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In Situ Monitoring of Cation-Exchange Reaction Shell Growth on Nanocrystals Annina Moser, Maksym Yarema, Weyde M.M. Lin, Olesya Yarema, Nuri Yazdani, and Vanessa Wood J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.7b08571 • Publication Date (Web): 11 Oct 2017 Downloaded from http://pubs.acs.org on October 17, 2017

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The Journal of Physical Chemistry

In Situ Monitoring of Cation-Exchange Reaction Shell Growth on Nanocrystals Annina Moser, Maksym Yarema, Weyde M. M. Lin, Olesya Yarema, Nuri Yazdani, and Vanessa Wood* Laboratory for Nanoelectronics, Department of Information Technology and Electrical Engineering, ETH Zurich, Gloriastrasse 35, CH-8092 Zurich, Switzerland *[email protected]

ACS Paragon Plus Environment

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ABSTRACT.

We demonstrate how in situ monitoring of the photoluminescence during shell growth around colloidal nanocrystals (NCs) can be used to develop a detailed and quantitative model for this process. We apply it here to study cation-exchange based growth of ZnS on a Cu-In-Se NC to form Cu-In-Se/ZnSe1-xSx alloyed NCs. We determine that this process begins with the Znprecursor binding to the outer layer of the NC followed by diffusion of Zn cations into successive atomic monolayers of the NC. At temperatures below 100°C, Zn cations can only diffuse into the outermost atomic monolayer of the Cu-In-Se NCs. At growth temperatures above 100°C, the second monolayer also becomes thermally accessible and can be filled with Zn cations. Our results provide an understanding of cation-exchange shell growth at the atomic level via optical analysis. The approach and mathematical model described here can be applied to other core/shell nanostructures and allows selection of optimal synthesis conditions to achieve desired core/shell design for a specific application.


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Introduction For over 20 years, it has been known that shell growth on colloidal NCs significantly improves photoluminescence (PL) efficiencies.1 Shells not only passivate electronic surface defects but also alter the electronic confinement, which can displace the electronic density away from the surface and allow versatile tuning of optical and electronic properties. Since then, a considerable amount of research has been devoted to developing new shell synthesis routes and expanding the range of semiconductor material combinations.2-3 Multiple4, gradient alloyed5 and “giant”6-7 shells have led to record PL efficiencies due to the suppression of non-radiative processes, such as Auger recombination.8 Using NCs in solid-state devices introduces additional challenges, such as high electric fields, charge injection, and transport.9-10 Tailoring of the band structure via shell growth can be used to overcome these challenges.9 Cation-exchange reactions are commonly used to grow a shell on NCs, providing a way to improve their optical and electronic properties while keeping their size and shape unchanged.11-14 Depending on the reaction parameters and the miscibility of phases, heterostructures or complex homogeneous alloys can be prepared via cation exchange, which are not otherwise accessible in a single synthesis step.15 It is understood that cation-exchange reactions in NCs involve diffusion of cations and are enabled by favorable thermodynamic properties.11 At higher reaction temperatures, more incoming cations can diffuse into the NC core13,

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and smoother interfaces (i.e., increased

alloying between core and shell materials) are observed.17 These effects point to thermally activated processes in the cation-exchange growth mechanism. While diffusion is well studied for bulk materials18 and thin films,19 cation-exchange reactions for colloidal NCs remain largely underexplored.12


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Several groups have previously performed studies of cation-exchange processes.13,

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For

example, Cossairt et al. analyzed the multistage Cd2+ diffusion mechanism in ZnTe nanorods by measuring the elemental composition of aliquots during the cation-exchange reaction.21 Here, we show how the multistage nature of a cation-exchange shell formation can be tracked using in situ PL measurements. The time evolution of the emission spectra during the synthesis is analyzed as a function of growth time, temperature and injection rate. As a model system, we select indium-rich group I-III-VI NCs. With the growth of a ZnS or ZnSe shell through the cation-exchange method, PL quantum yields of Cu-In-S, Cu-In-Se, Ag-In-S, and Ag-In-Se NCs16, 22-25 can be increased from 20-30% to exceed 80%,26-27 making them of interest for use in optical and optoelectronic devices28-30 and in bio-imaging applications.31-32 For the Cu-In-Se/ZnSe1-xSx NCs we focus on here, zinc and sulfur (or selenium) precursors can be added just after synthesis of Cu-In-Se (CISe) core-only NCs33 or in a second step after purification and re-dissolving of the CISe NCs.26-27, 31, 34 The average size of NCs remains nearly constant during the cation-exchange shell growth, which means the CISe cores are partially replaced by the shell, explaining the blue-shift of PL spectrum (i.e., shift towards higher energies).15-16,

34-35

Due to the multi-cationic nature of the Cu-In-Se/ZnSe1-xSx NCs, it is

challenging to quantify the type of shell (i.e., extent of alloying and shell thickness) that has been produced or understand how shell growth conditions influence these parameters. Our in situ PL approach allows us to identify distinct steps in the multistage cation-exchange growth of the ZnSe1-xSx shell and calculate the activation energies associated with each step. We determine that the cation-exchange growth mechanism consists of a surface reaction of the Zn precursor, followed by sequential solid-state diffusion of Zn atoms into successive atomic


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monolayers. These insights enable the selection of growth condition to achieve a specific atomic composition and core/shell alloying.

Figure 1. (A) Cation-exchange reaction scheme for Cu-In-Se/ZnSe1-xSx core/shell nanocrystal growth. (B) Schematic of the in situ photoluminescence setup and an example emission spectrum showing how a Gaussian fit can be used to extract a peak position, width, and height.

Experimental Section Materials. CuCl (anhydrous, ≥ 99.99 %), InCl3 (anhydrous, 99.999 %), diethylzinc (97 %), selenium (99.99 %) and tri-n-octylphosphine (TOP, 97 %) are purchased from STREM chemicals; sulfur (99.5 %), oleic acid (techn., 90%), toluene (≥ 99.7 %), ethanol (≥ 99.8 %) and methanol (≥ 99.9 %) from Sigma Aldrich; lithium bis(trimethylsilyl)amide (LiN[Si(CH3)3]2, 95%) and squalane (99 %) from Acros. General Remarks. All syntheses are carried out in an air-free environment using standard Schlenk line technique. Injection mixtures and stock solutions are prepared in a N2-filled glove


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box. A stock solution of 1 M sulfur in TOP (i.e., TOP:S) is prepared by dissolving respective amounts of sulfur in TOP at room temperature. Caution: diethylzinc and its solutions should be handled with care due to high toxicity and flammability of the substance. Scaled-up Synthesis of Cu-In-Se Core-Only Nanocrystals. To enable comparison of the different shell syntheses, we prepare a single batch of CISe NC cores. 6 g of Cu-In-Se NCs with a stoichiometry of Cu3In5Se9 and an average size of 3.3 ± 0.4 nm are fabricated using the upscaling technique described previously.36 In short, a solution containing 20 mmol of CuCl and 20 mmol of InCl3 in 120 mL of TOP, is heated to 305°C in a 1 L three-neck flask. An injection mixture, containing 80 mmol of Se and 120 mmol of LiN[Si(CH3)3]2 in 120 mL of TOP, is swiftly added from the connected funnel using underpressure-governed hot-injection technique.36 The temperature drops to 210-220°C and the reaction is terminated after 2 min of growth time by cooling the reactor to room temperature with pressurized air and later (at temperatures < 170°C) with a cold water bath. The obtained CISe NCs are purified three times with toluene, methanol, and ethanol as described in Ref. [36]. Finally, CISe NCs are dissolved in toluene with a concentration of 50 mg·mL–1 and stored in a glove box prior to synthesis of CISe-based core/shell NCs. The optical and structural properties of the NCs are shown in Figure S1. Synthesis of Cu-In-Se/ZnSe1-xSx Core/Shell Nanocrystals. The synthesis of Cu-In-Se/ZnSe1xSx

core/shell NCs is based on a previously reported cation-exchange recipe, whereby a Zn and S

precursor mixture is added to the solution of CISe NCs at elevated temperatures.34, 37 This results in substitution of Cu and In cations with Zn cations, as shown schematically in Figure 1(A). Briefly, a reaction mixture containing 1 mL of CISe NCs solution in toluene (50 mg of NCs in total), 10 mL of TOP and 2 mL of 1 M TOP:S is added to a N2-filled three-neck flask. The toluene is carefully removed by applying vacuum at room temperature for 5 minutes. The


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reaction flask is then backfilled with N2 and a syringe containing the Zn-precursor is connected to it. To investigate the effect of the total injection time, the syringe is filled either with 6 mmol of diethylzinc, 30 mL of TOP, and 6 mL of 1 M TOP:S or with 2.5 mmol of diethylzinc, 12.5 mL of TOP, and 2.5 mL of 1 M TOP:S. The reaction mixture is placed into a pre-heated squalane oil bath, enabling rapid heating to set temperatures ranging from 70°C to 174°C and temperature control within ± 2-3°C. The injection of Zn and S precursors is started as soon as the desired temperature is reached, and a syringe pump is used to maintain an injection speed of 0.7 mL·min–1 (approx. 0.115 mmol Zn·min–1). For the large injection volume (6 mmol of diethylzinc in 30 mL of TOP), this means a total injection time of 52 minutes. For the small injection volume (2.5 mmol of diethylzinc in 12.5 mL of TOP), the total injection time is 21 minutes. After the injection is complete, the reaction flask is cooled with a water bath and the solution of NCs is purified according to previous recipes.16, 34 Hereafter, the syntheses performed with 6 mmol or 2.5 mmol zinc precursor for temperatures between 70°C to 174°C and an injection speed 0.7 mL·min–1 will be referred to as the “long injection time temperature series” or “short injection time temperature series”, respectively At 120°C, syntheses with 6 mmol of diethylzinc in 30 mL of TOP and 6 mL of 1 M TOP:S in the syringe are additionally performed for injection speeds of 0.2 mL·min–1 and 2 mL·min–1. We refer to these as the “slow injection” and “fast injection” syntheses. In situ PL Measurements. During the shell growth, the reaction mixture is excited with a blue laser (CW 405 nm) through the flask side wall (Figure 1(B)). The squalane oil does not absorb at this wavelength. A fiber collects the luminescence and the PL spectra are recorded every second with an OceanOptics QE65000 spectrometer. The recorded PL spectra are broad, due to the two


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emission channels present in these I-III-VI nanocrystals.23 For simplicity, we fit the emission with a single Gaussian and extract the peak position, width, and height of all PL spectra. A full description of the data processing is given in Supporting Information. In particular, the raw in situ PL data is corrected for the temperature dependence of the NC emission. For temperature normalization, we perform reference heating measurements of core-only and core/shell NCs and observe a nearly linear dependence of the PL peak position on temperature (Figure S2(A)), consistent with previous reports.30, 38 We use this linear dependence (Table S1) to extrapolate the PL data to room temperature (Figure S2(B)). Figure 2 shows data from a representative in situ PL measurement. The temperature is maintained at 121.6 ± 1.5°C with an oil bath (Figure 2(A)). Both the PL peak energy and the PL width change throughout the shell growth (Figures 2(B,C)). For the PL peak position, a red-shift is visible during the first minute, followed by a strong blue-shift at longer reaction times. The PL peak width increases during the first 10 minutes of cation-exchange reaction, slowly decreasing afterwards until the end of the process. A time evolution schematic depiction of the emission spectrum is shown in Figure 2(D). We analyze these dependences to develop a model for shell growth via cation exchange. Characterization of Cu-In-Se/ZnSe1-xSx Core/Shell Nanocrystals. To complement the in situ PL measurements, all obtained NCs are characterized optically, structurally, and chemically. The absorption

measurements

are performed

with

an

Agilent

Cary 5000

UV-vis-NIR

spectrophotometer and the photoluminescence spectra of purified NCs in toluene (i.e., ex situ PL measurements) are taken with an OceanOptics QE65000 spectrometer using a CW 532 nm laser excitation. Energy-dispersive X-ray spectroscopy (EDX) data are measured with FEI Quanta 200 FEG and Hitachi S4800 SEM microscopes, operating at 30 keV. Transmission electron


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microscopy is performed on FEI Tecnai F30 FEG TEM microscope, operating at 300 keV. Size distribution analysis is carried out using ImageJ software, measuring sizes of > 100 NCs. X-ray diffraction measurements are made on a Rigaku SmartLab 9kW System with rotating Cu anode and 2D solid state detector HyPix-3000SL.

Figure 2. Results from an example in-situ PL measurement during shell growth with a stable temperature of 121.6 ± 1.5°C (A): (B) evolution of the emission peak position (Epeak), and (C) dependence of the emission peak full-width-at-half-maximum (FWHM) during Zn chalcogenide shell growth. (D) Schematic representation of the evolution of the PL peak position and width. The bold black, red, green, and blue spectra correspond to the reaction times marked with circles in (A-C).

Results and Discussion To begin our investigation of the cation-exchange mechanism, we consider the long injection time temperature series (Figure S3). Figure 3(A) shows temperature-normalized PL peak positions for each growth temperature as a function of reaction time. The 2D map (Figure 3(B))


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plots the shift of PL peak energy during the cation-exchange process relative to that of the coreonly CISe NCs. For temperatures below 150°C, an initial red-shift of the PL peak is visible, followed by a shift in peak positions to higher energies. At higher energies only a blue-shift of the PL peak is observed.

Figure 3. (A) PL peak position as a function of reaction time for the long injection time temperature series. (B) PL peak shifts from core-only CISe NCs mapped versus shell growth temperature and time. (C) Zoomed region of PL red-shifts and exponential decay fits for temperatures below 130°C. Inset (C): schematic illustration of Zn surface attachment, resulting in a slight red-shift of PL peak energy. (D) Arrhenius plot of the PL red-shift rate and corresponding activation energy for the surface reaction of Zn precursor.


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Surface Reaction and Subsequent Diffusion. We propose that this initial red-shift of the PL spectrum is due to an increase of the NC size, coming from the attached of Zn precursor molecules to the NC surface (inset of Figure 3(C)). Indeed, small red-shifts (tens of meV) are typically observed when shell material grows epitaxially around the core NCs.7 For all growth temperatures, the red-shift of PL peak position is 41 meV. This can be explained by the fact that our starting core NCs are the same for all shell growth syntheses and that the process is limited by the number of atomic sites that can be populated by Zn. The rate of the PL red-shift with time, which we fit with an exponential decay, exhibits a notable temperature dependence (Figure 3(C)). An Arrhenius-type fit of the PL peak red-shift rates vs. 1/T (Figure 3(D)) indicates an activation energy of 0.43 eV for the surface reaction of Zn precursor during the cation-exchange process (full description of fitting methods in Supporting Information). Furthermore, from the slow and fast injection syntheses at 120°C we observe that the red-shift of the PL peak energy occurs at similar rates for both normal and fast injection speeds (Figure S4). We conclude that, for the chosen reaction parameters of the temperature series, the surface reaction of Zn precursor is limited by surface reaction kinetics and not the availability of the precursor in the solution. For temperatures above 150°C the red-shift associated with the surface reaction is already from the beginning of the reaction outweighed by the strong shift to higher energies and therefore not visible in the PL spectra. This blue-shift of the PL emission spectrum during shell growth indicates an increased electronic confinement of the core NCs, which we associate with solid-state diffusion of Zn into the NC, replacing Cu and In ions. To analyze the PL peak position blue-shifts, we subtract the red-shift contribution attributed to the Zn surface reaction (Figure 4(A), calculation details in Supporting Information). The blue-shift is larger for higher growth temperatures, indicating more efficient


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diffusion processes with increasing temperatures. For high temperatures, curves of the blue-shift as a function of reaction time are best fit with a double exponential, whereas low temperature curves show single-exponential behavior (Figure 4(A)). This suggests two different diffusion processes. Arrhenius plots of the two sets of exponents reveal the activation energies of a lowand a high-energy process as 0.18 eV and 0.51 eV, respectively (Figure 4(B)). Their corresponding pre-exponential factors describe the contribution of each process at the different shell growth temperatures (Figure 4(C)).

Figure 4. (A) Relative blue-shift of PL peak position for the long injection time temperature series after subtraction of the red-shift contribution and single/double exponential growth fits. (B,C) PL shift rates and pre-exponential factors as derived from fits in (A). (D) Cationic composition of Cu-In-Se/ZnSe1-xSx core/shell nanocrystals for the short injection time temperature series. (E) Calculated number of zinc and sulfur atoms in a single nanocrystal. (F) Illustration of cation-exchange process, including surface attachment of Zn atoms and temperature-dependent solid-state Zn diffusion through subsequent atomic monolayers.


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Elucidating the Diffusion Mechanism. To elucidate the origins of the two distinct activation energies associated with Zn incorporation into the NCs, we determine the size and composition for all NCs from the two temperature series. Transmission electron microscopy (TEM) and X-ray diffraction data indicate that the size of Cu-In-Se/ZnSe1-xSx core/shell NCs is approximately equivalent to that of core-only CISe NCs (Figures S5 and S6) and that the crystal structure remains a zinc blende type34 before and after the growth of Zn chalcogenide shell (Figures S7 and S8). Figures 4(D) and S9(A) show the cationic compositions, calculated by energydispersive X-ray spectroscopy (EDX), as a function of reaction temperature for short and long injection times, respectively. While the content of Zn increases with higher temperatures, the indium and copper content decreases for both series. Sulfur content remains consistently smaller than that of Zn, indicating that the shell has a ZnSe1-xSx gradient composition. Combining EDX results with known densities of materials and size of NCs (3.15 nm), we estimate the number of Zn and S atoms per NC using the mathematical description given in Supporting Information (Figures 4(E) and S9(B)). This information enables us to rule out possible explanations for the two diffusion processes. Previous literature on I-III-VI thin films suggests several mechanisms for solid-state zinc diffusion.19 The two diffusion processes observed (Figures 4(A-C)) could be associated with e.g., different Zn diffusion mechanisms (interstitial vs. substitutional diffusion), differences between Zn-to-Cu and Zn-to-In substitution (i.e., ZnCu and ZnIn, respectively), or by competing cation- and anion-exchange processes. We rule out all of these explanations, considering that the high-energy process is insignificant below 100°C (Figure 4(C)). Since a significant amount of Zn cations are present in the structure at low temperatures (> 25% of cations, Figure 4(D)), the


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diffusion cannot be attributed to being solely interstitial. This indicates that substitutional diffusion is prevalent in all cases, so assigning the low-energy process to interstitial site diffusion and high-energy process to substitution is therefore not physical. Due to the larger ionic size of In3+, we assume that the ZnIn substitution process is more energy intense than ZnCu.39 However, even for low temperatures, a considerable amount of indium is replaced with zinc so linking the two processes to the zinc substitution mechanisms is not coherent with our observations. Finally, negligible amounts of sulfur below 150°C (Figure 4(E)) rules out anion-exchange as the cause of the blue-shift. Furthermore, comparing ex-situ PL emission energies with elemental composition shows a clear dependence on zinc content, while no trend is visible for sulfur content (Figure S10). We therefore need a different explanation for the two activation energies associated with Zn incorporation. In this study, CISe core NCs have a radius of less than three unit cells, which is significantly smaller than typical diffusion reaction zones. Indeed, diffusion is known to operate differently for nano-sized systems compared to bulk materials.40 For example, atomically sharp interfaces have been predicted for 2D thin films.41 In surface diffusion processes, each atomic monolayer may represent its own energetic barrier.42 We therefore propose that the low-energy process corresponds to the filling of the first atomic monolayer, while the high energy process corresponding to filling of the second monolayer. Assuming a subsequent filling of the atomic monolayers, the first monolayer is filled with Zn atoms for all but the lowest temperatures (Figures 4(E) and S9(B)). This would explain the modest blue-shift for all temperatures below 100°C. Filling of the second monolayer (i.e., the high-energy process) is only activated above 100°C, and the extent it is filled increases linearly with temperature, according to the behavior of the pre-exponential factors.


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As soon as the second layer becomes accessible, the filling level of the first slightly decreases, which can be seen as a decrease of the first pre-exponential factor between 100°C and 150°C (Figure 4(C)). We attribute this observation to finite solid-state diffusion kinetics (i.e., a move of Zn ions between first and second atomic monolayer). When approaching the maximum filling of the second monolayer, also the first is filled again to a higher degree. The dotted lines in Figure 4(E) mark the number of cations or anions present in the first outermost or two outermost atomic monolayers of NC (calculation details in Supporting Information). Intersections with these lines and the number of zinc correspond to the two temperatures 100°C and 150°C mentioned above, thus corroborating our proposed explanation of the stepwise diffusion. The evolution of PL peak width (full-width-at-half-maximum, FWHM) for the different reactions gives further evidence of a stepwise diffusion mechanism (Figure 5(A)). For temperatures below 100°C, the PL emission peak continues to broaden until the end of the reaction, whereas, for temperatures higher than 100°C, the PL peak exhibits a maximal broadening followed by gradual decrease of the PL peak width. The higher the growth temperatures, the shorter is the time at which the largest FWHM occurs. We attribute the PL peak broadening to the strain of the core/shell structure, originating from the lattice mismatch between core and shell materials.34 Inserting more Zn atoms increases strain due to an abrupt concentration gradient. Higher growth temperatures promote a relaxation of the structure via Zn diffusion, decreasing the lattice strain. Following this hypothesis, structural strain reaches a maximum when the first monolayer is completed with Zn atoms, while inner layers remain unchanged (which is the case at approx. 100°C, Figure 4(C)). When the filling of


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the second layer is kinetically accessible (above 100°C), the lattice strain can partially relax (Figure 5(B)).

Figure 5. (A) Full-width-at-half-maximum (FWHM) of the emission peak, mapped versus growth temperature and time for the long injection time temperature series. (B) Schematic explanation of PL width broadening through variable lattice strain due to the Zn diffusion through the atomic monolayers (MLs).

Conclusion In summary, we used in situ PL measurements to study the cation-exchange process during the growth of Cu-In-Se/ZnSe1-xSx core/shell NCs. At the beginning of the ZnSe1-xSx shell synthesis, a red-shift of emission spectra is observed, implying a size increase due to the surface attachment of zinc precursor molecules. The blue-shift of PL peak position can be associated with solid-state Zn diffusion through subsequent atomic monolayers of CISe NCs. While the surface monolayer of NCs can be exchanged with Zn ions at all shell growth temperatures between 70 and 174°C, the second outermost atomic CISe monolayer only becomes thermally accessible at elevated temperatures above 100°C. This multistage process is illustrated schematically in Figure 4(F).


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All stages of the cation-exchange shell formation are thermally activated processes. The activation energy of the surface Zn reaction is 0.43 eV, while the activation energies of Zn diffusion into the first and second outermost atomic monolayers are 0.18 eV and 0.51 eV, respectively. The initial surface process exhibits a higher energetic barrier compared to cationexchange in the surface atomic monolayer. This suggests that the surface reaction of Zn precursor is the limiting stage of the cation-exchange process and, once attached, Zn atoms easily diffuse into the outermost monolayer of the CISe NCs. Accordingly, the thickness of the ZnSe1xSx

shell is limited by temperature more than by reaction time or added zinc (Figure S10).

For luminescent nanocrystals, in situ PL measurements during cation-exchange shell growth enables a fast and detailed analysis of nanoscale diffusion without disturbing the system during shell growth synthesis. The influence of single reaction parameters on distinct diffusion processes can be accessed and activation energies calculated. A better understanding of cationexchange process kinetics at the atomic level can enable better engineering of core/shell nanocrystals for luminescence applications, including bio-labeling, optical down conversion, and electroluminescent devices.

ASSOCIATED CONTENT Supporting Information. Details of fitting methods and calculations, supporting Figures S1-S10 Author Contributions The manuscript was written through contributions of all authors. Funding Sources


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The authors acknowledge support from the Swiss National Science Foundation through the Quantum Sciences and Technology NCCR and an ETH Research Grant. ACKNOWLEDGMENT Authors thank David Norris for the SEM microscopy access and Mario Mücklich for technical lab assistance. TEM and EDX measurements are performed at the Scientific Center for Optical and Electron Microscopy (ScopeM) of the Swiss Federal Institute of Technology, Zurich. References 1. Dabbousi, B. O.; Rodriguez-Viejo, J.; Mikulec, F. V.; Heine, J. R.; Mattoussi, H.; Ober, R.; Jensen, K. F.; Bawendi, M. G. (CdSe)ZnS Core-Shell Quantum Dots: Synthesis and Characterization of a Size Series of Highly Luminescent Nanocrystallites. J. Phys. Chem. B 1997, 101, 9463-9475. 2. Reiss, P.; Protière, M.; Li, L. Core/Shell Semiconductor Nanocrystals. Small 2009, 5, 154-168. 3. Chaudhuri, R. G.; Paria, S. Core/Shell Nanoparticles: Classes, Properties, Synthesis Mechanisms, Characterization, and Applications. Chem. Rev. 2012, 112, 2373-2433. 4. Talapin, D. V.; Mekis, I.; Götzinger, S.; Kornowski, A.; Benson, O.; Weller, H. CdSe/CdS/ZnS and CdSe/ZnSe/ZnS Core-Shell-Shell Nanocrystals. J. Phys. Chem. B 2004, 108, 18826-18831. 5. Xie, R.; Kolb, U.; Li, J.; Basché, T.; Mews, A. Synthesis and Characterization of Highly Luminescent CdSe-Core CdS/Zn0.5Cd0.5S/ZnS Multishell Nanocrystals. J. Am. Chem. Soc. 2005, 127, 7480-7488. 6. Chen, Y.; Vela, J.; Htoon, H.; Casson, J. L.; Werder, D. J.; Bussian, D. A.; Klimov, V. I.; Hollingsworth, J. A. “Giant” Multishell CdSe Nanocrystal Quantum Dots with Suppressed Blinking. J. Am. Chem. Soc. 2008, 130, 5026-5027. 7. Mahler, B.; Spinicelli, P.; Buil, S.; Quelin, X.; Hermier, J.-P.; Dubertret, B. Towards Non-Blinking Colloidal Quantum Dots. Nat. Mater. 2008, 7, 659-664. 8. Cragg, G. E.; Efros, A. L. Suppression of Auger Processes in Confined Structures. Nano Lett. 2010, 10, 313-317. 9. Bozyigit, D.; Yarema, O.; Wood, V. Origins of Low Quantum Efficiencies in Quantum Dot LEDs. Adv. Funct. Mater. 2013, 23, 3024-3029. 10. Kagan, C. R.; Lifshitz, E.; Sargent, E. H.; Talapin, D. V. Building Devices from Colloidal Quantum Dots. Science 2016, 353, 885-894. 11. Beberwyck, B. J.; Surendranath, Y.; Alivisatos, A. P. Cation Exchange: A Versatile Tool for Nanomaterials Synthesis. J. Phys. Chem. C 2013, 117, 19759-19770. 12. De Trizio, L.; Manna, L. Forging Colloidal Nanostructures via Cation Exchange Reactions. Chem. Rev. 2016, 116, 10852-10887.


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In Situ Monitoring of Cation-Exchange Reaction ... - ACS Publications

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