LETTER pubs.acs.org/JPCL
Band Alignment and Internal Field Mapping in Solar Cells Yafit Itzhaik,† Gary Hodes,† and Hagai Cohen*,‡ †
Department of Materials and Interfaces and ‡Department of Chemical Research Support, The Weizmann Institute of Science, Rehovot 76100, Israel
bS Supporting Information ABSTRACT: The internal fields and band offsets developing at individual interfaces, a critical aspect of device performance, are generally inaccessible by standard electrical tools. To address this problem, we propose chemically resolved electrical measurements (CREM) capable of resolving the internal details layer-by-layer. Applied to nanoporous photovoltaic cells, we thus extract a realistic band diagram for the multi-interfacial structure and, in particular, resolve the two p-n-like junction fields built spontaneously in the device. The lack of homogeneity common to many of these nanoporous cells is exploited here to “see” deep into the cell structure, beyond the typical depth limitations of the surface-sensitive technique. Further information on the cell operation under “real” working conditions is achieved by studying the charge trapping at each specific layer under optical and electrical stimuli. Our methodology overcomes a missing link in device characterization and in fundamental studies of nanoscale solid-state devices. SECTION: Electron Transport, Optical and Electronic Devices, Hard Matter
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olar cells are generally designed to promote charge separation via the spontaneous development of internal fields and band offsets. An approximated band diagram of the cell may be constructed from the surface and bulk properties of the individual materials used, in particular, their work function (WF), top of the valence band (VB), and bottom of the conduction band (CB) energies.1 However, it is well known that very often the results of such an analysis do not reflect that actual junction.2,3 Electrostatic changes taking place at subsurface domains when “pieces are put together” usually remain unknown, because all electrical tools available are limited to probing the integral volume between electrodes, with no direct access to internal details. A critical gap in characterization capabilities is therefore encountered. A recent technique based on X-ray photoelectron spectroscopy (XPS), chemically resolved electrical measurements (CREM),4 proposes an answer to the above limitations (see the Supporting Information for background material). Exploiting the inherent sensitivity of electron spectroscopy to the electrostatic potentials developing at the sample, CREM offers unique advantages for the study of charge transport and trapping mechanisms across heterostructures.59 The present work proposes a CREM-based systematic derivation approach to heterostructure energetics, corrected for both changes induced during device construction and those induced by the probe itself during characterization. Improved consistency of results is thus achieved, allowing detailed characterization of the device internal fields and their influence on band offsets. The first step in our procedure to map the band structure of a device is to in situ measure the WF of the fresh surface10 under “irradiation-free” conditions (a CREM-based method with no X-ray and essentially no eFG flux), providing a reference surface potential of the bare sample.11 Core level line positions are then r 2011 American Chemical Society
measured under low X-ray flux (source power of 15 W), including repeated window(s) and a second WF measurement for quick inspection of short time-scale variations. This set of measurements provides a detailed reference, to be compared with subsequent measurements under varying charging conditions (see “electrical loops” in the Supporting Information), which, eventually, allows elimination of the beam-induced electrostatic changes at each specific layer. The latter is important not only for the understanding of a particular sample but also for retaining consistency between different samples. Effects of light illumination1214 and of selective charge injection (electrons from an electron flood gun, eFG, and holes under increased X-ray flux)15,16 are then studied by following the core-level line shifts under controlled charging conditions. Measurements were performed on a slightly modified Kratos Ultra-DLD XPS spectrometer, using a monochromatic Al Kα source at 75 W. XPS chemical analysis was performed for every sample, complementing the CREM electrical study. The TiO2/InOH-S/Sb2S3/KSCN/CuSCN solar cells were prepared in a procedure described elsewhere.17 Samples representing every step in cell construction were studied, including one with a very thin (ca. 30% pore filling, probably not homogeneously distributed)18 CuSCN layer. The basic cell contains (see Figure 1a) a spin-coated nanoporous TiO2 electron conductor (grown on a spin-coated dense TiO2 on conductive FTO, fluorine-doped tin oxide, substrate), a chemically bath deposited (CBD) Sb2S3 absorber, and a solution-infiltrated CuSCN hole conductor. The CBD In-OH-S layer is used as a nucleation enhancer for the Sb2S3 (it allows a more even coating of the TiO2 with Sb2S3 but does not markedly Received: September 18, 2011 Accepted: October 26, 2011 Published: October 26, 2011 2872
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Figure 1. (a) Plan view SEM image of the cell showing the granular structure and the presence of holes through which inner domains can be probed (two examples are circled). (b) Schematic illustration of the granular multilayer structure, assuming incomplete coverage at each of the various deposition steps, as we believe to be likely.
affect the cell efficiency). The KSCN solution treatment of the absorber was found to improve the cell performance and conductivity, for which various reasons have been suggested, including improvement of contact between absorber and CuSCN or doping of the CuSCN by SCN.19 A general drawback of XPS regards the limited depth sensitivity of the technique, typically down to ∼15 nm. In the cell structure described here, however, residual signals from underlying layers could still be identified when, due to nonperfect coverage, scattered small areas of the underlying porous structure could be seen even through the full CuSCN deposit (Figure 1a). This can be considered to be an advantage, typical of porous systems, where coverage tends to be uneven: It allows analysis of “internal” layers that would normally be inaccessible by CREM. Such lateral nonuniformity may introduce electrostatic variations, typically on a length-scale given by the Debye length of the compound and, therefore, there is no guarantee that signals from uncovered (or poorly covered) domains do indeed reflect those domains that are well-covered. However, as discussed below, the exposed areas are smaller than the Debye length of the titania and the (relatively intrinsic) Sb2S3. Therefore, such exposed regions appear to represent indeed the potential of the bulk of the porous structure. Figure 2a presents raw VB spectra recorded from representative stages of the TiO2/In-OH-S/Sb2S3/KSCN/CuSCN multilayer construction. A gradual buildup of the VB top toward the Fermi level is clearly observed in the Figure. These spectra, as well as the in situ WF measurements, provide for each sample its top-surface information with no (or very limited, as with VB) access to the local data at inner domains. Correction terms are therefore needed for translating these (intermediate) surface data into those of corresponding interfaces in the full cell. These correction terms are extracted here from corresponding core electron lines. To do that, we rely on the fact that the measured
Figure 2. (a) Raw valence band spectra of the various steps in cell construction: the bare TiO2 surface (black) and after In-OH-S (blue), Sb2S3 (red), and CuSCN (pink) deposition, demonstrating the gradual buildup of density of states toward the Fermi level at EB = 0 eV. The inset depicts the Ti 2p3/2 line for the bare TiO2 surface (I), after Sb2S3 deposition (II) and after CuSCN deposition (III), showing (in addition to trivial suppression in intensity) the peak shifts due to evolution of internal fields. (b) Full cell band diagram based on the WF and valence band measurements of intermediate stages in sample preparation, subsequently corrected for the core level line shifts (with correction terms up to 600 meV). The data of the bottom of the conduction band are based on ex situ optical absorption measurements.
kinetic energy (Ekij) of a core level j at an emission site i is sensitive to the local potential (Vi) Ekij ¼ hυ EBj þ eVi þ :::
ð1Þ
where hυ is the photon energy, eVi is the electrostatic energy term, and EBj the binding energy of the inspected electronic state (additional terms irrelevant to the present discussion, e.g., the instrumental WF, are omitted intentionally from eq 1). The electrostatic term, eVi, usually missing from XPS analyses, is the main input exploited by the CREM method. For practical purposes, when a common shift is measured for the various elements in a given compound, for example, Ti and O in the TiO2, or Cu, S, C, and N in the CuSCN, chemical modifications (changes in EB) may be excluded and the shift (ΔEk) can be interpreted as a purely electrostatic potential change to be applied to all energy levels including the valence bands and the local WF. As an example, the Ti 2p3/2 lines shown in the inset to Figure 2a, 2873
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Figure 3. (a) ΔEk values measured for various core levels at separate steps in the device construction: bare TiO2, In-OH-S, Sb2S3 before annealing, after annealing, KSCN, initial stage of CuSCN deposition, full deposition of CuSCN, the latter sample being used as a reference for all data sets. (b) Schematic diagrams explaining the creation of p-n-like junction fields, with the actual potential values, measured in the present system, indicated: Top scheme presents the isolated cell components (before contact), and the bottom scheme presents in two stages the changes (ΔEk) after bringing a new component into contact, first (I) the absorber and second (II) the top electrode. Note that the latter scheme (II) represents schematically the final potential profile across the double junction of the cell. The curves (broken lines) connecting measured data points (red dots) are guesses and may deviate from the real potential profile. This is particularly so for the field in the TiO2, where we have little idea at present of the real profile shape.
all associated with the same chemical state (Ti in TiO2), present energy shifts due to changes in the local electrostatic potential, as analyzed in detail below. Figure 2b shows the resultant energy band diagram of the full cell: TiO2/In-OH-S/Sb2S3/KSCN/CuSCN. The (in situ) WF data are presented by the vacuum level (Evac) and the conduction band positioning is based on ex situ optical band gap measurements. The Figure shows convincingly the level of detail available by the technique. Remarkably, consistent results are obtained under small, controlled changes in cell preparation, which indicates the integrity of the method and repeatability achieved for identical steps in different samples.
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Whereas the above band diagram proposes rich and valuable information, it is also worth inspecting the correction values themselves. Figure 3a summarizes the kinetic energy variations (ΔEk), as determined for each sample (values already free of the instrumentally-induced “charging-like” shifts) with reference to the full cell (that of thick CuSCN); the only exception is the C 1s reference at EB = 285.0 eV, chosen arbitrarily due to marked chemical changes involved in the last step of cell construction (see the broader presentation of the data in the Supporting Information file, Figure S4). Interestingly, some of the deposition steps involve slight Ek changes only, whereas pronounced ΔEk values are observed after two specific stages: the deposition of Sb2S3 and of CuSCN. We attribute the latter ones to the development of internal junction fields, first at the junction below and, second, at the junction on top of the Sb2S3 absorber. As illustrated in Figure 3b, an electric field is expected to evolve across each of these junctions, including a change in the bottom junction taking place during the formation of the top one. This “text book”-like behavior for p-n junctions, which is now measured directly, provides important information on losses in the device and, in particular, the large fraction of open circuit voltage (VOC) that falls on the TiO2 itself.20 Note that curve II in Figure 3b, bottom, depicts a potential profile across the cell that is more realistic than the schematic band diagram in Figure 2b. Measured data points are indicated in Figure 3b as red dots, whereas the full curve is just an educated guess. Fingerprints of the nonuniform system morphology are revealed from the data in Figure 3a. First, an indium chemical shift, ΔEB = 100110 meV, appears during the KSCN deposition (note, e.g., its shift by 150 meV, whereas Ti and Sb undergo much smaller EB changes, 30 and 60 meV, respectively). As indicated by the XPS-derived Sb/Ti intensity ratio (see Supporting Information) and the KSCN stoichiometry, this chemical change arises from K-adsorption at open In-OH-S domains, between islands of Sb2S3 (the band diagram in Figure 2b is already corrected for this chemical shift). Simultaneously, preferential deposition of KSCN occurs on the Sb2S3 surface areas (see the Supporting Information). Second, Figure 3a depicts two stages in the CuSCN deposition: an early stage with a relatively small density of CuSCN grains, that is, a stage where inner domains can be probed more easily, and a later (final) stage with practically all titania pores filled + a layer of CuSCN on top of the porous structure. Clear potential changes (see Ti, In, Sb, and even K in Figure 3a) are induced at the bottom layers already during the first stage, with negligible changes during additional CuSCN deposition. On the other hand, the surface potential, as probed by the WF (see Supporting Information, Figure S4), changes first by only 100 meV, whereas it builds up an additional 400 meV upon filling the pores. These seemingly contradicting results21 arise from the surface heterogeneity realized during early CuSCN deposition stages, with domains of high and low WF, on and between the CuSCN grains, respectively. Accordingly, during the last stage in Figure 3a (see also Supporting Information, Figure S4), there are yet significant Ek changes at the top species: Cu, N, S, and, especially, K (the C line could not be used due to overlapping signals). Part of these changes reflect chemical shifts, but e100 meV arise from an electrostatic component across an intermixed region of CuSCN containing K+ and additional SCN. Notably, the electrostatic changes at inner layers (Ti, In, Sb) under nonuniform surface coverage, for example, the initial CuSCN deposition stage, appear to be “surprisingly” uniform, with little or no observable core-level line broadening (i.e., insignificant lateral 2874
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Figure 4. (a) Electrical loop plotting the incremental (Sb minus In) potential build-up at the Sb2S3 layer, before and after annealing (note the removal of hysteresis by the annealing process). (b) Response to light as recorded from the various core levels in a cell with (dark colors) and without (light colors) the KSCN treatment. Note that switching the light on (brown) and off (blue) should ideally yield identical shifts of opposite signs, unless local charge trapping occurs.
variations in electrostatic potential). This fact suggests that the Debye-length characteristic to those domains is larger than the characteristic distance between CuSCN islands. It further provides solid support for our argument above on the reliability of data probed through much smaller holes in the thick CuSCN electrode. Top layer species (Cu, K, N, and S) show more pronounced line broadenings during early stages of their deposition. Careful analysis of these line shapes suggests electrostatic variations of 50300 meV at the initially deposited CuSCN particles, attributed to the crystallite size distribution characterizing early stages of junction buildup. Eventually, these variations become rather irrelevant for the band diagram of the full cell, Figure 2b (where the pores are essentially filled), and the Cuint potential is chosen to match the value estimated for large grains, which also matches the value inferred from the K data. Beyond the ground-state band diagram information, charge trapping under external stimuli can have a major effect on device performance. The present methodology enables detailed, domain specific, inspection of such effects under light illumination and/or selective charge injection. We show two examples of how omission of certain optimization steps in the cell preparation leads to undesirable charge trapping characteristics (which are paralleled by a reduction in cell performance). The first example, the effect of annealing the Sb2S3 absorber layer, is presented in Figure 4a, showing hysteresis in core level energies under electrical loops: a step-by-step variation of the charge injection conditions, from positive to negative charge and backward (see the Supporting Information). The Figure presents differential data: incremental potential changes specific to
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the domain under inspection, with the underlying contributions (on which this layer floats) subtracted. Despite the noisy data, hysteresis is clearly observed in the as-deposited, amorphous Sb2S3 sample (Figure 4a, top curve), pointing to enhanced electron capturing at the Sb2S3 layer.5 After annealing to crystalline stibnite, the hysteresis practically disappears (Figure 4a, bottom curve), reflecting a major reduction in trapping as a result of the annealing process (in agreement with our direct cellperformance tests). Selective trapping of holes at K-containing domains could similarly be inferred from electrical loops performed at a later stage of cell construction (not shown). Figure 4b presents core level shifts of two cells, with and without the KSCN solution treatment, taking place upon switching a halogen lamp on and off.12,16 In the cell that has undergone the KSCN treatment, a smooth and reversible photoresponse analogous to band flattening in semiconductors is observed, reflecting the efficient evacuation of photoexcited electrons to ground and of holes to the top CuSCN electrode, with no pronounced charge trapping at inner interfaces.16 In contrast, the cell prepared without the KSCN treatment presents irreversibility (compare “light on” vs “light off” data) due to electron trapping, mainly at the interface with the p-doped top electrode. Whereas the overall improvement achieved by the annealing and the KSCN treatment can be easily deduced from our direct cell performance tests, identifying trap location and type is uniquely provided by the present CREM method. Consequently, distortions in the cell energetics (band offsets too) due to charge trapping under nonequilibrium conditions can be deduced from Figure 4a,b and be added as a complementing input to the band diagram data. In summary, our method (1) provides improved (corrected) band diagrams of nanoporous cells; (2) demonstrates a unique achievement, the detailed extraction of internal fields within multilayer heterostructures; and (3) provides information on charge trapping at specific locations. We can measure both the inherent fields associated with the device structure as well as those changes taking place under operation, namely, under charge or photon injection. This combination of data is of great importance for the understanding of transport issues in light sensitive systems and multilayered heterostructures in general.
’ ASSOCIATED CONTENT
bS
Supporting Information. The CREM method: technical and general comments, description of the electrical loop experiments, comments on the nonplanar morphology and its XPS characterization, and an extended presentation of the line-shift data. This material is available free of charge via the Internet at http://pubs.acs.org.
’ ACKNOWLEDGMENT We thank Tatyana Bendikov for complementary XPS work. This research was supported by a research grant from Rowland and Sylvia Schaefer and by the Israel Ministry of Science. ’ REFERENCES (1) Chasse, T.; Wu, C.-I.; Hill, I. G.; Kahn, A. Band Alignment at Organic-Inorganic Semiconductor Interfaces: α-NPD and CuPc on InP(110). J. Appl. Phys. 1999, 85, 6589–6592. (2) Freeouf, J. L.; Woodall, J. M. Schottky Barriers: An Effective Work Function Model. Appl. Phys. Lett. 1981, 39, 727–729. 2875
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(3) Niles, D. W.; Margaritondo, G. Heterojunctions: Definite Breakdown of the Electron Affinity Rule. Phys. Rev. B 1986, 34, 2923–2925. (4) Cohen, H. Chemically Resolved Electrical Measurements using X-ray Photoelectron Spectroscopy. Appl. Phys. Lett. 2004, 85, 1271–1273. (5) Rozenblat, A.; Rosenwaks, Y.; Cohen, H. Hot-Electron Characteristics in Chemically Resolved Electrical Measurements of Thin Silica and SiON Layers. Appl. Phys. Lett. 2009, 94, 213501-1–213501-3. (6) Doron Mor, I.; Hatzor, A.; Vaskevich, A.; van-der Boom-Moav, T.; Shanzer, A.; Rubinstein, I.; Cohen, H. Controlled Surface Charging as a Depth Profiling Probe for Mesoscopic Layers. Nature 2000, 406, 382–385. (7) Filip-Granit, N.; van der Boom, M. E.; Yerushalmi, R.; Scherz, A.; Cohen, H. Sub-Molecular Potential Profiling Across Organic Monolayers. Nano Lett. 2006, 6, 2848–2851. (8) Demirok, U. K.; Ertas, G.; Suzer, S. Time-Resolved XPS Analysis of the SiO2/Si System in the Millisecond Range. J. Phys. Chem. B 2004, 108, 5179–5181. (9) Suzer, S.; Sezen, H.; Dana, A. Two-Dimensional X-ray Photoelectron Spectroscopy for Composite Surface Analysis. Anal. Chem. 2008, 80, 3931–3936. (10) Cohen, H.; Nogues, C.; Zon, I.; Lubomirsky, I. e-beam Referenced Work Function Evaluation in an X-ray Photoelectron Spectrometer. J. Appl. Phys. 2005, 97, 113701/1–5. (11) This is a version of the technique described in ref 10, where the only recorded signal is the sample current. (12) Cohen, H.; Sarkar, S. K.; Hodes, G. Chemically Resolved Photovoltage Measurements in CdSe Nanoparticle Films. J. Phys. Chem. B 2006, 110, 25508–25513. (13) Sarkar, K.; Kronik, L.; Hodes, G.; Cohen, H. Defect Dominated Charge Transport in Si-Adsorbed CdSe Nanoparticle Layers. J. Phys. Chem. C 2008, 112, 6564–6570. (14) Sezen, H.; Suzer, S. Dynamical XPS Measurements for Probing Photoinduced Voltage Changes. Surf. Sci. 2010, 604, L59–L62. (15) Vilan, A.; Bendikov, T. A.; Cohen, H. Secondary Electron Emission Control in X-ray Photoelectron Spectroscopy. J. Electron Spectrosc. Relat. Phenom. 2008, 162, 99–105. (16) Cohen, H. Chemically Resolved Electrical Measurements in Organic Self-Assembled Molecular Layers. J. Electron Spectrosc. Relat. Phenom. 2010, 176, 24–34. (17) Itzhaik, Y.; Niitsoo, O.; Page, M.; Hodes, G. Sb2S3-Sensitized Nanoporous TiO2 Solar Cells. J. Phys. Chem. C 2009, 113, 4254–4256. (18) The full CuSCN deposition essentially fills the pores together with an additional ∼1 μm above the porous structure. (19) Larramona, G.; Chone, C.; Jacob, A.; Sakakura, D.; Delatouche, B.; Pere, D.; Cieren, X.; Nagino, M.; Bayon, R. Nanostructured Photovoltaic Cell of the Type Titanium Dioxide, Cadmium Sulfide Thin Coating, and Copper Thiocyanate Showing High Quantum Efficiency. Chem. Mater. 2006, 18, 1688–1696. (20) The potential value at the back end of Figure 3b, corresponding to the TiO2/FTO interface, is determined at high accuracy from the bare titania surface, using CREM electrical loops (see the Supporting Information). Possible capacitive contributions at the back contact and the interface itself were thus found to be small and of a sign opposite to that of the device junction fields discussed above. On the other hand, details of the potential drop across the granular TiO2 layer (e.g., its characteristic length compared with the single grain-size) are as yet unknown. (21) Note the close agreement between the WF changes and the C data in the Supporting Information file, Figure S4, expected to hold as long as the top C contamination is of similar chemical state for the various samples.
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