Enhanced Multiple Exciton Generation in PbS|CdS Janus-Like

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Enhanced Multiple Exciton Generation in PbS| CdS Janus-Like Heterostructured Nanocrystals Daniel M. Kroupa, Gregory F. Pach, Márton Vörös, Federico Giberti, Boris D Chernomordik, Ryan W. Crisp, Arthur J Nozik, Justin C. Johnson, Rohan Singh, Victor I. Klimov, Giulia Galli, and Matthew C. Beard ACS Nano, Just Accepted Manuscript • Publication Date (Web): 14 Sep 2018 Downloaded from http://pubs.acs.org on September 14, 2018

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Enhanced Multiple Exciton Generation in PbS|CdS Janus-Like Heterostructured Nanocrystals Daniel M. Kroupa,1,2# Gregory F. Pach,1# Márton Vörös,4,5 Federico Giberti,5 Boris D. Chernomordik,1 Ryan W. Crisp,1,3 Arthur J. Nozik,1,2 Justin C. Johnson,1 Rohan Singh,6 Victor I. Klimov,6 Giulia Galli,4,5,7 and Matthew C. Beard1* 1. 2. 3. 4. 5. 6. 7.

Chemistry & Nanoscience Center, National Renewable Energy Laboratory, Golden, Colorado 80401, United States Department of Chemistry and Biochemistry, University of Colorado, Boulder, Colorado 80309, United States Department of Physics, Colorado School of Mines, Colorado 80401, United States Materials Science Division, Argonne National Laboratory, Lemont, Illinois 60439, United States Institute for Molecular Engineering, University of Chicago, Chicago, Illinois 60637, United States Chemistry Division, Los Alamos National Laboratory, Los Alamos, New Mexico 87545, United State Department of Chemistry, University of Chicago, Chicago, Illinois 60637, United States *Corresponding Author ([email protected]) # Equal Contributors

Abstract Generating multiple excitons by a single high-energy photon is a promising third generation solar energy conversion strategy. We demonstrate that multiple exciton generation (MEG) in PbS|CdS Janus-like hetero-nanostructures is enhanced over that of single-component and core/shell nanocrystal architectures, with an onset close to two times the PbS band gap. We attribute the enhanced MEG to the asymmetric nature of the hetero-nanostructure that results in an increase in the effective Coulomb interaction that drives MEG and a reduction of the competing hot exciton cooling rate. Slowed cooling occurs through effective trapping of hotholes by a manifold of valence band interfacial states having both PbS and CdS character, as evidenced by photoluminescence studies and by ab initio calculations. Using transient photocurrent spectroscopy, we find that the MEG characteristics of the individual nanostructures are maintained in conductive arrays and demonstrate that these quasi-spherical PbS|CdS nanocrystals can be incorporated as the main absorber layer in functional solid-state solar cell architectures. Finally, based upon our analysis, we provide design rules for the next generation of engineered nanocrystals to further improve the MEG characteristics. Keywords: multiple excition generation, carrier multiplication, quantum dot, nanocrystal, solar cell, transient absorption spectroscopy TOC Graphics

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As conventional single-junction solar cells begin to approach their fundamental theoretical efficiency limits, future generation solar energy technologies need to focus on achieving efficiencies above the Shockley-Queisser (SQ) limit. As first suggested by Nozik in 2002, multiple exciton generation (MEG), also known as carrier multiplication (CM), can be enhanced in semiconductor nanocrystals (NCs) compared to parental bulk solids.1 This process is analogous to impact ionization in bulk semiconductors and describes the conversion of a single high-energy photon into multiple electron-hole pairs (excitons). Solar cells utilizing MEG can, in principle, surpass the SQ limit for solar energy conversion.2-4 Following the first spectroscopic observation of MEG/CM in PbSe NCs in 2004,4 multiple studies have provided direct evidence for MEG enhanced photon-to-electrical-current conversion by demonstrating solar cells with peak external quantum efficiencies (EQE) greater than 100%.5-7 Recently we studied MEG in a hydrogen evolution reaction and observed an EQE for hydrogen generation exceeding 100%.8 MEG has also been experimentally verified in photodetectors9 and other photoelectrochemical systems.10 However, in order for MEG to have a significant impact on future solar energy conversion technologies, new materials exhibiting higher efficiencies are necessary. In particular, the threshold should be as close to two times the band gap (Eg) as possible. Fundamental studies of MEG have been performed for many material systems including 2D graphene,11 1D-carbon nanotubes,12 and semiconductor NCs of varying composition and dimensionality, such as 0D quantum dots (QDs),13-19 quasi-1D quantum rods (QRs),20 2D nanosheets21 and core/shell structures.22 Historically, lead chalcogenide QDs (PbE; E=S, Se, or Te) have been the most widely studied materials for MEG.4, 23-32 Two important characteristics of the process are the threshold photon energy, , at which the quantum yield (QY) for exciton generation begins to exceed unity and the electron-hole pair creation energy, , which is the excess photon energy required to generate an additional electron-hole pair after . The electron-hole pair creation energy can be related to the intraband energy-loss (“cooling”) rate 2 ACS Paragon Plus Environment

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due to non-MEG processes (rcool) and the characteristic time of the individual impact-ionization collision (τMEG) by εeh = rcoolτMEG.33-34 Since smaller values of εeh correspond to stronger MEG, the boost in the MEG performance is possible by desigining a nanostructure so as to slow down intraband cooling and/or enhance Coulomb interactions responsible for exciton-generating collisions. PbE QDs typically exhibit of ~2.6 - 3 Eg while PbE QRs have shown enhanced MEG over QDs exhibiting a similar but an almost two-fold enhancement in , attributed to increased Coulombic coupling in the quasi-1D structure.20 Hetero-nanostructures have also been used to improve MEG performance. Specifially, Cirloganu et al. 35 used wavefunction engineering in PbSe/CdSe core/shell QDs to “break” symmetry between the conduction and the valence band (and hence reduce hvth) and to slow down cooling of shell localized “hot” holes (reduces εeh). This allowed for pushing very close to the fundamental energy-conservation-defined limit of ~2.2 , and simultaneously achieve a nearly four-fold increase in the multiexciton yield compared to PbSe QDs with a similar band gap. A slowed hole relaxation was realized due to electronic decoupling between the core and the shell states occurring at large shell thicknesses and the development of a sizable energetic gap between the shell and the core states. The small core size, on the other hand, allowed for an increased Coulomb interaction among valence band carriers, which further helps drive MEG. A downside of the PbSe/CdSe core/shell structures, which limits their utility in photovoltaic technologies, is that photogenerated holes are localized in the core with the shell acting as a barrier for hole extraction, ultimately leading to low photocurrents in conventional solar cell architectures.36 Recently we developed the synthesis of Janus-like PbE|CdE NCs (E=S,Se) that have a spherical shape and a spatially asymmetric composition.37 In this study, we use femtosecond transient absorption (TA) spectroscopy to probe the MEG characteristics of PbS|CdS Janus-like heterostructured NCs. We find MEG to be highly efficient with characteristics similar to those of the PbSe/CdSe core/shell QDs reported previously. Importantly, as distinct from the core/shell structures, the Janus-like NCs provide access to both electrons and holes, which allows for incorporation in photoconductive devices. In accordance, we demonstrate that MEG is retained in electronically coupled Janus-particle arrays using transient photocurrent spectroscopy38 and show power conversion efficiencies of nearly 3% in unoptimized, proof-of-principle solar cells. Results and Discussion PbS|CdS Janus-like heterostructure NCs were synthesized following a partial cation exchange method described by Zhang et al.37 Varying the reaction time and Pb-to-Cd precursor ratios allow for control over the extent of exchange of the as-prepared CdS QDs, while varying the reaction temperature and size of the original QDs dictates the dimensions of the resulting heterostructured NCs. The extent of exchange is inferred using a combination of absorption spectroscopy, X-ray diffraction (XRD) (Fig. 1b), X-ray fluorescence (XRF) spectroscopy, and 3 ACS Paragon Plus Environment

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transmission electron microscopy (TEM) (Fig. 1d). After the cation exchange reaction, the absorbance spectrum of the original CdS QDs (blue-trace, Fig. 1a) gains new features below 2.5 eV (black-trace, Fig. 1a) indicative of the newly formed PbS domain. The cation exchange process can be quite uniform, as shown by the sharp first exciton absorption peak associated with the PbS domain (Fig. 1a, inset), as well as a sharp lattice boundary between the PbS and CdS domains (Fig. 1d). Additional absorption and emission spectra, and XRD patterns of NCs with varying Pb:Cd ratios are reported in supporting figures Figs. S1 and S2.

Figure 1. Representative optical and structural characterization of original CdS QDs and ion-exchanged Janus PbS|CdS NCs. (a) UV-Vis-NIR absorption spectra of original CdS QDs (blue) and final PbS|CdS heterostructured NCs (black). (b) An X-ray diffraction pattern of PbS|CdS NCs shows peaks associated with crystalline CdS and PbS domains. (c) TEM images of starting CdS QDs used for cation exchange and (d) resulting PbS|CdS NCs. The PbS|CdS NCs are seen to be approximately half the size (~5 nm) of the initial CdS QDs (~10 nm), suggesting that CdS QDs partially dissociate during the cation exchange process.


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Visible and infrared (IR) photoluminescence (PL) spectroscopy are used to reconstruct an approximate energy level diagram for the PbS and CdS domains in PbS|CdS Janus-like NCs (Fig. 2a,b). The original CdS QDs show a sharp PL peak attributed to band-edge emission as well as a broader band at lower energies corresponding to trap-state emission (Fig. 2a, offset trace), presumably due to under-coordinated surface sulfur (S) anions39-42 The PL of the Janus-like NCs, when photoexcited below the CdS absorption onset (hvex = 2.0 eV), shows a peak at ~1.04 eV indicative of PbS band-edge emission (Fig. 2a, green trace). The PL spectrum broadens and extends into the visible region when the excitation is above the CdS absorption onset (hv ex = 3.1 eV; Fig. 2a, blue trace). This spectrum can be decomposed into three bands: the PbS band-edge emission (Fig. 2a, dashed red trace), “indirect” emission due to a CdS conduction band (CB) to a PbS valence band (VB) transition (Fig. 2a, dashed blue trace) and an emission peak which we assign to a transition from the PbS CB edge to the VB interface state (Fig. 2a, dashed green trace); see the discussion below. The PL emission at higher energies of PbS|CdS Janus-like NCs, observed when the photoexcitation is set above the absorption onset of the CdS domain is a strong indicator of the assigned type-I band offset. In Fig. 2b, we display an approximate energy diagram of Janus PbS|CdS NCs along with optical transitions responsible for the PL features described above. The PbS domain states are denoted by the grey lines and labeled as 1Ee, 1Eh, 2Ee, 2Eh, etc. The CdS domain states are shown by the pink lines. In the PL data we find evidence of slowed carrier cooling since the PL emission is broadened to higher energies (well above the band-gap of the PbS NC domain) when photoexcitation occurs above the CdS absorption onset. This suggests that electrons residing in the CdS CB states have slowed relaxation to the bottom of the PbS CB band edge, while holes are also “stabilized” in higher-energy, above-band-edge states likely associated with the PbS/CdS interface. A similar effect has been reported in PbSe/CdSe QD films where hot electrons transfer between the two materials and slow their cooling to the PbSe band edge.43 In our picture (Fig. 2b) we denote this by a barrier between the PbS and CdS domains. Near the CB minimum, electronic states are primarily composed of cation orbitals (Pb, Cd) while near the VB maximum the states are composed of the anion orbitals (S). Therefore, mixing of electronic states across the PbS/CdS interface is expected to be considerably stronger in the VB than in the CB. This could explain the formation of mixed PbS/CdS interfacial VB states that are only weakly coupled to PbS-only VB-edge state, and thus can serve as fairly long-lived “hot” hole traps. Our assignment of a band alignment at the PbS/CdS interface to type-I (Fig. 2b) is consistent with literature reports on PbS|CdS multipod heterostructures,44 as well as prior studies on PbS/CdS core/shell particles.45,46 Experimental evidence for a type-I alignment stems from femtosecond TA spectroscopy experiments, which find that electrons selectively photoexcited into the PbS domain do not undergo charge transfer to the CdS domain (expected for a type-II energy alignment; supporting information Fig. S3). Further experimental evidence of a type-I band alignment is the sharp excitonic absorbance feature (inset, Fig. 1a) attributed to the PbS domain of the heterostructure. In contrast, a poorly defined first exciton absorption band is 5 ACS Paragon Plus Environment

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typially seen for quasi-type-II PbSe|CdSe Janus-like NCs and PbSe/CdSe core/shell QDs where conduction band-edge electrons are able to delocalize over the entire heterostructure due to close proximity between CB minima in both domains.35,36,47

Figure 2. PbS|CdS Janus-like NC energetics. (a) The original CdS QD PL spectrum (black trace, offset vertically for clarity) shows band-to-band emission at 2.59 eV and a defect band at 2.0 eV. For the Janusparticles the PL spectrum depends upon pump-photon energy. With excitation below the CdS absorption onset (hvex = 2.0 eV), the emission spectrum (green trace) shows mainly PbS band-to-band emission. With excitation above the CdS absorption edge (hvex = 3.1 eV), the emission spectrum broadens and extends into the visible (blue trace). The spectrum can be decomposed into band-to-band emission at 1.04 eV (dashed-red trace), CdS CB to PbS VB “indirect” emission (dashed blue trace) at 1.15 eV, and interface trap state emission (dashed green trace) centered at 1.46 eV. (b) An approximate energy level diagram for 50:50 PbS|CdS heterostructured NCs recreated based the PL spectra in (a). It shows a type-I band alignment between the PbS and CdS domains. Colored lines correspond to the emission peaks shown in part (a). The dashed green lines show the interface state manifold. Finally, gray lines represent the core PbS states while the pink lines are for CdS.

MEG Characteristics To study MEG in the Janus-like NC heterostructures, we use TA spectroscopy to probe the NC interband bleach that arises due to a state-filling effect, as previously described48 (representative TA data and MEG analysis are reported in supporting information Fig. S6). Figure 3 shows QYs of PbS|CdS Janus-like NCs of varying composition compared to other material systems reported previously in literature. We studied 50:50, 60:40, and 90:10 PbS|CdS NCs. For PbS|CdS NCs with less than 50% Pb content, TA MEG analysis proved difficult because of broadening/washing-out of the first exciton absorption peak likely due to high NC size polydispersity and/or quasi-type-II band alignment. For 90:10 PbS|CdS NCs, we observed MEG characteristics that were similar to those of pure PbS NCs (Fig. 3, purple and orange circles).29, 35 For the 50:50 PbS|CdS NCs, we find that is very close to the energy-conservation-defined limit of 2 . Further, we observe a rapid increase in the photon-to-exciton with increasing pump energy before plateauing near 1.2. The QY then increases again above 3Eg.


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Figure 3. Summary of quantum yields (QYs) of photon-to-exciton conversion for various NC systems. QY as a function of normalized photon energy (hv/Eg) for Janus PbS|CdS NCs of various compositions studied here (solid circles) in comparison to literature results for PbS QDs, PbSe QRs and PbSe/CdSe core/shell 20, 29, 35 NCs (open circles). The dashed blue trace represents the expected MEG QY that can be expected when 25% of photoexcited carriers undergo MEG with 0.98.

Previously, we derived an expression for QYs of photon-to-exciton conversion assuming a competition between carrier cooling and impact-ionization-like collisions leading to MEG. In that analysis, the number of e-h pairs generated per absorbed photon can be expressed as follows:49-50

1 %

!"

#

# ! "

!"

⋯

where is the rate of producing (i + 1) excitons from (i) hot excitons and

(1)

is the cooling rate via all non-MEG processes. We assume the cooling rate of single- and multi-excitons are

!

' equivalent. The quantity can be related to a previously introduced τMEG by & ,

while in the case of intraband relaxation dominated by longitudinal optical (LO) phonon emission, ! can be connected to ( ! by ! )*+ ' ( ! , where hvLO is the LOphonon energy. Each successive term in Eq. 1 is only valid when energy conservation is met, thus when the photon energy is greater than 3 all three terms may contribute. When ) , 4 higher order terms are necessary and can be included by expanding the series. To evaluate Eq. 1, we assume

!

%

is independent of excess photon energy51 while increases quadratically above its %

threshold energy, and is defined by the relationship: - 7 ACS Paragon Plus Environment

% # ฀ ). ⁄/ "

where

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%

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%

/. ) − / , and P describes the competition between and

!.

Equation 1 can

then be expressed as a function of excess energy and P (MEG characteristics for a variety of P values are plotted Fig. S7). The MEG efficiency can be expressed as, -⁄1 -.50 In the case where ≪ 1 the QY vs. ) would not exhibit a staircase characteristic but rather the QY would exceed 1 at a threshold photon energy, ) > 2 and then the QY would increase linearly (see Fig. S7). Therefore, the fact that, in our data, the MEG QY exhibits the staircase behavior but with overall QYs less than 2 (between 2Eg and 3Eg) suggests that some fraction of the initial photoexcited carriers undergo MEG with near 100% efficiency while the remaining photoexcited carriers do not undergo MEG or do so with lower efficiency. In deriving Eq. 1 we assume that all photons absorbed with energy greater than 2Eg can undergo MEG. However, we can model our data using Eq. 1, by allowing for only a fraction of the excited carriers to undergo MEG. Using such an analysis, we find that ~ 25% of the carriers undergo MEG with 0.98 (Fig. 3, dashed-blue trace). To explain these results, we consider that there is a distribution of electron-hole states that are produced according to the transition dipole matrix element, which is summed over all possible transitions. The energy of an absorbed photon is distributed between the conduction and the valence bands according to ) /

. . where . is the excess energy (energy greater than the band edge) of the electron (e) or hole (h). For efficient MEG, . or . should be equal or greater than (minimum

energy to create an electron-hole pair). In spherical PbS and PbSe QDs, due to nearly equal effective masses in the valence and conduction bands, optical selection rules suggest that only states where the excess energy is equally distributed ( . . are allowed and thus the minimum is 3 . In contast, achieving a threshold of 2Eg requires that the excess energy /

/

reside in either the photoexcited electron or hole ( . and . 0). The Janus-particle nature of these NCs modifies this selection rule to allow for a lower threshold energy and a higher efficiency. We propose here that a faction (25%) of absorbed photons produce photoexcited holes with . ≥ which is near or above the threshold (Fig. 4, left panel, black

circles) while the remainder of photoexcited carriers are excited into states with either insufficient excess energy (Fig. 4a, gray circles) or the excess energy is supplied to the electron, which cools quickly to the bandedge.


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Figure 4. Proposed mechanism for enhanced MEG in PbS|CdS NCs. (a) Excitation with photon energies near 2Eg produces both excitation (grey circles) where neither the electron nor hole has sufficient energy to initiate MEG and excitations where there is enough excess energy in the hole to undergo MEG. (b) Hot holes quickly localize to the PbS|CdS interface states (dashed-green lines). (c) Those carriers with sufficient excess energy can eventually undergo MEG, as cooling to the band edge is hindered.

In our model, photoexcited electrons quickly cool to the CB edge (Fig. 4b). While photoexcited holes quickly (within ~1 ps) localize within PbS|CdS interface states (Fig. 4b) rather than cooling to the VB band edge. Such fast trapping times are consistent with literature reports for CdS NCs in which under-coordinated sulfur surface/interface atoms introduce mid-gap states, resulting in efficient trapping and surface localization of photoexcited holes with >99% efficiency on a ps time-scale.52-55 Those carriers trapped at the interface and with sufficient excess energy eventually undergo MEG with near 100% efficiency while the carriers that do not have sufficient energy either cool to the band-edge or undergo radiative or nonradiative recombination (Fig. 4c).


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Figure. 5. (a) Density Functional Theory density of states (DOS) calculations of small model-size Pb61Cd82S128Cl30 NCs. The inverse participation ratio (IPR, shown as a vertical-bar spectrum) highlights the density of the localized states near the valence band-edge. (b) Selective square moduli of single particle electronic state wave functions corresponding to states labelled on the left hand side. Dark gray spheres represent lead, small blue spheres represent cadmium, yellow atoms represent sulfur, and large blue spheres represent chloride.

Nature of Interface States To gain further insights into the nature of the interface states, we performed density functional theory (DFT) calculations on small model-size chloride-terminated PbS|CdS NCs with composition Pb61Cd82S128Cl30 corresponding to a total diameter of 2.5 nm and an epitaxial interface (structural models and detailed computational methods and results can be found in supporting information, Fig. S4). We find that the lowest energy CB states reside primarily within the PbS domain of the NC (Fig. 5b, right), in agreement with the assignment of the nearIR PL feature. Note, there are CdS localized surface states near the CB edge (Fig. 5a) which could act as electron traps. For the VB region, we find a multitude of localized interfacial states extending ~1 eV below the VB edge, which have both PbS and CdS character (Fig. 5b, left). These states are similar to those observed computationally at the interface of PbSe/CdSe nanoheterostructures.47 States deeper in the VB are delocalized across the entire nanostructure and penetrate into both CdS and PbS domains (Fig. 5b). The spatial localization of the states was measured by computing the inverse participation ratio (IPR) (Fig. 5a). States with high IPR are localized while delocalized states have low IPRs. It is not straightforward to quantitatively extrapolate our computational results to the experimentally relevant NC sizes (~4 nm diameter). However, we can gain some insights by recalling that in our recent study on PbSe QDs embedded in CdSe matrices we found that the 10 ACS Paragon Plus Environment

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interfacial gap states remained at the same energetic position with increasing QD size and they eventually merged with the continuum of delocalized occupied states (see SI of Ref. 47). On the other hand, the band-edge delocalized states were strongly size dependent. Therefore, with increasing QD diameter the delocalized PbS VB states are expected to become shallower than the localized interface states. We thus suppose that the localized interface states calculated here could appear near the delocalized PbS VB edge states and would extend towards the CdS VB edge. These interface states (represented by dashed-green lines, Fig. 2b) are likely responsible for asymmetric broadening of the PL band on its high-energy side (Fig. 2a, dashedgreen lines). Due to their mixed PbS and CdS character they do not strongly overlap with the band-edge states of pure PbS, thus supressing intraband relaxation and facilitating MEG. Indeed, the computational analysis of MEG56,57 indicates that the inclusion of the localized interface states greatly enhances the MEG rate (supporting information, Fig. S8). There are at least three scenarios that can explain the production of a hole with sufficient excess energy (/ 2 to undergo MEG. (1) The presence of an internal electric field. In DFT calculations, we find our model Janus nanoparticle has a nonzero dipole moment of ~35 Debye, suggesting that there is an internal electrostatic field, consistent with previous literature.44 This field can modify absorption selection rules allowing for the photoexcitation of states where the photon excess energy is transferred to the hole.58 The presence of an internal field can also increase the MEG rate.59,60 Eshet et al. calculated MEG efficiencies for type-II QR heterostructures and found that a built-in field is important for MEG for two reasons59 i) the density of multiexciton states at ± above (below) the PbS domain conduction (valence) band increases with increasing built-in electric field, and ii) the built-in field drastically increases the effective Coulombic coupling between photoexcited charge carriers at photon energies at and above 2 . (2) Electronic coupling between the CdS and PbS domains forms new states near 2Eg. There is evidence for such states in our DFT calculations which indicate that states deep in the valence band have both PbS and CdS character. (3) Excitation of a pure-phase PbS intrinsic state with excess energy deposited solely in the valence band hole, followed by the rapid capture at the CdS|PbS interface rather than rapid cooling to the bandedge. Such states in pure-phase PbSe QDs were recently observed by Spoor et al.60,61 These authors find optically allowed higher lying transitions that involve the band-edge electron (L4-6) or band-edge hole (L5-7). Excitation of the L5-7 transitions that lie close in energy to 2Eg would produce a hole with sufficient excess energy to undergo MEG. In the case of CdS|PbS heterostructures, interfacial states can capture the hot-hole providing a pathway to subsequently undergo MEG. The competition between carrier cooling via phonon emission or other relaxation channels that can compete with MEG has received considerable attention in the literature. Theory predicts that quantum-confinement should induce a phonon bottleneck in semiconductor NCs, decreasing the hot exciton cooling rate, 6778, compared to bulk semiconductors. In practice, intraband cooling is extremely fast in NCs (ps to subpicosecond timescale), which can be attributed to multiphonon scattering processes where the exciton couples to vibrational modes 11 ACS Paragon Plus Environment

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of the crystal lattice27 or the surface bound organic ligands.62,63 For example, a study comparing the MEG efficiency of PbS and PbSe QDs concluded that 6778 was faster for PbS because it has a stronger electron-phonon coupling.27 Another example of slowed 6778 for enhanced MEG is the quasi-type II PbSe/CdSe core/shell heterostructure QD system.35 Slowed hot-hole cooling was attributed to decoupling of higher-energy VB states, confined primarily to the shell region, from the lower-energy states mostly localized in the core. The hole cooling is further reduced due to a wide energy gap developing between the core- and shell-based states at large shell thicknesses (quasi phonon bottleneck).

Figure 6. Time-resolved PL spectroscopy of NCs. Time-resolved PL kinetics detected for emission in the visible spectral range for (a) original CdS NCs and (b) 50:50 PbS|CdS Janus-like NCs. Same for NIR emission (for (c) completely cation-exchanged PbS QDs and (d) 50:50 PbS|CdS Janus-like NCs. We find that the NIR emission from the PbS domain remains unchanged for the Janus-like NCs compared with the completely exchanged PbS NCs; however, the visible emission associated with a trapped hole is much longer lived for the Janus-like NCs compared to that for the original CdS QDs, suggesting slowed hole cooling in the heterostructure NCs due to the emergence of the PbS|CdS interface states. Transient PL data for additional samples can be found in Fig. S10.

Here we find additional mechanism for slowing carrier cooling: interactions with localized interfacial states. The cooling rates in Janus-like NCs were studied through time-resolved PL (TRPL) spectroscopy. PL lifetimes from Janus-like NCs were obtained in both the visible and near-IR (NIR) spectral ranges (See Fig. S11 for time-integrated PL spectra) and are compared to lifetimes of original CdS QDs and completely cation-exchanged PbS QDs. The CdS QD PL transient (Fig. 6a) is a result of a convolution of two distinct radiative recombination mechanisms: a fast band-to-band pathway with a lifetime of & 13.3 ns and a longer lived trapstate pathway with a lifetime of & ≅ 120 ns. In contrast, TRPL lifetimes of visible emission from PbS|CdS Janus-like NCs (from the interface states) are nearly an order of magnitude longer— 12 ACS Paragon Plus Environment

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regardless of whether the excitation energy is above (Fig. 5, blue trace) or below the CdS band gap (Fig. 6b, red trace). Examining the NIR PL intensity of the completely cation-exchanged PbS QDs (Fig. 6c) reveals a lifetime of & 2.17 μs while the 50:50 PbS|CdS Janus-like samples show similar lifetimes, which vary only slightly with excitation energy (Fig. 6d). Therefore, the interfacial states observed in the Janus-particles have character of both the PbS (long lifetime) and CdS (spectral position) components and thus are distinct from the pure-phase CdS surface traps. These states clearly slow the carrier relaxation to the PbS band-edge states. Furthermore, excitation at 2 eV, which is below the CdS band-edge, still produces PL from these interfaces states but at an overall lower emission efficiency (supporting Fig. S5). PbS|CdS optoelectronic devices Unlike previously mentioned PbSe/CdSe core/shell QDs, Janus-like NCs present the opportunity to be incorporated as the active layer in optoelectronic devices. This is due to the fact that within Janus-like NCs, either of the photogenerated charge carriers can access a particle surface and can, therefore, be extracted in solar cells. This is a significant advantage over PbSe/CdSe core/shell structures where photogenerated holes are localized in the QD core and cannot be extracted. To verify that the MEG effect can lead to enhanced photocurrent in films of Janus NCs, we apply a transient photocurrent (TPC) spectroscopy implemented with optically gated Auston-switch-like structures. These devices represent a planar transmission line with a narrow gap bridged by a photoconducting materials.64 Excitation of the photocondutor launches a photocurrent transient whose temporal evolution is controlled by the dynamics of photoinjected carriers. This technique provides high (~50ps) temporal resolution and it has been applied previously to detect and quantify MEG yields in PbSe-QD films 38 and to investigate mechanisms of early time photoconductance in these systems.65 To incorporate Janus PbS|CdS NCs in an Auston-switch architecture, we deposit them layer-bylayer onto a quartz substrate with gold plating on its back side. To increase inter-NC electronic coupling, each successive NC layer is treated in a 1 mM solution of EDT in acetonitrile. The top gold contacts, separated by a 100 µm gap, are applied by thermal evaporation. The TPC is measured with excitation below (Fig. 7a, red trace) and above (Fig. 7a, blue trace) the MEG threshold using 100-fs pulses at 1.55 and 3.1 eV, respectively. As in previous TPC studies of MEG, this effect is observed as a fast intial relaxation component associated with Auger recombination of biexcitons produced by single photons. In fact, the “fast” time constant observed in TCP measurements with 3.1 eV excitation (~170 ps) is consistent with the biexciton lifetime (~130 ps) inferred from purely optical measurements of these Janus NCs dispersed in solution. These observations indicate that the MEG characteristics are maintained in photoconductive arrays of Janus PbS|CdS NCs. We also incorporated Janus NCs into solar cells and tested their performance under simulated solar illumination. To fabricate these devices, we deposited PbS|CdS NCs onto TiO2-coated FTO/glass substrates using layer-by-layer dipcoating in which each layer was treated in a 1 mM 13 ACS Paragon Plus Environment

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solution of 1,2-ethanedithiol (EDT) in acetonitrile to displace native oleate ligands. A MoOx/Au top contact was then thermally evaporated to complete the device. Top performing solar cells achieved nearly 3% efficiency with a 20.5 mA/cm2 short-circuit current density and 307 mV open-circuit voltage (Fig. 7b). The EQE (Fig. S12) peaks above 75% for NCs with a 0.9 eV band gap. The fact that the EQE does not exceed 100% for these proof-of-principle devices is likely due to imperfect carrier transport and collection at the electrodes. We expect that future efforts on optimizing the device architecture and the NC surface chemistry would lead to EQEs in exess of 100% as was previously observed for pure-phase QDs.5-8

Figure 7. Photocurrent-based measurements of films of PbS|CdS Janus-like NCs (a) Representative TPC traces recorded using 1.55 eV (red) and 3.1 eV (blue) excitation with low fluence (~0.1) 100 fs pulses. Traces are normalized so as to match their long-time (single-exciton) signals, which allows us to clearly visualize the fast Auger decay of biexcitons produced by MEG in the case of 3.1 eV excitation. (b) A representative current density-vs.-voltage characteristic of a solar cell achieving a nearly 3% power conversion efficiency, based on EDT-treated 50:50 PbS|CdS Janus-like NCs. Inset shows the solar-cell architecture.

Conclusion and Outlook From the above investigation of the PbS|CdS Janus-like heterostructures, we can derive several design rules for the next generation of engineered nanostructures that may achieve even higher MEG efficiency. Namely, the heterojunction nature and a slowing of the carrier cooling are essential components of these structures. Hetero-nanostructures show promise for continued exploration for MEG applications having demonstrated a pathway to break the selection limitations imposed on single-phase PbE QDs. Additionally, Janus-like particles contain a spatial asymmetry not present in core/shell structures, which can enhance the Coulombic coupling similar to that observed in 1D quantum confined structures, and further, facilitate e-h spatial separation which slows down recombination and simplifies carrier extraction. Slowing carrier relaxation due to interaction with a manifold of interfacial states is an approach for redirecting photogenrated carriers towards the desired MEG channel. One possible future direction might 14 ACS Paragon Plus Environment

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involve lowering the VB of the heterostructure material. For example, in the CdSe/PbSe system it would be desirable for the VB of the CdSe to be low enough, relative to that of PbSe, so as to interact with hot-carriers (relative to the PbSe VB) and further achieve slowing of the carriers.43 Thus Janus-like heterostructures of CdSe/PbS, CdSe/PbSe, and CdTe/PbTe are all of potential interest. Another potential area of interest would be the introduction of specific localized states within the VB or CB through either ligand/QD interactions or the introduction of specific impurity atoms with localized states high in the VB or CBs. Through exploration of NC core compositions and surface chemistry, coupled with improved solar cell device architectures, we believe Janus-like heterostructure NCs can provide a platform for the continued development of highly efficient PV technology utilizing MEG. Materials and Methods A. Materials PbCl2 (99.999%), oleylamine (tech. grade, 70%), oleic acid (tech. grade, 90%), 1-octadecene (tech. grade, 90%), cadmium oxide, (NH4)2S 40-48 wt. % in water, tetrachloroethylene (TCE, ≥99.9%), acetonitrile (99.8%), 1,2-ethanedithiol (EDT, ≥98.0%), molybdenum(VI) oxide (99.99%), gold (99.999%), hexane (95%), and ethanol (≥99.5%) were purchased from Sigma Aldrich. All the chemicals were used as received. B. Nanocrystal Synthesis and Purification PbS|CdS Janus-like heterostructure NCs were synthesized according to Zhang et al.37 First, CdS QDs were synthesized and purified using a modified procedure from Peng et al.66 The CdS QDs, suspended in 1-octadecene at ~40 mg/mL, were swiftly injected into a hot (140 - 170 °C depending on desired final NC size) Pb precursor solution, which consisted of 10 mL of oleylamine and ~3 mmol PbCl2 (varied depending on desired final NC composition ratio). The reaction flask was submerged in an ice water bath directly after injection of the CdS QDs, followed by injection of 10 mL hexanes. At 40 °C, 8 mL of oleic acid was injected for improved solution stability of the partially exchanged PbS|CdS Janus-like NCs. The product was washed two times under ambient conditions using hexane/ethanol as the solvent/antisolvent pair, and was stored as a solid. The NCs were suspended in hexane for thin film formation or tetrachloroethylene for spectroscopic measurements. C. Nanocrystal Characterization UV-Vis Absorbance: Optical absorbance spectra were collected using a Shimadzu-3600 spectrometer. Steady-State NIR Photoluminescence (PL): NIR PL spectra were collected using a home-built PL instrument consisting of an excitation source, emission monochromator (Photon Technologies International), and a two-stage thermocouple-cooled extended InGaAs detector (Judson). The higher energy excitation source used to excite both PbS and CdS domains was a chopped Hg-Xe lamp passed through a monochromator and band-pass filter. The lower energy excitation source used to exclusively excite the PbS domain was a 780 nm LED (ThorLabs M780L3), which was 15 ACS Paragon Plus Environment

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driven by a 15 V square wave at 25 Hz using a Stanford Research Systems (SRS) DS335 function generator. The detector signal was amplified using a SRS SR530 lock-in amplifier, and all spectra were corrector for grating and detector efficiencies using a calibrated lamp. Photoluminescence Excitation (PLE): PLE spectra were collected using a modified Horiba JobinYvon Fluorolog 3. The excitation source was a chopped 450 W Xe lamp passed through a monochromator and band-pass filter. The detector unit consisted of a liquid nitrogen-cooled Si CCD coupled to an iHR320 emission monochromator. X-ray diffraction (XRD): XRD was performed on a Bruker D8 Discover diffractometer using Cu Kα radiation (λ =1.54 Å). Transmission electron microscopy (TEM): TEM studies were completed using FEI T30 at 300 kV. TEM grids were prepared by dropcasting a dilute solution of NCs in a hexane mixture onto the carbon coated copper grids. X-ray Fluorescence (XRF): XRF was performed on a Fischer XDV-SDD XRF Time-resolved Photoluminescence (TRPL): TRPL measurements were made using a supercontinuum fiber laser (Fianium, SC-450-PP) operating at 100 kHz as the excitation source. The excitation wavelength was chosen using an acousto-optic tunable filter to be 450 nm or 600 nm with of pulse energy of approximately 1 nJ. A streak camera for NIR detection (Hamamatsu C11293-02, 859-1600 nm) or visible detection (Hamamatsu C10910-04, 400-850 nm) was used to detect time-resolved spectra. The instrument response function depends on the time window size, but at best is 100 ps. Decays were fitted using a multiexponential function convoluted with an instrument response function, which was measured by scattering the excitation beam from a Ludox solution. Transient Photocurrent Spectroscopy (TPC): Devices for TPC were fabricated by spincoating PbS|CdS Janus-like NCs in a layer-by-layer fashion on top of quartz. Each successive layer was treated by briefly (~1-2 s) soaking films in a 1 mM solution of EDT in acetonitrile. This process was repeated until film thickness reached ~150 nm and films were then annealed in an oxygenfree environment at 120 C for 20 minutes. The gold contacts were then deposited by thermal evaporation. TPC measurements were made by exciting the device with ~100 fs pulses obtained from an amplified Ti:sapphire laser system. Pulses with photon energy of 3.1 eV were obtained by frequency doubling of the laser output at 1.55 eV using a BBO crystal. The transient photocurrent signal was measured through a 20 GHz sampling oscilloscope. Steady-State Photoluminescence (PL) Spectroscopy: Steady-state PL measurements were made using a Thorlabs fiber-couple LED as the excitation source pulsed at 10 Hz using a ThorLabs DC2200 LED driver. Visible detection was made using an Ocean Optics OceanFX spectrometer and NIR detection was made using an Ocean Optics NIRQuest spectrometer. Visible and NIR spectra were stitched with a LabVIEW program developed in-house using an Ocean Optics HL2000-HP blackbody lamp for broad-band calibration. D. Transient Absorption (TA) Spectroscopy


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TA spectra were collected using a home-built instrument. A Continuum Integra-C regeneratively amplified Ti:sapphire laser with ~3.5 W, 1 kHz, and ~100 fs pulse width output at ~800 nm is split into two beams; pump and probe. The pump beam is directed into a Palitra-Duo-FS:FS optical parametric amplifier that is capable of producing excitation wavelengths from 27022,000 nm and is modulated at 500 Hz through an optical chopper to block every other laser pulse. The probe beam passes through a multipass delay stage that can achieve up to ~4.5 ns of pump-probe delay, and is focused onto a sapphire crystal to produce a broadband Vis-NIR spectrum from 800-1600nm. The probe is passed through a continuously variable neutral density filter and a fraction is picked off to be used as a reference, which accounts for probe beam intensity fluctuations and improves signal-to-noise. The pump and probe beams are then overlapped at the NC solution sample, which is vigorously stirred to eliminate photocharging effects. NIR photodiode arrays (Ultrafast Systems) are used to detect the probe and reference beams for data acquisition. E. Solar Cell Fabrication Solar cells were fabricated on glass/FTO substrates supplied from Thin Film Devices. A TiO2 solgel layer is spincoated onto FTO substrates to form a 70 nm layer following previous reports.67 Janus-like NCs are then deposited on top of the TiO2 using a layer-by-layer dipcoating method. Substrates are dipped in a solution of Janus NCs that are dispersed in hexane at ~10 mg/mL before being dipped into a solution of 1 mM EDT in acetonitrile. This process is repeated until the desired film thickness is achieved (typically ~200 nm). Films are then annealed at 120 °C for 20 minutes in an oxygen-free environment. Finally, a back contact of MoOx (20 nm) and Au (100 nm) is deposited by thermal evaporation. F. Solar Cell Characterization Current-voltage characterization of solar cells was performed in a nitrogen-filled glovebox on a Newport solar simulator calibrated using a Si reference diode. Devices were masked to an area of 0.059 cm2. External quantum efficiency (EQE) measurements were done using a Newport Oriel IQE200 in a nitrogen-filled glovebox. Supporting Information The Supporting Information is available free of charge on the ACS Publications website at DOI: Absorbance, PL, PLE, and XRD of Janus-like nanocrystals with varying Pb:Cd ratios, hyperspectral transient absorption surface plots and representative transient absorption data with MEG analysis, structural model and computational results detailing energy alignment, exciton quantum yield plotted for systems of differing MEG efficiency, computed MEG rates for Januslike nanocrystals with and without interface states as well as a varying electric field, additional TRPL plots with time-integrated PL spectra, EQE from the device shown in Fig. 7b. Acknowledgements We acknowledge support from the solar photochemistry program within the Division of Chemical Sciences, Geosciences and Biosciences, Office of Basic Energy Sciences, Office of Science within the US Department of Energy for the synthesis of the Janus particles, ultrafast 17 ACS Paragon Plus Environment

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measurements of MEG and the photolumenesence measurements of the Janus-particles. Support for the TPC MEG studies and the development and characterization of PV devices was provided by the Center for Advanced Solar Photophysics (CASP), an Energy Frontier Research Center funded by the US Department of Energy, Office of Science, Office of Basic Energy Sciences. G.F.P. acknowledges support from the Global R&D program (1415134409) funded by KIAT, MOTIE. M.V. was supported by laboratory Directed Research and Development (LDRD) funding from Argonne National Laboratory, provided by the Director, Office of Science, of the U.S. DOE under Contract No. DE-AC02-06CH11357. This research used computational resources provided by the Los Alamos National Laboratory Institutional Computing Program, which is supported by the U.S. DOE National Nuclear Security Administration under Contract No. DEAC52-06NA25396, and the National Energy Research Scientific Computing Center (NERSC) which is supported by the Office of Science of the U.S. DOE under Contract No. DE-AC0205CH11231.The computational efforts of FG and GG were supported by MICCoM, as part of the Computational Materials Sciences Program funded by the U.S. Department of Energy, Office of Science, Basic Energy Sciences, Materials Science and Engineering Division through Argonne National Laboratory, under Contract No. DE-AC02-06CH11357. Part of this work was authored in part by Alliance for Sustainable Energy, LLC, the manager and operator of the National Renewable Energy Laboratory for the U.S. Department of Energy (DOE) under Contract No. DEAC36-08GO28308. The views expressed in the article do not necessarily represent the views of the DOE or the U.S. Government. The U.S. Government retains and the publisher, by accepting the article for publication, acknowledges that the U.S. Government retains a nonexclusive, paidup, irrevocable, worldwide license to publish or reproduce the published form of this work, or allow others to do so, for U.S. Government purposes.

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