Letter pubs.acs.org/JPCL
Photoexcited Carrier Dynamics of Cu2S Thin Films Shannon C. Riha,§,† Richard D. Schaller,‡ David J. Gosztola,‡ Gary P. Wiederrecht,‡ and Alex B. F. Martinson*,§ §
Materials Science Division and ‡Nanoscience and Technology Division, Argonne National Laboratory, 9700 South Cass Avenue, Lemont, Illinois 60439, United States S Supporting Information *
ABSTRACT: Copper sulfide is a simple binary material with promising attributes for low-cost thin film photovoltaics. However, stable Cu2S-based device efficiencies approaching 10% free from cadmium have yet to be realized. In this Letter, transient absorption spectroscopy is used to investigate the dynamics of the photoexcited state of isolated Cu2S thin films prepared by atomic layer deposition or vapor-based cation exchange of ZnS. While a number of variables including film thickness, carrier concentration, surface oxidation, and grain boundary passivation were examined, grain structure alone was found to correlate with longer lifetimes. A map of excited state dynamics is deduced from the spectral evolution from 300 fs to 300 μs. Revealing the effects of grain morphology on the photophysical properties of Cu2S is a crucial step toward reaching high efficiencies in operationally stable Cu2S thin film photovoltaics.
SECTION: Spectroscopy, Photochemistry, and Excited States
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exceeding that of an equivalent planar device.15 While there are numerous advantages to nanostructuring, a basic understanding of device operation is clearly lacking and device efficiencies remain well below the 10% mark reported the seminal 1981 report.6 We seek, therefore, to get a clearer picture of the underlying factors that govern the performance of Cu2S in thin film solar cells. In particular, this letter elucidates the photophysical properties of Cu2S thin filmsmetrics that are essential to utilizing Cu2S in solar energy conversion devices. Using ultrafast spectroscopy, we map the relaxation pathways of photogenerated charge carriers. We deduce that the main source of annihilation occurs at grain boundaries. Similar to that observed in CIGS and CdTe devices,22,23 the results suggest that large grains may also be the key to obtaining good device efficiencies in Cu2S-based thin film solar cells. Aside from a handful of reports on the photophysics of Cu2S nanocrystals,24−30 many of which suggest significant substoichiometry as the main determining factor, this is the first time that the photoinduced dynamics of Cu2S thin films have been reported. Cu2S thin films were prepared either directly through atomic layer deposition (ALD)31noted herein as Cu2S-director by vapor-based cation exchange of ALD-grown ZnS thin films noted herein as CE-Cu2S. Details of the film preparation can be found in the Experimental Methods. Briefly, Cu2S-direct films were deposited at 165 °C by cycling between doses of
toichiometric copper(I) sulfide (Cu2S) is a promising material for thin film solar cells owing to its 1.2 eV direct band gap, >104 cm−1 absorption coefficient, nontoxicity, and elemental abundance.1−4 Indeed, 10% power efficient Cu2Sbased thin film solar cells were fabricated in the early 1980s through a simple topotaxial exchange with CdS thin films.5−7 While this was quite remarkable, given Si solar cell efficiencies were 11% at the time, the success was short-lived due to the operational instability of the device that degraded in performance over a period of a few weeks.8−11 The result was abandonment of this system in favor of the more stable Cu(In,Ga)S2 system. Recently there has been a resurgence of interest for stabilizing a Cu2S system12−17 owing to its elegant simplicity, extraordinary generation potential,18 and unintentional presence in other thin film PV. Attempts have targeted Cu2S directlystarting with pure and stoichiometric Cu2S,16,17 dopant inclusion,13 and passivation16,17as well as looking into alternative device structures, including nanoscale devices14,15 and replacing CdS with a stable oxide such as TiO2 or ZnO.19−21 In 2008 Alivisatos and co-workers reported a CdS nanorod/Cu2S nanocrystal solar cell with a 1.6% power conversion efficiency.14 More recently, Tang et al. described a solution-processed route to solar cells based on core−shell CdS/Cu2S nanowires.15 In that Letter, the authors argue that the solution-based cation exchange of single crystal CdS nanowires leads to relatively few interface defects, resulting in excellent charge separation and minimal minority charge carrier recombination. Based on the 1.29 μm2 active area of these nanowire devices, a power conversion efficiency of ∼5.4% was calculated as well as both an open circuit voltage and fill factor © XXXX American Chemical Society
Received: October 15, 2014 Accepted: November 3, 2014
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reaction in the presence of ZnS films.33 Thimsen et al. noted that the copper concentration approaches that for stoichiometric Cu2S after a number of CuAMD exposures for a given ZnS film thickness. Through intentional use of this approach, we prepared CE-Cu2S thin films for comparison to the directly deposited thin films. Briefly, CE-Cu2S thin films were prepared by exposing ALD-ZnS thin films to numerous 2.0 s CuAMD vapor pulses at 165 °C. X-ray fluorescence (XRF) and energy dispersive X-ray spectroscopy (EDS) measurements could not detect Zn following the extensive cation exchange. Furthermore, Hall effect measurements were conducted inside an inert atmosphere glovebox. The p-type carrier concentration of the thin films was on the order of 1015−1016 cm−3, and the mobility was near 5 cm2 V−1 s−1, resulting in conductivity of order 10−3 Ω1 cm−1. Such carrier concentrations would not be possible if the Cu2S was significantly substoichiometricas copper vacancies result in p-type dopingor n-type Zn dopants prevalent. Therefore, these electronic properties suggest the ratio of copper to sulfur in the vapor-based cation exchange samples is 2:1. Similar to the directly deposited films, upon air exposure, the carrier concentration quickly rose and was above 1019 cm−3 within hours. This phenomenon is not uncommon and is attributed to an increase in the number of Cu-vacancies due to surface oxidation, resulting in a copper-poor Cu2−xS film.16,17 Figure 1b shows that the as-made CE-Cu2S films consist of much smaller, 50 nm sized grains. Annealing the CECu2S thin films in the ALD chamber for 1 h at 250 °C with H2S doses every 2 min significantly increased the grain size (average ∼175 nm), which were also found to be continuous through the film (i.e.; smaller grains were not observed at the substrate interface), as shown in Figure 1c,e, respectively. CE-Cu2S thin films were found to be low-chalcocite by X-ray diffraction (XRD) and show a mixture of high and low-chalcocitetemplated reflections upon annealing (Figure S1). Figure 2a shows the absorption profile for 100 nm thick Cu2S thin films prepared by the aforementioned methods along with postdeposition treatments. There are two main features in the visible portion of the absorption profile for all films. The absorption onset near 1000 nm corresponds to the direct band gap4determined here from the Tauc plot in Figure 2bof 1.2 eV. A second strong increase in the absorption coefficient
bis(N,N′-disecbutylacetamidinato)dicopper(I) (CuAMD) and H2S precursors. Figure 1a,d show the topology and side profile
Figure 1. SEM images of (a) Cu2S-direct, (b) CE-Cu2S as-made, (c) CE-Cu 2S annealed, (d) side−Cu2 S-direct, (e) side−CE-Cu2 S annealed. Scale bars are 500 nm.
of a Cu2S thin film prepared directly by ALD. A wide distribution of grain diameters are present, all of which appear to extend through the entire thickness of the film. The average grain diameter was ∼150 nm, regardless of thickness in the range explored here (35 to 300 nm). X-ray diffraction of the Cu2S-direct thin film (Supporting Information Figure S1, blue trace) reveals the films to be highly oriented low chalcocite, which has been templated by the growth of high chalcocite (00l) planes parallel to the substrate, as previous discussed.31,32 The weak reflection at 40.85° 2θ corresponds to the (225) plane of low chalcocite. We previously reported that Cu2 S thin films were inadvertently produced through a vapor-based cation exchange
Figure 2. (a) Absorption profile and (b) Tauc plot of various copper sulfide thin films. Two spectral features are observed in the spectra: a sharp onset around 1.2−1.3 eV corresponding to the band gap, and a second strong absorption around 1.9−2.1 eV, which is often attributed to a second direct band gap. 4056
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Figure 3. Transient absorption spectroscopy data from a cation-exchanged film. Plot (a) is a contour plot showing the 2D spectra evolution with time. Four components (b) were used to fit the transient absorption spectrum, which includes an initial, short-lived photoinduced absorption (gray, solid line), a higher energy photoinduced bleach (purple, dash-dot) that transitions into a lower energy photoinduced bleach (blue, dotted), and a long-lived photoinduced absorbance (orange, dash). A band energy diagram (c) shows the different transitions from which each spectral component may derive.
films are often broader, overlapping, and therefore challenging to interpret. Therefore, we deconvolute the spectra with Glotaran, a global analysis software program designed for timeresolved spectroscopy data.37 Component analysis revealed that four spectral components, which displace each other in sequence, were required to fit the TA data from all Cu2S thin films (Figure 3b). The first component is a very short-lived induced absorption centered around 1000 nm. As this component decays, two negative features appear, with the higher energy component decaying into the lower energy component (Figure S3). After 10s of ps this second feature is replaced by a long-lived photoinduced absorption. Figure 3c depicts a mechanism that illustrates the multiple decay pathways of the photoexcited charge carriers consistent with the observed spectral evolution. Following the pump pulse (405 nm), carriers are excited into a multitude of conduction band states. This feature decays within the first picosecond, most likely due to relaxation into the lower conduction band energy levels (intraband relaxation). According to density functional theory there are two conduction band densities of states near the band edge,4,34 which may explain the second and third component bleach features that are only marginally separated in energy and time delay. To deduce the origin of these spectral bleaches, we turn
was observed at higher energy (1.9−2.1 eV) that may be attributed to transitions to higher conduction bands (refer to Figure S2).4,34 In addition to the two spectral features in the visible region, near-infrared (NIR) absorption is often observed in nonstoichiometric Cu2S, which is attributed to free carrier absorption. Cu-vacancies are formed as the material oxidizes in air, which increases the free carrier concentration and leads to the NIR absorption.17,35,36 To illustrate this, we took UV− vis−NIR data of the various thin films either immediately after removal from an inert atmosphere or after sitting in air for a few weeks. For illustration purposes, we show the effect for an annealed CE-Cu2S thin film that has been exposed to air in Figure 2a. Clearly the NIR absorption is much higher compared to the other thin films that were protected from the environment. Transient absorption (TA) spectroscopy, combined with time-resolved photoluminescence spectroscopy, is a powerful route to uncovering the mechanisms and time constants for photoexcited charge carrier dynamics. Figure 3a shows a representative 2D TA spectral evolution, with the hot colors corresponding to a pump-induced absorption increase and the cool colors indicating an absorption decrease or “bleach”. Compared to molecular species, the spectral features of thin 4057
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Figure 4. Decay times of the four spectral components as a function of grain area for 3 film thicknesses: 35 nm (blue diamond), 80 nm (purple circle), 282 nm (orange triangle). We observe a linear relationship between (b) τ2 and (c) τ3 and grain area, while there is little correlation between grain area and (a) τ1 and (d) τ4. This corroborates the hypothesis that charge carriers from the band edge are being annihilated at trap states near grain boundaries.
to time-resolved, streak camera detection of photoluminescence (TRPL, Figure S4). Similar to the TA spectra, two features were observed in the TRPL data with small energy separation and with lifetimes (∼7 and 30 ps) strikingly similar to those determined from TA. A negative TA signal (bleach) can result from more light reaching the detector due to less sample absorption or additional light from photon emission. Therefore, negative components 2 and 3 were attributed to stimulated emission. The higher energy bleach has an average decay time of ∼6 ps, resulting from stimulated emission in competition with relaxation into the lower energy conduction band level. The lower energy bleach decays after an average of 30 ps, again due to stimulated emission in addition to relaxation of electrons into the penultimate transient state. In addition to the prominent bleaches in components 2 and 3, we observe a small positive absorption around 1150 nm (see also Figure S2) that is consistent with the energy separation (∼1.1 eV) of the first and second steady-state absorption peaks, at 910 and 510 nm, respectively. Therefore, we tentatively assign these positive absorption features, which are clearly convoluted with other features in this wavelength region, to electron excitation from the first (twinned) conduction band manifolds to the second major density of states in the conduction band. In order to assign the final induced absorption component with a lifetime of ∼500 ns, a pump power TA study was performed. A linear relationship was observed, Figure S5, between the pump power
and the intensity of the photoinduced absorption. On the other hand, no correlation could be made between the lifetime and pump power. These observations are consistent with discrete traps states below the conduction band edge.28 Furthermore, the long-lived photoinduced absorption has characteristic energy consistent with the difference between a level slightly below the first CB level and the second CB manifold. Next a systematic study was performed to uncover any correlations between the photoexcited carrier lifetimes and Cu2S thin film structural and chemical properties. First we analyzed the effect of well-known film oxidationbased on the exposure time to ambient airon the excited state dynamics. The control samples were sealed in an air-free quartz cuvette in the glovebox that was removed just prior to the TA measurements. After the initial TA spectrum was recorded with the sample sealed in the quartz cuvette, the seal was cracked, and additional TA spectra were taken in air. As shown in Figure S6, little change in the time constants were observed, indicating that oxidation (and the corresponding change in carrier concentration that span 3−4 order of magnitude) plays surprisingly little role in determining the excited state lifetimes of the photoexcited carriers. In addition, we also looked at the effect of film preparation method (ALD direct versus vapor-based cation exchange), annealing conditions, salt activation, Na ion treatments, and film structure (i.e., thickness, grain diameter, grain area). Of the 4058
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N2 purge. (Safety note: H2S is a poisonous gas but not flammable at 4% concentration.) Vapor-Based Cation Exchange. First, thin films of ZnS were deposited at 200 °C in the Savannah 200 with diethylzinc (DEZ) and 4% H2S. The sequence for deposition was 0.015 s DEZ, 8 s N2 purge, and 0.1 s H2S, 8 s N2 purge. Following deposition, the reactor was coated with 200−300 nm of Al2O3 to prevent the CuAMD precursor from being unproductively consumed by exchange reaction with the reactor walls. Next, a vapor-based cation exchange was performed in the Savannah 200 at 165 °C with 2.0 s CuAMD (heated to 150 °C) pulses. Typically, 200 pulses were needed to completely convert 100 nm thick ZnS films to Cu2S. Annealing Conditions. Cu2S thin films prepared by both methods were subject to different annealing conditions. In vacuo annealing was performed in the Savannah 200 in the range of 250−325 °C for 15−60 min, with 0.1 s 4% H2S pulses every 2 min. Other annealing experiments were carried out in the N2-filled glovebox on a hot plate with a metal box to contain the heat. The temperature range for these inert atmosphere annealing conditions were 300−500 °C for a period of 15−60 min. Similar annealing experiments were also performed in ambient air on a benchtop. Na+ and Cl− Treatments. In addition to exploring different annealing conditions, Na+ and Cl− treatments were also tested to determine the effect on grain size and charge carrier recombination. In one method, Cu2S thin films were submersed in anhydrous methanol solutions containing various metal chloride salts for 10 s, dried on a Kim wipe, and annealed in vacuo at 325 °C for 60 min. In another method, anhydrous methanol solutions containing various metal chlorides or NaOH was spin coated at 2000 rpm on the Cu2S thin films, followed by in vacuo annealing at 325 °C for 60 min. All steps for each method were carried out in the N2-filled glovebox. Characterization. Hall measurements were recorded at room temperature with an Ecopia HMS-3000 Hall measurement system. The probe current was selected such that the voltage signal was between 0.1 and 1 V. SEM images were collected using a Hitachi S-4700-II SEM with an EDS detector. Reflectance-corrected UV−vis−NIR was collected on a Varian Cary 5000 with an integrating sphere accessory (DRA-2500). XRD was taken on a Philips X′Pert Pro MRD diffractometer using Cu Kα (λ = 1.5418 Å) with operation conditions set to 30 kV/30 mA. A 60 mm graded parabolic W/Si mirror with a 0.8° acceptance angle and a 1/8° divergence slit were used to condition the incident beam. A 0.27° parallel plate collimator and a flat pyrolytic graphite monochromator were positioned in front of the PW3011/20 sealed proportional point detector to collect the reflected beam. Transient absorption spectroscopy was performed using an amplified Ti:sapphire laser system (Newport Spectra-Physics Spitfire Pro) centered at 800 nm, with a 1.67 kHz repetition rate and a 120 fs pulse. The laser is equipped with a TOPAS optical parametric amplifier (OPA). The output of the OPA was set at either 400 or 800 nm and served as the excitation pulse. The white light continuum probe was generated by focusing 10% of the amplifier output into a 2 mm thick sapphire window. The pump and probe beams were directed to an Ultrafast Systems Helios transient absorption spectrometer equipped for probe detection in the near-infrared (NIR) spectral region from approximately 830 nm to 1450 nm (using either Si or InGaAs detectors). The pump beam was focused to a spot size of approximately 200 nm diameter, with the probe beam slightly smaller than the pump beam. TA data
variables examined, only grain structure significantly and systematically altered the photoexcited carrier lifetimes. Figure 4 displays the decay times for each of the four components mentioned above plotted as a function of grain area, which was measured using scanning electron microscopy (SEM). Panels a−c show the effects of grain area for 3 different film thicknesses: 35, 80, and 282 nm. Accurately measuring τ4 required a different time scale to be analyzed (note the scale in panel d is in nsmeasured with the Ultrafast Sytems EOS detectorwhile panels a−c are in psmeasured using the Ultrafast Systems Helios detector) and therefore, only one film thickness was measured. Clearly grain area does not affect τ1 and τ4, as might be predicted according to their assignment as thermal cooling and trap-relaxation processes. On the contrary, there is a strong correlation between the grain area and τ2 and especially τ3 at every thickness. That the relaxation time increases with increasing grain area further suggests that the fourth component in the transient spectra is related to a trap state at the grain boundary. To the extent that τ3 also scales with thickness for an equivalent grain area, we further hypothesize that the film surface and/or interface with substrate provide an additional relaxation route. Overall we find that minority carrier lifetimes in these Cu2S thin films are dominated by picosecond relaxation to trap-like states on the grain surfaces. This is not entirely surprising as other thin film absorbers exhibit an equivalent effect, most notably CdTe and CIGS.22,23 However, optimization of grain growth and surface passivation treatments have largely alleviated the deficiency in these materials. While our initial studies include some treatments that have been effective in related systems (salt activation and Na ion treatment), our experiments are far from exhaustive. Therefore, additional studies on grain growth and surface treatments of Cu2S thin films are warranted. This study provides insights into why such high device efficiencies may have been achieved in CdS/Cu2S heterostructure devices fabricated through the original 1970s topotaxial exchange process. CdS and Cu2S are well latticed matched and single-crystal Cu2−xS films can be made through an exchange reaction with CdS single crystals.38 Because CdS thin films with micron-size grains were used, it can be inferred that the lateral grain size of the resulting Cu2S films in these devices may also be of order of microns.7 This may also support the impressive efficiency of the CdS/Cu2S nanowire device, given that the Cu2S films that formed were templated from micron-length single crystal CdS nanowires.15 In contrast to the topotaxial approach, methods of fabricating stand-alone thin film Cu2S solar cells result in nanometer-sized Cu2S grains.14 Therefore, we posit that highly efficient and stable Cu2S thin film photovoltaics will not only require a deviation from CdS as the n-type mate but also a novel fabrication route that significantly augments grain size and prioritizes surface passivation.
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EXPERIMENTAL METHODS Cu2S Thin Film PreparationDirect Deposition. Direct deposition of Cu2S thin films was carried out using an UltratechCambridge Nanotech Savannah 200 ALD coupled to a N2-filled glovebox. The deposition temperature and bis(N,N′-disecbutylacetamidinato)dicopper(I) precursor (CuAMD-purchased from DOW Chemical Company) temperature were 165 and 150 °C, respectively. Films were grown by alternating 2.0 s CuAMD and 0.1 s 4% H2S pulses, separated in time with a 13 s 4059
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(7) Hall, R. B.; Meakin, J. D. Design and Fabrication of HighEfficiency Thin-Film CdS−Cu2S Solar-Cells. Thin Solid Films 1979, 63, 203−211. (8) Al-Dhafiri, A. M.; Russell, G. J.; Woods, J. Degradation in CdS− Cu2S Photovoltaic Cells. Semicond. Sci. Technol. 1992, 7, 1052−1057. (9) Matsumoto, H.; Nakayama, N.; Yamaguchi, K.; Ikegami, S. Stability and Mechanism of Degradation in CdS−Cu2S Ceramic Solar Cells. Jpn. J. Appl. Phys. 1976, 15, 1849−1850. (10) Matsumoto, H.; Nakayama, N.; Yamaguchi, K.; Ikegami, S. Improvement of Stability in CdS−Cu2S Ceramic Solar Cells. Jpn. J. Appl. Phys. 1977, 16, 1283−1284. (11) Rastogi, A. C.; Salkalachen, S. X-ray Photoelectron Spectroscopy Studies of Copper Diffusion Behaviour and Related Degradation Phenomena in Thin Film Cds: Cu2S Solar Cells. Sol. Cells 1983, 9, 185−202. (12) Lotfipour, M.; Machani, T.; Rossi, D. P.; Plass, K. E. αChalcocite Nanoparticle Synthesis and Stability. Chem. Mater. 2011, 23, 3032−3038. (13) Machani, T.; Rossi, D. P.; Golden, B. J.; Jones, E. C.; Lotfipour, M.; Plass, K. E. Synthesis of Monoclinic and Tetragonal Chalcocite Nanoparticles by Iron-Induced Stabilization. Chem. Mater. 2011, 23, 5491−5495. (14) Wu, Y.; Wadia, C.; Ma, W.; Sadtler, B.; Alivisatos, A. P. Synthesis and Photovoltaic Application of Copper(I) Sulfide Nanocrystals. Nano Lett. 2008, 8, 2551−2555. (15) Tang, J. Y.; Huo, Z. Y.; Brittman, S.; Gao, H. W.; Yang, P. D. Solution-Processed Core-Shell Nanowires for Efficient Photovoltaic Cells. Nat. Nanotechnol. 2011, 6, 568−572. (16) Martinson, A. B. F.; Riha, S. C.; Thimsen, E.; Elam, J. W.; Pellin, M. J. Structural, Optical, and Electronic Stability of Copper Sulfide Thin Films Grown by Atomic Layer Deposition. Energy Environ. Sci. 2013, 6, 1868−1878. (17) Riha, S. C.; Jin, S.; Baryshev, S. V.; Thimsen, E.; Wiederrecht, G. P.; Martinson, A. B. F. Stabilizing Cu2S for Photovoltaics One Atomic Layer at a Time. ACS Appl. Mater. Interfaces 2013, 5, 10302−10309. (18) Wadia, C.; Alivisatos, A. P.; Kammen, D. M. Materials Availability Expands the Opportunity for Large-Scale Photovoltaics Deployment. Environ. Sci. Technol. 2009, 43, 2072−2077. (19) Liu, G.; Schulmeyer, T.; Thissen, A.; Klein, A.; Jaegermann, W. In Situ Preparation and Interface Characterization of TiO2/Cu2S Heterointerface. Appl. Phys. Lett. 2003, 82, 2269−2271. (20) Bessekhouad, Y.; Brahimi, R.; Hamdini, F.; Trari, M. Cu2S/TiO2 Heterojunction Applied to Visible Light Orange II Degradation. J. Photochem. Photobiol., A 2012, 248, 15−23. (21) Burgelman, M.; Pauwels, H. J. Theoretical Advantages of pn+Type Cu2S-ZnO Solar Cell. Electron. Lett. 1981, 17, 224−226. (22) Singh, L.; Saxena, M.; Bhatnagar, P. K. Short Circuit Current Variation of CIGS Solar Cells with Grain Boundaries Recombination Velocity. Indian J. Pure Appl. Phys. 2004, 42, 841−844. (23) Paudel, N. R.; Young, M.; Roland, P. J. Post-Deposition Processing Options for High-Efficiency Sputtered CdS/CdTe Solar Cells. J. Appl. Phys. 2014, 115, 064502. (24) Artemyev, M. V.; Gurin, V. S.; Yumashev, K. V.; Prokoshin, P. V.; Maljarevich, A. M. Picosecond Absorption Spectroscopy of Surface Modified Copper Sulfide Nanocrystals in Polymeric Film. J. Appl. Phys. 1996, 80, 7028. (25) Klimov, V. I.; Karavanskii, V. A. Mechanisms for Optical Nonlinearities and Ultrafast Carrier Dynamics in CuxS Nanocrystals. Phys. Rev., B: Condens. Matter 1996, 54, 8087−8094. (26) Klimov, V. I.; Haring-Bolivar, P.; Kurz, H.; Karavanskii, V. A. Optical Nonlinearities and Carrier Trapping Dynamics in Cds and CuxS Nanocrystals. Superlattices Microstruct. 1996, 20, 395−404. (27) Klimov, V.; Haring Bolivar, P.; Kurz, H.; Karavanskii, V.; Krasovskii, V.; Korkishko, Y. Linear and Nonlinear Transmission of CuxS Quantum Dots. Appl. Phys. Lett. 1995, 67, 653. (28) Lou, Y.; Chen, X.; Samia, A. C.; Burda, C. Femtosecond Spectroscopic Investigation of the Carrier Lifetimes in Digenite Quantum Dots and Discrimination of the Electron and Hole
was probed in the NIR region (830−1600 nm). For probing time scales longer than 3 ns via transient absorption, we switched to an Ultrafast Systems EOS spectrometer, which uses a fiber laser continuum for the probe. Time-resolved photoluminescence spectroscopy was collected using 35 fs excitation pulses at 400 nm focused to a 400 μm diameter spot. Emitted photons were collected with a lens and directed to a 150 mm spectrograph and streak camera detector that operated in a photon counting mode.
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ASSOCIATED CONTENT
S Supporting Information *
XRD, TRPL data, and TA pump power dependence. This material is available free of charge via the Internet at http:// pubs.acs.org.
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AUTHOR INFORMATION
Corresponding Author
*E-mail:
[email protected]. Present Address †
(S.C.R.) University of Wisconsin-Stevens Point, Department of Chemistry, D129 SCI, 4001 Fourth Avenue, Stevens Point, WI 54481. Notes
The authors declare no competing financial interest.
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ACKNOWLEDGMENTS S.C.R. was supported in part by the Department of Energy (DOE) Office of Energy Efficiency and Renewable Energy (EERE) Postdoctoral Research Awards under the EERE Solar Program administered by the Oak Ridge Institute for Science and Education (ORISE) for the DOE. ORISE is managed by Oak Ridge Associated Universities (ORAU) under DOE contract number DE-AC05-06OR23100. The research was performed at Argonne National Laboratory, a U.S. Department of Energy Office of Science Laboratory operated under Contract no. DE-AC02-06CH11357 by UChicago Argonne, LLC. Use of the Center for Nanoscale Materials was supported by the U.S. Department of Energy, Office of Science, Office of Basic Energy Sciences, under Contract No. DE-AC0206CH11357. The electron microscopy was accomplished at the Electron Microscopy Center at Argonne National Laboratory, a U.S. Department of Energy Office of Science Laboratory operated under Contract No. DE-AC0206CH11357 by UChicago Argonne, LLC.
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REFERENCES
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The Journal of Physical Chemistry Letters
Letter
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