Article pubs.acs.org/JACS
Giant PbSe/CdSe/CdSe Quantum Dots: Crystal-Structure-Defined Ultrastable Near-Infrared Photoluminescence from Single Nanocrystals Christina J. Hanson,† Nicolai F. Hartmann,† Ajay Singh,† Xuedan Ma,† William J. I. DeBenedetti,‡,○ Joanna L. Casson,§ John K. Grey,∥ Yves J. Chabal,⊥ Anton V. Malko,‡ Milan Sykora,§ Andrei Piryatinski,∇ Han Htoon,† and Jennifer A. Hollingsworth*,† J. Am. Chem. Soc. 2017.139:11081-11088. Downloaded from pubs.acs.org by KAOHSIUNG MEDICAL UNIV on 09/29/18. For personal use only.
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Materials Physics and Applications Division: Center for Integrated Nanotechnologies, Los Alamos National Laboratory, Los Alamos, New Mexico 87545, United States ‡ Department of Physics, The University of Texas at Dallas, Richardson, Texas 75080, United States § Chemistry Division, Los Alamos National Laboratory, Los Alamos, New Mexico 87545, United States ∥ Department of Chemistry, The University of New Mexico, Albuquerque, New Mexico 87131, United States ⊥ Department of Materials Science and Engineering, The University of Texas at Dallas, Richardson, Texas 75080, United States ∇ Theoretical Division, Los Alamos National Laboratory, Los Alamos, New Mexico 87545, United States S Supporting Information *
ABSTRACT: Toward a truly photostable PbSe quantum dot (QD), we apply the thick-shell or “giant” QD structural motif to this notoriously environmentally sensitive nanocrystal system. Namely, using a sequential application of two shell-growth techniques partial-cation exchange and successive ionic layer adsorption and reaction (SILAR)we are able to overcoat the PbSe QDs with sufficiently thick CdSe shells to impart new single-QD-level photostability, as evidenced by suppression of both photobleaching and blinking behavior. We further reveal that the crystal structure of the CdSe shell (cubic zinc-blende or hexagonal wurtzite) plays a key role in determining the photoluminescence properties of these giant QDs, with only cubic nanocrystals sufficiently bright and stable to be observed as single emitters. Moreover, we demonstrate that crystal structure and particle shape (cubic, spherical, or tetrapodal) and, thereby, emission properties can be synthetically tuned by either withholding or including the coordinating ligand, trioctylphosphine, in the SILAR component of the shell-growth process.
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INTRODUCTION
Despite their potential, practical application of PbSe QDs has been stymied by a lack of long-term stability in air. Specifically, PbSe QDs kept in air, at room temperature, and, especially, in the presence of light are known to exhibit rapid photooxidative degradation, including both loss in PL intensity and spectral blue-shifting as the effective QD size shrinks with oxidation of surface lead and selenium.12−14 Significant effort has been dedicated to enhancing the stability of PbSe QDs. Approaches have included inorganic shelling, e.g., to form PbSe/CdSe core/ shell QDs,15 and inorganic or organic surface treatments. The latter have entailed halide passivation using PbX2 (X = Cl, Br, I) precursors during PbSe synthesis as well as postsynthesis exposure to molecular chloride16 or halide salts,17 or improved organic passivation using phosphonic acid ligands.18 Enhanced stability has been confirmed by following PL and absorption
Lead chalcogenide quantum dots (QDs) are infrared active semiconductor nanocrystals with wide-ranging potential applications, including field effect transistors,1 photodetectors,2 photovoltaics,3 light emitting diodes,4 and as gain media for realization of amplified spontaneous emission5,6 and lasing.7 In particular, PbSe QD photoluminescence (PL) wavelengths can be tuned from the near-infrared8 to the mid-infrared (∼4000 nm) by simple particle-size tuning.9 The extreme sizedependent spectral tunability (compared, for example, to CdSe QDs) is attributable to PbSe’s large exciton Bohr radius (46 nm).9 Emission in the near-infrared is compatible with common fiber-optics-based telecommunication technologies (1300−1550 nm) and overlaps with a second optical transparency window for biological tissue (up to ∼1350 nm),10 while mid-infrared sources are targeted for novel lasers needed in medical diagnostics, environmental monitoring, and military applications.11 © 2017 American Chemical Society
Received: April 12, 2017 Published: July 18, 2017 11081
DOI: 10.1021/jacs.7b03705 J. Am. Chem. Soc. 2017, 139, 11081−11088
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cycles; Figure 2; Supporting Information, Synthetic Methods and Table S1). Partial cation exchange of lead for cadmium yields CdSe shells with a zinc-blende (zb) crystal structure,15 selectively forming {111} rocksalt (rs)-PbSe/zb-CdSe interfaces.29 However, we find that shell growth by the SILAR process can be tuned to form either cubic zinc-blende or hexagonal wurtzite (w) CdSe, depending on the identity of the Se precursor. Namely, if Se dissolved in 1-octadecene (Se-ODE) was used, the PbSe/CdSe/CdSe g-QDs were cubic in shape and possessed predominantly zinc-blende crystal structure (Figure 1a, b, and g; 10.1 ± 0.8 nm across diagonal of cubes, Figure S2), but if, instead, Se dissolved in trioctylphosphine (Se-TOP) was employed, the resulting thick-shell nanocrystals were quasispherical or hexagonal bipyramidal in shape (Figure S3) and exhibited wurtzite crystal structure (Figure 1c, d, and g; 8.1 ± 0.9 nm diameter, Figure S4). Thus, despite a common rock salt PbSe core, different SILAR shell growth conditions produced distinct shapes and crystal structures. It has been noted previously for the CdSe/CdS core/thickshell system that the crystal structure of the core dictates the crystalline phase of the shell.30−32 In the instances describing wurtzite shell growth on a wurtzite core, this result is unremarkable, as wurtzite is the thermodynamically favored phase for both CdSe and CdS.26,33,34 However, in the case of zinc blende heteroepitaxial growth, different rationales have been put forth to explain the influence of core structure on shell growth, which encompass both thermodynamic and kinetic arguments, such as slow shell growth (afforded by slow addition of precursor and low reactivity of the S precursor)32 and low reaction temperature (enabled by “dynamic” surface passivation and sufficiently reactive precursors),31 respectively. Here, we find a situation where the shell can either grow epitaxially (zb-CdSe on zb-CdSe) or organize into the bulk thermodynamic phase (w-CdSe on zb-CdSe). The only difference between the reaction conditions affording the two results in our case is the absence or presence of TOP in the selenium precursor. We surmise that the presence of this additional coordinating solvent (representing from ∼1% of the total reaction volume after 1 SILAR shell-addition cycle to ∼30% after 10 SILAR shell-addition cycles; see the Supporting Information for details) can have several effects that may lead to growth of a hexagonal shell. First, TOP can modify Se (and, possibly, Cd) reaction rates by stabilizing the solvated form of the precursor(s) and, once Se has added to the growing nanocrystal surface, passivating Se surface sites, effectively reducing reactivity toward Cd adatoms.35,36 Second, TOP and other strongly coordinating solvents can render the metal− chalcogenide (Cd−Se) bond-making process reversible,36 favoring the formation of the lowest energy surface. Finally, TOP may further lower the surface energy of the wurtzite facets compared to zinc blende facets. Under the above conditions, the structure-directing ability of the core is compromised in favor of the lowest-energy, thermodynamic phase. In contrast, presumably faster and/or irreversible (less-dynamic) shell growth in the absence of TOP permits the shell structure to be dictated by that of the core, despite the relatively high temperatures employed for SILAR growth. We have observed similar nonepitaxial shell growth previously in the case of zb-InP/w-CdS core/thick-shell QDs.27 Although not employing TOP, a coordinating primary amine (oleylamine) comprised a significant percentage of the solvent
trends as a function of time and storage conditions and by demonstrating improved device performance.18 Notwithstanding improvements in the environmental stability of PbSe QDs, to date, PL properties have only been investigated for QD ensembles. Single-dot-level emission properties have been experimentally inaccessible as a result of rapid photobleaching and the very long PL lifetimes (≥1 μs) that are characteristic of these colloidal nanocrystals.19,20 Even in the case of relatively faster emitting PbS QDs (∼500−1000 ns solution-phase lifetimes21), radiative recombination of excitons is sufficiently slow to result in a weak PL signal, which when combined with low QD quantum yields (QYs), nanocrystal instability, and high detector background noise, e.g., in the case of InGaAs array or avalanche photodiode (APD) detectors, prevents measurement of single QD emission.22 Observation of single-dot-level PbS/CdS QD emission has required advanced detectors, such as the superconducting nanowire single photon detector (SNSPD).21 The ability to measure and, ultimately, make use of photons emitted by single PbSe QDs would be significantly enhanced if rapid photobleaching could be suppressed. For example, we have observed that, when these QDs are spread onto glass substrates at low concentrations and interrogated at room temperature with laser excitation, they photobleach within seconds, rendering both single-QD PL studies and applications impractical. It is now well-known that photobleaching and even fluorescence intermittency (blinking) is suppressed or eliminated in the case of some core/shell QDs for which the protective shell is very thick.23−27 Here, we show for the first time that the notoriously photounstable PbSe quantum emitter can be rendered ultraphotostable with the addition of a thick CdSe shell, exhibiting nonphotobleaching emission at the level of QD clusters and even for single nanocrystals spread bare on a glass slide and interrogated for over 1 h using high laserexcitation power densities (50 to >300 W/mm2). Unexpectedly, we find that the shell crystal structurezinc blende or wurtzitedetermines the efficiency of PbSe/CdSe g-QD PL and, thereby, our ability to detect single-dot emission. We further show that cubic or hexagonal shell growth, as well as shell shape, can be synthetically directed.
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RESULTS AND DISCUSSION Thick-shell or “giant” PbSe/CdSe QDs (PbSe/CdSe g-QDs) were synthesized using sequentially applied partial-cation exchange15 and modified successive ionic layer adsorption and reaction (SILAR)26,28 shell-growth procedures. Briefly, for example, PbSe QD cores (5.6 ± 0.5 nm diameter, 1561 nm 1S exciton absorption peak; Figure S1) suspended in toluene were mixed with excess cadmium oleate, heated to 90 °C, and allowed to react for 2.5 h. This process resulted in a partial exchange of Pb for Cd, shrinking of the PbSe core, and formation of a CdSe shell (Supporting Information, Synthetic Methods).15 In this way, the total particle diameter remains approximately constant, while the core size as determined by known diameter/PL-energy correlations becomes smaller. Here, the new reduced core size calculated on the basis of the shifted PL peak position was ∼4.1 nm (with corresponding postexchange PL peak maximum: ∼1410 nm), implying a CdSe shell thickness of 1 nm zb-CdSe shells (1.7 nm). Subsequent SILAR shelling using the Se-TOP precursor yielded PbSe/zb-CdSe/w-CdSe nanocrystals similar to those observed for the smaller-core preparations; however, the w-CdSe formed as tetrapoidal arms (Figure 1e,f). We speculate that arm formation, rather than conformal shell growth, was promoted in this case due to the presence of well-established zinc-blende CdSe facets, only characteristic of sufficiently thick CdSe shells (i.e., ∼5 MLs vs 20% of the 4000 s observation time (Figure 5c). By comparison, 0% of a population of the well-known, high-quality CdSe/ZnS QDs
Figure 4. Theoretical contour plots showing shell thickness and core radius dependencies for PbSe/CdSe (a) electron−hole overlap integral and (b) optical bandgap energy (as 1S wavelength in μm). Regions where specific types of electronic structure are expected to prevail are also indicated. Band diagrams for PbSe/CdSe QDs representing two core/shell combinations predicted to have quasi type II character: (c) larger core (4 nm diameter) and thinner shell (0.75 nm) (indicated by a square) and (d) smaller core (2 nm diameter) and thicker shell (3 nm) (indicated by a circle). For parts c and d, the dashed lines mark 1S electron and hole state energies that determine the optical gap (indicated by arrow). The associated charge carriers’ density distributions are superimposed, indicating a quasi type II localization regime.
Figure 5. (a) Video still image of photoluminescence (PL) from single rs-PbSe/zb-CdSe/zb-CdSe g-QDs dispersed on a glass substrate (total shell thickness: 3.0 nm; 8.6 ML); the scale bar represents 1 μm. (b) Long PL−time trace for the g-QD circled in part a. The red dashed line indicates the background level. (c) On-time histogram (122 gQDs analyzed). 11085
DOI: 10.1021/jacs.7b03705 J. Am. Chem. Soc. 2017, 139, 11081−11088
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CdSe QDs were suspended in hexanes, with concentration assessed as before by determining absorbance as the PbSe 1S peak.2,3 The PbSe/CdSe core/shell QDs synthesized using the cation exchange approach for shell growth were subjected to further shell growth using the method of successive ionic layer adsorption and reaction (SILAR) growth. Separately, stock solutions of 0.2 M Se in ODE and 0.2 M Cd-oleate in ODE (1:4 molar ratio Cd:oleic acid) were prepared and used as precursors for the SILAR monolayer-bymonolayer additions.4−6 Alternatively, for SILAR growth employing a Se-TOP precursor, a 0.2 M Se solution in TOP was prepared by dissolving Se shot (0.79 g) in TOP (50 mL) by first degassing the mixture under a dynamic vacuum at 90 °C for 1 h and then heating at 180 °C until a clear solution was obtained. ODE (5 mL) and OA (1 mL) were then degassed under a dynamic vacuum at 100 °C for 1 h. The temperature was lowered to 30 °C and the flask opened to flowing argon. PbSe/CdSe QDs (2.5 mL of a 3.91 × 10−6 M solution in hexanes) prepared by cation exchange were injected into the ODE/ OA mixture. Hexanes were removed under a dynamic vacuum, such that the QD concentration was 1 nM. Under flowing argon, the QD solution was slowly heated to 240 °C over ∼10 min. The Cd precursor solution was added in an amount equivalent to 1/2 monolayer of CdSe shell, followed by the equivalent amount of Se precursor, waiting 10 min in between each addition. Precursor addition quantities were calculated for each subsequent monolayer based on the known QD concentration and assuming that each addition of Cd and Se precursor afforded a full monolayer of shell growth and corresponding increase in QD size.4−6 PbSe/CdSe/CdSe QDs (or tetrapods) were extracted at different points in the SILAR growth process and precipitated by adding ethanol. Cycles of suspension and precipitation using hexanes and ethanol, respectively, were applied to remove excess ligands. QDs were suspended in TCE for absorption and emission measurements and in hexanes for the preparation of samples for structural analysis by TEM. Structural Characterization. Nanocrystal size and shape were assessed by transmission electron microscopy (TEM) using a JEOL 2010 instrument operating at 200 kV. High-resolution TEM and annular dark-field scanning TEM (ADF-STEM) images were acquired using an FEI-Titan 80-300 instrument operating at 200 kV. X-ray diffraction was collected on a Rigaku Ultima III diffractometer that uses a Cu Ka (1.5406 A) X-ray source. The Jade 9.0 data analysis program was used for semiquantitative analysis of the relative contributions of wurtzite and zinc blende phases to the XRD powder pattern. Specifically, a whole pattern fitting and Rietveld refinement method was employed. R-fitted values of 20% for similar experimental conditions.25 Examples of blinking studies for QDs emitting in the infrared above 1 μm (toward, for example, the telecommunications relevant wavelengths of 1.3 and 1.55 μm) are rare due to the inherent challenges pertaining to both the materials (weak emission, slow radiative recombination rates, and/or rapid photobleaching) and the techniques used to observe them.19,25 In fact, this report represents only the second demonstration of suppressed blinking behavior for IRactive colloidal nanocrystals, with an InP/CdS core/thick-shell QD being the first.27 And, as previously discussed for CdSe/ CdS and InP/CdS gQDs, the thick shell is a composition for which the resulting core/shell electronic structure has type II character. In these instances, the unusual single-dot level optical stability is observed.50
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CONCLUSION In summary, we have shown that high-temperature SILAR shell growth can be employed for nominally thermally sensitive PbSe QDs if a shell is first applied using the low-temperature cation exchange technique. Furthermore, the structure and resulting optical properties of the core/thick-shell nanocrystals are determined by the choice of Se precursor, i.e., the presence or absence of TOP. Successful application of a thick CdSe shell possessing the cubic zinc-blende crystal structure yields highly stable nanocrystals for which both photobleaching and blinking are suppressed. The enhanced stability and access to PL that is tunable above 1 μm imply potential applications in telecommunications, near-IR photodetectors, and optical tags.
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EXPERIMENTAL SECTION
Materials. 1-Octadecene (ODE, 90%), oleylamine (OAm, 70%), cadmium oxide (CdO, 99.99%), and trioctylphosphine (TOP, 90%) were purchased from Acros Organics. Lead(II) oxide (PbO, 99.9998%), selenium shot (Se, 99%), and oleic acid (OA, 90%) were purchased from Alfa Aesar. Tetrachloroethylene (TCE, 99+% spectroscopic grade), 1,2-dichlorethane (DCE, 99.8%, anhydrous), and diphenylphosphine (DPP, 98%) were purchased from Sigma-Aldrich. All chemicals were used without further purification. Synthesis. PbSe QDs were synthesized and subjected to partial exchange of Pb for Cd using a previously described method.1 For example, 600 mg of PbO, 6 mL of OA, 8 mL of ODE, and 8 mL of TOP were degassed at 110 °C for 2 h. The resulting mixture was then heated under flowing argon to 205 °C to obtain an optically clear solution before cooling to 150 °C. Eight mL of a 1 M TOPSe solution and 50 μL of DPP were injected swiftly into the lead solution, causing an immediate drop in temperature to ∼125 °C and affording nucleation and growth of PbSe nanocrystals. After 2 min, the reaction was terminated by removing the flask from heat and submerging the flask in an ice bath. The QDs were precipitated from the growth solution by adding acetone, and then subjected to two cycles of suspension and precipitation using toluene and acetone, respectively. The processed QDs were dispersed in a minimum amount of toluene and QD concentration determined by measuring the 1S absorbance peak, from which the molar absorptivity can be calculated.2,3 In this case, the solution concentration was 1.07 × 10−5 M. A 10 mL portion of these QDs [5.6 ± 0.5 nm in diameter as determined by transmission electron microscopy (TEM) analysis] was injected into 16 mL of 0.2 M Cd-oleate in ODE (1:4 molar ratio Cd:oleic acid) and allowed to stir at 90 °C for 2.5−12 h, resulting in a CdSe shell ∼0.75−0.90 nm thick by the sacrificial replacement of Pb for Cd. Growth of thicker shells by cation exchange could be realized using the same experimental parameters but starting with larger PbSe cores and exposing these to the Cd solution for longer times (≥12 h). The cation exchanged product was subjected to two cycles of precipitation and resuspension using hexanes and ethanol, respectfully. The PbSe/ 11086
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was fitted with a double exponential function. For each sample, five different positions were measured to obtain an average lifetime. Single-QD measurements were obtained using a ProEM 512 EMCCD (Princeton Instruments) camera with a long pass filter of 900 nm to ensure that only emission in the near-infrared contributed to the observed signal. The detected signal was split from the excitation with a 442 dichroic beam splitter (Semrock Di02-R442-25x36). Samples were excited using a 405 nm continuous wave (CW) laser (330 W/ mm2). To ensure that single QDs were being observed, a QD stock solution was successively diluted, and emission from the resulting solutions spread onto glass substrates was observed until a series of dilutions was reached for which the number of emissive spots on the substrate scaled linearly with the dilution factor. The final ultradilution used for the widefield blinking/bleaching measurements shown here yielded 20 QDs/50 μm2 area. The ultradilute solutions are prepared just prior to deposition onto a glass substrate to avoid QD clustering resulting from ligand destabilization.
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REFERENCES
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ASSOCIATED CONTENT
S Supporting Information *
The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/jacs.7b03705.
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Article
Experimental and modeling details and results of size analyses using TEM images (PDF)
AUTHOR INFORMATION
Corresponding Author
*
[email protected] ORCID
Nicolai F. Hartmann: 0000-0002-4174-532X Ajay Singh: 0000-0002-5168-7522 Xuedan Ma: 0000-0002-3163-1249 Yves J. Chabal: 0000-0002-6435-0347 Han Htoon: 0000-0003-3696-2896 Jennifer A. Hollingsworth: 0000-0003-3099-1215 Present Address ○
W.J.I.D. is currently at Department of Chemistry, Cornell University, Ithaca, New York, 14853.
Notes
The authors declare no competing financial interest.
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ACKNOWLEDGMENTS The work was supported primarily by a Division of Materials Science and Engineering, Office of Basic Energy Sciences (OBES), Office of Science, U.S. Department of Energy (DOE), grant, 2009LANL1096. A.S. was supported by a Los Alamos National Laboratory Director’s Postdoctoral Fellowship and C.J.H. by Laboratory Directed Research and Development funds. Work of the UT Dallas group (W.J.I.D., Y.J.C., A.V.M.) was supported by a DOE OBES grant, DE-SC0010697. Work was performed at CINT, a DOE, OBES Nanoscale Science Research Center & User Facility, with aspects of the work supported by a CINT User Project (U2013A0134). We thank Darrick Williams for obtaining and analyzing XRD patterns as well as Chris Sheehan for conducting STEM-in-SEM. High resolution TEM and low resolution ADF-STEM images were obtained at the National Center for Electron Microscopy (NCEM), Lawrence Berkeley National Laboratory, a user facility supported by the DOE, OBES, under contract no. DEAC02-05CH11231. 11087
DOI: 10.1021/jacs.7b03705 J. Am. Chem. Soc. 2017, 139, 11081−11088
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DOI: 10.1021/jacs.7b03705 J. Am. Chem. Soc. 2017, 139, 11081−11088
