTO. This effect has been also observed for thin TiO2 films on FTO 18,55. The ND values obtained from the Mott-Schottky analysis are therefore not representative for the thin ALD-SnO2 films. We also deposited thicker films of 100-200 nm, however, the fitting of impedance data for these thicker amorphous layers was not very reliable most probably due to the low conductivity of the films (see Table 1). Fig. 7 summarizes our findings on the flatband ad conduction band positions in our thin ALD SnO2 films. The redox potential of [Fe(CN)6]3-/4- (0.59 VRHE at pH 2.5) is more positive (energetically lower) than the conduction band edge of a-SnO2 but energetically slightly higher than ECB of c-SnO2. Hence, the a-SnO2/[Fe(CN)6]3-/4- contact is rectifying. Thus it is expected to suppress the oxidation of ferrocyanide and solely allowing the reduction of ferricyanide at potentials negative of EFB. As shown in Fig. 1, this rectifying behavior is perfectly achieved by our thin ALD a-SnO2 films. However, the c-SnO2/[Fe(CN)6]3-/4- contact is not

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rectifying which explains the observed metal-like behavior with the [Fe(CN)6]3-/4- redox couple in Fig. 2. Measurements with the MV2+/+ redox couple pinpoint the conduction band position to ≈0.17 V vs. RHE and Hall measurements on c-SnO2 can be used for the determination of the conduction band potential that is ≈0.52 V more positive than that of a-SnO2.

Figure 7: Energetic alignment of amorphous (a-SnO2) and crystalline (c-SnO2) with redox couples MV2+/+ (at pH 10.6), [Fe(CN)6]3-/4- (at pH 2.5) and [Ru(bpy)3]3+/2+ (at pH 2.5).

3. CONCLUSIONS Using voltammetric and XPS studies, we show that the ALD-SnO2 films are entirely pinhole-free when deposited on FTO or other substrates such as gold. The blocking was excellent for the as-grown SnO2 films down to thicknesses of 2 nm, which compared favorably to ALD-made titania buffer layers. Even after calcination at 450 °C the perfect blocking function

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is entirely preserved. It is remarkable that amorphous SnO2 ALD layers do not crack nor form pinholes upon annealing although they undergo crystallization as shown by TEM. The crystallization of ALD SnO2 films leads, however, to a 0.5 V shift in the flatband potential as was determined by impedance spectroscopy and confirmed in CV measurements with several model redox couples. Furthermore, choosing MV2+/+ redox couple in suitable pH conditions, we accurately probed the conduction band position of a-SnO2. This analysis also revealed limitations of the Hall measurements on amorphous materials providing charge carrier densities of ND ≈ 1.3·1016 cm-3 versus the calculated ones from the conduction and flatband positions (ND ≈ 1.9·1018 cm-3). The latter charge carrier concentration is close to the ones measured for the crystalline ALD SnO2 by Hall effect (ND ≈ 3·1018 cm-3). Our electrochemical analysis suggests a difference of 0.5 V in flatband potentials and a maximum difference of ≈0.52 V in the conduction band positions of amorphous SnO2 compared to crystalline SnO2, that has to be taken into account in photovoltaic or photoelectrochemical devices (e.g. for water splitting) in terms of ideal energy band alignment.

4. EXPERIMENTAL SECTION FTO glass (TEC 15, 15 Ohm/sq, G2E) was first wiped with acetone and then cleaned for 10 min in piranha solution (H2SO4/H2O2 = 3:1) prior to loading into the ALD chamber. Deposition of SnO2 was carried out at 118 °C in a home-built ALD setup with a shower-head design47 using tetrakis(dimethylamino)tin(IV) (TDMASn, 99.99%-Sn, Strem Chemicals, Inc.) and ozone. A constant growth rate of 0.065 nm/cycle was confirmed by ellipsometry (Sopra GES 5E). TDMASn was held at 65 °C. Ozone was produced by an ozone generator (AC2025, IN USA Inc.) fed with oxygen gas (99.9995% pure, Carbagas) producing a concentration of 13% ozone in O2. Nitrogen was used as a carrier gas (99.9999%, Carbagas) with a

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flow rate of 10 sccm. In some cases, the SnO2 films were deposited also on a 150 nm thick Au layer that was sputtered (Alliance Concept DP650 sputtering system) on FTO using a 10 nm sputtered Cr interfacial layer for chemical stability. The thicker blocking layers were also measured by Alfa step profilometer (Tencor Instruments). In all cases, the electrodes with blocking layers were used either as made or after subsequent calcination in air at 450 oC for 15 minutes with a heating ramp of 5 °C/min. Atomic force microscopy (AFM) images were obtained using Multimode Nanoscope IIIa (Bruker, USA) instrument. SEM images were acquired on a Zeiss Merlin microscope. Transmission electron microscopy (TEM) measurements were carried out with a Tecnai Osiris (FEI, USA). Cross-section lamellae for TEM were prepared with a Zeiss NVision 40 CrossBeam with focused ion beam (FIB). X-ray photoelectron spectroscopy (XPS) was studied on Au-supported samples using Omicron Nanotechnology instrument equipped with a monochromatized AlKα source (1486.7 eV) and a hemispherical analyzer operating in constant analyzer energy mode with a multichannel detector. The CasaXPS program was used for spectra analysis. Hall measurements were carried out with an Ecopia HMS 3000 Hall setup (Microworld) with a magnetic field of 0.54 T in a 4point configuration measuring an area of 1 cm2. The SnO2 layer was deposited on quartz substrate and contacted via soldered contacts of InSn (Microworld). Five samples per condition (as-deposited and calcined) were measured and the average values are reported. Electrochemical experiments were carried out in a one-compartment cell using Autolab Pgstat-30 equipped with the FRA module (Metrohm) controlled by the GPES-4 software. Ag/AgCl (sat’d KCl) was used as the reference electrode and a Pt wire as the counter electrode. The electrolyte solutions were purged with Ar and the measurement was carried out under Ar atmosphere in a closed electrochemical cell. Impedance spectra were measured at varying potentials, which were scanned typically from 1.3 V to -0.3 V vs. Ag/AgCl and back.

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Impedance spectra were evaluated using Zview (Scribner) software fitting with a Randles circuit where the impedance of a CPE equals: ZCPE = B(iω)-β

(7)

with ω being the EIS frequency and B, β the frequency-independent parameters of the CPE (0.8≤ β ≤1; experimental values were from 0.92 to 0.98). The capacitance, C is calculated from ZCPE by:

C=

( RCT ⋅ B)1/ β RCT

(8)

Here, RCT is the charge-transfer resistance, which is parallel to CPE. The circuit is completed with a series resistance, RS, characterizing the ohmic resistance of electrodes, electrical contacts and electrolyte solution. Alternatively, we also fitted with a Warburg resistance (when present; for a survey of circuits see Ref.56). Electrolytes, solvents and redox-active molecules were of the standard quality (p.a. or electrochemical grade) purchased from Aldrich or Merck and used as received.

Supporting Information X-ray photoelectron spectra, AFM images and additional electrochemical data. This information is available free of charge via the Internet at http://pubs.acs.org/. Acknowledgements MG thanks for the financial support from the Swiss National Science Foundation, CCEM-CH in the 9th call proposal 906: CONNECT PV, the SNSF NRP70 "Energy Turnaround" and the King Abdulaziz City for Science and Technology (KACST). LS and MG thank the Swiss National Energy Office for financial support under the PEChouse 3 project. LS acknowledges

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the support from the European FP7 FET project PHOCS (no. 309223). LK acknowledges the support from the Grant Agency of the Czech Republic (contract No. 13-07724S). Thanks are due to Dr. Pavel Janda and Dr. Hana Tarabkova for AFM measurements.

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