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HgSe/CdE (E=S,Se) Core/Shell Nanocrystals by Colloidal Atomic Layer Deposition Laxmi Kishore Sagar, Willem Walravens, Jorick Maes, Pieter Geiregat, and Zeger Hens J. Phys. Chem. C, Just Accepted Manuscript • Publication Date (Web): 05 Jun 2017 Downloaded from http://pubs.acs.org on June 7, 2017

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The Journal of Physical Chemistry

HgSe/CdE (E=S,Se) Core/Shell Nanocrystals by Colloidal Atomic Layer Deposition Laxmi Kishore Sagar,†,‡ Willem Walravens,†,‡ Jorick Maes,†,‡ Pieter Geiregat,†,‡ and Zeger Hens∗,†,‡ Physics and Chemistry of Nanostructures, Ghent University, Ghent, Belgium, and Center for Nano and Biophotonics, Ghent University, Belgium E-mail: [email protected]

Abstract Core/shell colloidal quantum dots (QDs) can have a more efficient photoluminescence than core only QDs due to better surface passivation and a localization of the charge carriers in the core. Here, we demonstrate the synthesis of HgSe/CdS, HgSe/CdSe and HgSe/CdSe/CdS core/shell QDs at room temperature by a succession of self-limiting half reactions. A detailed structural characterization through transmission electron microscopy analysis shows that shell growth happens in an additive manner. Each sequence of two half reactions results in a ≈ 0.5 nm increase in shell thickness, meaning that approximately one monolayer is grown at the time. In all cases, shell growth eliminates the n-type doping that is intrinsic to HgSe QDs. In the case of HgSe/CdSe core/shell QDs, shell growth is strain-free as both materials have an almost identical lattice parameter. CdS shells on the other hand lead to the most ∗

To whom correspondence should be addressed Physics and Chemistry of Nanostructures, Ghent University, Ghent, Belgium ‡ Center for Nano and Biophotonics, Ghent University, Belgium †

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pronounced increase of the photoluminescence quantum yield, reaching 16.5% after mild annealing for HgSe/CdS QDs emitting at around 0.69 eV (1.8 µm).

Colloidal semiconductor nanocrystals or quantum dots (QDs) offer size-tunable optoelectronic properties, where a desired wavelength range can be addressed by the appropriate choice of material and nanocrystal dimensions. 1–3 For applications at short-wave infrared wavelengths, for example, differently sized lead chalcogenide nanocrystals give access to nanomaterials with a band gap ranging from 800 to 2500 nm. 4–6 Driven by highly precise synthesis methods and the often unique properties of these nanoscale materials, extensive research was conducted into the use of PbS and PbSe QDs for photovoltaic energy conversion, 7–10 photodetection, 11,12 and infrared imaging. 13 However, due to their long radiative lifetime and eightfold degenerate band-edge states, PbS and PbSe QDs are less interesting for applications involving the emission or amplification of infrared light. 14–16 A less explored alternative for the Pb-chalcogenides are nanocrystals of the mercury chalcogenides, such as HgSe and HgTe. As bulk materials, these are semi-metals, yet when the crystal dimensions are reduced to nanometer sizes, a band gap opens up that can be tuned through the entire infrared range. 17 While mainly studied in view of photodetection, the lower degeneracy of their band-edge states and their shorter radiative lifetime make them appealing alternatives to PbS and PbSe for applications involving light emission and amplification. 18 Over the last 20 years, the formation of core/shell QDs has proven to be a key enabler for light emitting applications based on QDs. Inorganic shells can provide excellent surface passivation of the core QD and allow for a further tuning of the opto-electronic properties. 19,20 While a broad range of shell growth procedures has been developed for Zn, Cd and Pb-chalcogenide QDs, literature shows little examples of Hg-based core/shell QDs. A notable example involves the formation of CdS/HgS/CdS quantum dot/quantum wells and CdTe/HgTe inverted core/shells by the exchange of surface adsorbed Cd2+ for Hg2+ in aqueous media and reaction with hydrogen sulfide or telluride gas, respectively. 21–23 A similar approach was pursued more recently to synthesize HgTe/CdS core/shell QDs, 24 while a more 2 ACS Paragon Plus Environment

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Table 1: Lattice mismatch (%) of possible II-VI semiconductors with HgSe at 27 ◦ C. Positive sign (+) indicates tensile strain, negative sign (-) indicates compressive strain. For all materials, data pertain to the zinc blende polymorph. Shell Lattice mismatch HgSe (%)

ZnS + 11.01

ZnSe + 6.74

ZnTe - 0.32

CdS + 4.27

CdSe 0

CdTe - 6.57

involved aqueous/non-aqueous approach has been followed to form (Cd,Hg)Te/ZnS QDs. 25 The latter studies, however, focused on optimizing photoluminescence quantum yield starting from a single set of small core nanocrystals such that the thus formed core/shell QDs fail to cover a broad spectral range in accordance to HgSe and HgTe core QDs. 26–29 Taking the seminal case of CdSe quantum dots, established shelling methods involve either a layer-by-layer approach through successive exposure to cation and anion precursors or the use of core quantum dots as seeds for the heterogeneous nucleation and growth of the shell material. 30,31 Such shelling reactions typically use elevated temperatures – up to 300 ◦ C or more – and work best when the lattice mismatch between core and shell is minimal. In the case of CdSe-based core/shell QDs, for example, the relatively small 4.3% lattice mismatch between CdSe and CdS translates into epitaxial shells that can be several nanometers thick, whereas ZnS shells (lattice mismatch of ≈ 11.0%) remain restricted to one or two monolayers at best. Using lattice mismatch as an initial criterion to select appropriate shell materials for HgSe core QDs, Table 1 shows that the zinc blende polymorphs of CdSe, ZnTe and CdS stand out since their lattice parameters differ by less than 5% from that of HgSe. As ZnTe is prone to oxidation and often leads to core/shell QDs with a staggered band alignment, we consider CdSe and CdS to be the most interesting materials for shell growth around HgSe QDs. However, the tendency of mercury and cadmium chalcogenides to form solid solutions rather than phase separated core/shell structures can compromise the formation of such HgSe/CdE (E=S,Se) core/shell QDs by a mere duplication of established shelling reactions. 32–35 This may in part explain why all the above discussed procedures involve room temperature or close-to room temperature reactions. 3 ACS Paragon Plus Environment


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Here, we demonstrate the formation of HgSe/CdSe and HgSe/CdS core/shell QDs using a room temperature colloidal layer-by-layer process, similar to the so-called colloidal atomic layer deposition (c-ALD) as originally developed by Ithurria et al. 36 and recently applied for the formation of PbS/CdS core/shell QDs in an additive layer-by-layer fashion. 37 In particular, we show that successive exposure of HgSe QDs to Cd and S or Se precursors leads to a progressive increase of the nanocrystal diameter by ≈ 0.50 nm/cycle; an observation confirming that a single cycle of two half reactions results in the formation of one CdE monolayer. The formation of HgSe-based core/shell QDs with controlled shell thickness enables us to systematically study the influence of shell growth on the optical and structural properties of the resulting core/shell QDs. X-ray diffraction (XRD) analysis points to a rather unique strainless core/shell system in the case of HgSe/CdSe in line with their nearly equal lattice parameters. Remarkably, both CdSe and CdS shell growth remove the intrinsic n-type doping of larger HgSe QDs and improve the photoluminescence quantum yield by almost 2 orders of magnitude as compared to the original core QDs. In this way, this work provides a crucial next step in the understanding of HgSe-based QDs and their application for infrared optoelectronics.

Experimental Section Chemicals. Mercury chloride (HgCl2 , 99.5%), selenourea (98%), dodecane thiol (DDT, 98%), tetracholoroethylene (TCE, 99.0%, ACS reagent), formamide (99.5%, ACS reagent), ammonium sulfide (20% solution in water) and cadmium acetate dihydrate (CdAc2 .2H2 O, 98%) were obtained from Sigma-Aldrich. Sodium selenide (Na2 Se, 99.8%) was obtained from Alfa-Aesar. Oleylamine (OLA, 80-90%) was obtained from Acros Organics. All the solvents like toluene, methanol and isopropanol were purchased from VWR. HgSe synthesis. HgSe QDs were prepared by following the procedure of Deng et al. 26 Typically, 54 mg (0.2 mmol) of HgCl2 was dissolved in 8 mL of oleylamine. The reaction mixture 4 ACS Paragon Plus Environment

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was flushed under nitrogen atmosphere and heated to 110 ◦ C for an hour. Afterwards, the temperature was lowered and 0.2 mmol of selenourea dissolved in 2 mL of oleylamine was injected to start the HgSe formation. To obtain 3 nm QDs, a growth temperature of 60 ◦ C and a growth time of 2 minutes was used, to obtain 4.25, 4.7 and 5 nm QDs, a growth temperature of 80 ◦ C and growth times of 6, 10 and 15 minutes were used, respectively, and to obtain 5.7 and 6 nm QDs, a growth temperature of 110 ◦ C and growth times of 3 and 10 minutes were used, respectively. In all cases, the reaction was quenched by simultaneous injection of 3 mL of toluene and 1.2 mL of DDT and cooling with a water bath. The QDs were purified 3 times with toluene/methanol as the solvent/anti-solvent combination and stored in toluene.

HgSe/CdS synthesis. We synthesized HgSe/CdS core/shell QDs using a colloidal atomic layer deposition (c-ALD) approach as initially developed by Ithurria et al. 36 Typically, a solution of 3 mg of HgSe cores in 1 mL of toluene and 100 µL of oleylamine was combined with 1.5 mL of formamide to form a liquid-liquid biphasic system. To grow a single layer of CdS, 40 µL of ammonium sulfide (20% ammonium sulfide solution in formamide, 5 mg, 0.073 mmol) was first added to this system, after which the biphasic solution was thoroughly shaken, the polar phase was discarded and the non-polar phase was washed twice with formamide to extract residual ammonium sulfide. Next, 1.5 mL of formamide was added to make a biphasic system once more, together with 60 µL of formamide solution containing 1.6 mg (0.0061 mmol) of cadmium acetate. This solution was again thoroughly shaken, the polar phase was discarded and the non-polar phase was washed twice to extract residual cadmium acetate. By repeating this S and Cd exposure, different shells can be grown one after the other. While we typically grew up to 5 shells, nothing restricts a further continuation of the procedure to grow any number of shells desired. Before recording the absorption spectra, the samples were purified twice using toluene as the solvent and a methanol/isopropanol mixture as the non-solvent to get rid of excess OLA and other organic impurities.



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HgSe/CdSe synthesis. We synthesized HgSe/CdSe core/shell quantum dots in a similar fashion as HgSe/CdS core/shell quantum dots, except that 20% sodium selenide in formamide was used as the Se precursor.

Annealing of the core/shell quantum dots. Annealing was achieved by adding 300 µL of dodecanethiol (DDT) to a fraction of HgSe/CdS or HgSe/CdSe QDs in toluene. The resulting mixture was heated at 60 ◦ C or 100 ◦ C on a hot plate, with a continuous stirring for 45 minutes. Afterwards, the sample was cooled to room temperature and purified with toluene/methanol to remove excess organic impurities. All the experiments were carried out under inert atmosphere.

Structural Characterization. XRD samples were prepared by drop casting a layer of the desired core or core/shell material from a hexane:heptane (80:20) solution on a glass substrate. Measurements were performed on a ARL XTRA Powder Diffractrometer. Transmission electron microscopy (TEM) images were recorded on an aberration corrected JEOL 2200-FS operated at 200 kV. High angle annular dark field (HAADF) images were recorded in Scanning TEM mode with a spot size of 0.7 nm and a camera length of 60 cm. For STEM-HAADF analysis, 4.7 nm HgSe core QDs were overcoated by either 3 CdSe shells or 3 CdS shells.

Optical Characterization. For UV-Vis absorption spectroscopy, known amounts were taken from reaction aliquots or stock solutions of HgSe core and HgSe/CdE core/shell QDs dispersed in toluene and kept in a nitrogen-filled glovebox, dried under nitrogen atmosphere, redispersed in tetracholoroethylene (TCE) and analyzed using a Perkin-Elmer Lambda 950 UV-Vis-NIR spectrophotometer. Absorbance at wavelengths beyond 3 µm was analysed using a Nicolet 6700 Fourier Transform Infrared Spectrometer. Photoluminescence (PL) was measured using a FL920 Edinburgh spectrofluorometer with a 450 W Xenon lamp for excitation. Depending on the emission wavelength, the photoluminescence is measured using a


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Figure 1: (a) Absorption spectra of different sized HgSe QD dispersions. The inset shows the level diagram used to interpret the spectra, where (full lines) H-S and (dashed lines) H-P are used to indicate the first and second interband transition and (dotted lines) S-P the intraband transition. The spectra make clear that with increasing size, the band-edge transition H-S collapses, the intraband transition S-P appears and the interband transition H-P persists, an evolution attributed to the accumulation of electrons in the lowest conduction band level S. All spectra have been normalized at 3.1 eV (outside the range of energy shown in the figure) and offset vertically for clarity. (b-e). Bright field TEM images of the different QD samples used in (a). In each micrograph, the dimension mentioned refers to the estimated average QD diameter as obtained from analysing TEM images; the error on the estimated diameter is in all cases smaller than 0.05 nm. liquid N2 cooled InGaAs PMT or InSb detector. The emission spectra of core and core/shell QDs was measured by continuous wave excitation at 0.8 µm. Quantum yield (QY) measurements with the InSb detector were performed using a 5 nm PbS QD sample as a reference. The photoluminescence quantum yield of the reference was determined using an integrating sphere an the InGaAs PMT.

Results and Discussion HgSe Quantum Dots (QDs) Figure 1a shows the absorption spectra, normalized at 3.1 eV, of differently sized HgSe QDs as synthesized according to the procedure of Deng et al. (see Experimental Section). The formation of HgSe QDs is attested by the TEM images shown in Figures 1b-e. 26 The



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spectrum of the 3 nm HgSe QDs features the typical band-edge transition, labeled here as H-S and indicated with full lines, at λH−S ≈ 1.02 eV (1.21 µm) together with an additional feature, indicated with a dashed line, at around 1.5 eV. Turning to the 4.3 nm HgSe QD spectra, it can be seen that both features exhibit the redshift that comes with a reduction of the quantization energy, appearing in this case at around 0.85 eV (1.45 µm) and 1.15 eV (1.08 µm). However, the H-S transition seems to have lost intensity and a pronounced absorption band, indicated with a dotted line, at around 0.35 eV (3.5 µm) shows up. These observations are well known for HgSe QDs and have been attributed to the filling of the lowest conduction band level, which we will label as S, with electrons when the QD size exceeds a critical value. 26 This interpretation accounts for the bleach of the band-edge transition and makes that the IR absorption band can be assigned to an intraband transition between the lowest conduction-band level S and the next conduction-band state, which we will label as P. Since the energy of this S-P transition is about equal to the energy difference between the low energy (H-S) and the high energy interband feature, we assign the latter to the H-P transition. These assignments have been summarized in the level scheme added to Figures 1a. Increasing the size of the HgSe QDs further to 5.0 and 5.7 nm results in the complete disappearance of the H-S transition, a further redshift of the H-P transition and an increased absorbance due to the intraband S-P transition. Interestingly, the electron occupation in the conduction-band state S can be estimated from the ratio between the band-edge and the intraband absorption. In this case of thiol ligands, we find that electronfilling of the lowest conduction-band level starts when the QD band-gap is at around 0.95 eV (1.3 µm) and is complete for HgSe QDs with band gaps below 0.7 eV (1.7 µm), see Supplementary Information Figure S1.

Layer-by-Layer Shell growth With colloidal quantum dots (CQDs), additive shell growth is typically achieved by high temperature methods involving either seeded growth or the successive introduction of anion 8 ACS Paragon Plus Environment

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Figure 2: (a) Absorption spectra of (4.3 nm size) HgSe QD dispersions (black) prior to and (colored) after successive colloidal LBL cycles with CdSe shelling. The spectra have been normalized at 1.90 eV. (b) Absorption spectra of HgSe (4.3 nm) QD dispersions (black) prior to and (colored) after successive colloidal LBL cycles with CdS shelling. The spectra have been normalized at 2.47 eV. and cation precursors. 30,38 Both methods proved very effective to synthesize, for example, brightly emitting CdSe/CdS, CdSe/ZnS and CdTe/CdS core/shell QDs with quantum yields reportedly reaching 90%. 39–42 Especially in the case of PbS and PbSe QDs, superseded shell growth through cation exchange is typically used as an alternative to grow PbS/CdS or PbSe/CdSe core/shell QDs by exposing the original core QDs to a reaction mixture containing cadmium carboxylates at temperatures of 100 ◦ C or more. 43,44 Since the HgSe QDs are synthesized at relatively low temperatures (50-110 ◦ C), the more conventional, high temperature shell growth procedures may be problematic. When we explore either of the aforementioned shelling procedures to grow a CdE shell on HgSe QDs, we indeed always observed a marked deterioration of the size dispersion due to undesired Ostwald ripening. Similar to the case of PbS/CdS, 37 we therefore opted for a room temperature colloidal atomic layer deposition (c-ALD) process to grow CdE shells around HgSe core QDs. This approach, originally developed by Ithurria et al., 36 relies on the successive exposure of QDs to cadmium and chalcogen precursors to grow CdS or CdSe by self-limiting half cycle reactions. Figure 2a and b show the absorption spectra of a dispersion of 4.3 nm HgSe QDs before and after up to three successive CdSe or CdS c-ALD cycles, respectively. Absorption 9 ACS Paragon Plus Environment


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spectra have been normalized at 1.90 and 2.47 eV for CdSe and CdS shells, respectively. At these energies the respective shells are still largely transparent, such that the absorbance can be seen as an approximate measure for the amount of HgSe. As outlined in the Experimental Section, each cycle involves the successive exposure of HgSe QDs dissolved in a toluene/oleylamine mixture to formamide solutions of cadmium acetate and either sodium selenide or ammonium sulfide. Focusing first at the short wavelength part of the normalized absorption spectra, it appears that successive c-ALD cycles induce a gradual increase of the absorbance at wavelengths slightly below either the CdSe or CdS bulk band gap. This effect is most pronounced after 3 c-ALD cycles and yields a first indication of CdSe or CdS – CdE in short – shell growth around the original core HgSe QDs. As the HgSe core synthesis yields partially n-doped HgSe QDs (see Figure 1a), the spectrum of the original core QDs features both the typical H − S band-edge transition at λH−S ≈ 0.85 eV (1.45 µm) and a pronounced mid infrared absorption feature at around 0.35 eV (3.5 µm) that results from the S − P intraband transition. This latter feature disappears after the first CdSe or CdS deposition cycle, which suggests that the core QDs become intrinsic. In line with this conclusion, the band-edge transition gains intensity while also showing a shift by 0.06 eV (0.12 µm) to lower energies. In fact, these changes already take place after the first half cycle where HgSe QDs are exposed to Se2− or S2− anions (see Supplementary Information, Figure S2), an observation that agrees with previous literature. 26 As shown in Figure 2a-b, subsequent CdSe or CdS addition cycles result in a minor further redshift of the interband transition by about 0.01 eV (0.03 µm), whereas the normalized absorbance slightly drops and the intraband transition remains absent. The undoing of the n-type doping upon shell growth manifests itself more clearly in the case of larger HgSe QDs. These feature a fully bleached band-edge absorption that is entirely recovered upon, for example, CdSe shell growth (see Supplementary Information, Figure S3). Importantly, as HgSe doping is size-dependent, the band-edge transition will be bleached more for the larger HgSe QDs in a given batch of partially-doped HgSe QDs.


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Figure 3: (a-b) Bright field TEM image of initial 6.0 nm HgSe QDs and the HgSe/CdSe QDs obtained after 3 c-ALD cycles. (c,d) (top) size histograms of (red) the original HgSe QDs and the HgSe/CdSe QDs obtained after 1,3,4 and 5 c-ALD cycles together with (bottom) the average nanocrystal diameter as a function of the colloidal ALD cycle, the vertical error bars represent the standard deviation of the QD distribution. (e) STEM-HAADF images of the 4.7 nm HgSe QDs exposed to 3 c-ALD cycles, indicating a concentric structure where a higher Z contrast core is surrounded by a lower Z contrast shell of 1.0-1.5 nm thick. (f) x-ray diffractograms of the initial HgSe QDs (red) and the HgSe/CdSe QDs (blue) after 3 colloidal ALD cycles. The vertical (red) lines indicate the diffraction pattern of bulk HgSe. This results in an artificial blueshift of the remaining band-edge absorption feature that will be undone upon shell growth. Hence, the redshift of the band-edge transition upon shelling cannot be simply interpreted as an effect of charge-carrier delocalization, which is well-known in the seminal case of CdSe/CdS core/shell QDs. In fact, the minor shifts of the band-edge transition upon further shell growth suggest that charge-carrier delocalization is limited. Interestingly, the CdSe and CdS c-ALD growth cycles can be combined to grow more complex shells around given HgSe core QDs. As an example, absorption spectra of HgSe QDs, and successively formed HgSe/2CdSe and HgSe/2CdSe/CdS core/shell QDs are shown in the Supplementary Information (Figure S4). It appears that the further addition of CdS does not markedly influence the H − S band-edge transition. On the other hand, it leads to a pronounced increase of the absorbance at shorter wavelengths, which increases with each additional CdS growth cycle. This finding points once again towards the formation of a CdS shell in a layer-by-layer manner. 11 ACS Paragon Plus Environment


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To confirm CdE shell growth around the original HgSe core QDs, we investigated the structural properties of the HgSe core QDs exposed to successive CdSe or CdS c-ALD cycles using transmission electron microscopy (TEM). Figure 3a-b show bright field TEM images obtained before and after 3 CdSe c-ALD cycles, starting from 6.0 nm HgSe QDs. The images attest the increase in QD size without the undesired formation of pure CdSe nanocrystals by secondary nucleation. Importantly, the shape of the resultant HgSe/CdSe QDs is retained after shell growth with a very good size dispersion (≈ 5 %). Size histograms analyzed from TEM images of the original HgSe and the HgSe/CdSe QDs after each of the 5 c-ALD cycles are given in Figure 3c. These clearly show a gradual increase of the QD diameter dQD by ≈ 0.5 nm per cycle, see Figure 3d, a value in good agreement with the CdSe lattice parameter of 0.608 nm. Figure 3e shows a scanning TEM high angle annular dark field (HAADF) image of HgSe/CdSe QDs that were formed by exposing 4.7 nm HgSe QDs to 3 c-ALD cycles. Clearly, all QDs feature a concentric structure of a bright core surrounded by a 1.0-1.5 nm dark ring, in agreement with the formation of a core/shell structure where the heavier element (Hg) makes the core appear brighter. Finally, Figure 3f shows powder x-ray diffractograms of HgSe and HgSe/CdSe QDs, made by growing CdSe in 3 c-ALD cycles around 6.0 nm HgSe cores. The XRD pattern of the HgSe QDs matches the indicated reference pattern of zinc blende HgSe. Moreover, HgSe/CdSe core/shell QDs have an almost identical diffraction pattern, with slightly narrower diffraction peaks. As zinc blende HgSe and CdSe have nearly identical lattice parameters of 0.608 nm, this suggests that the CdSe shell grows coherently on the HgSe core, which results in a mere extension of the nanocrystals without inducing significant strain. Figure 4 represents the results of a similar structural analysis of HgSe/CdS core/shell QDs. In this case, bright field TEM images indicate that the shape of the originally quasispherical HgSe core QDs becomes more irregular after c-ALD growth. However, by analyzing the projected surface area, we still estimate an increase of dQD by ≈ 0.55 nm/cycle and the formation of a core/shell structure is again confirmed by HAADF-STEM. Although HgSe


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Figure 4: (a-b) Bright field TEM image of initial 5.7 nm HgSe QDs and the HgSe/CdS QDs obtained after 3 c-ALD cycles. (c) size histograms of (red) the original HgSe QDs and the HgSe/CdS QDs obtained after 1,3,4 and 5 c-ALD cycles together. (d) the average nanocrystal diameter as a function of the colloidal ALD cycle, the vertical error bars represent the standard deviation of the QD distribution. (e) STEM-HAADF images of the 4.7 nm HgSe QDs exposed to 3 c-ALD cycles, indicating a concentric structure where a higher Z contrast core is surrounded by a lower Z contrast shell of 1.0-1.5 nm thick. (f) x-ray diffractograms of the initial HgSe QDs (red) and the HgSe/CdS QDs (blue) after 3 colloidal ALD cycles. The vertical (red) lines indicate the diffraction pattern of bulk HgSe. The inset shows the shift to larger angles of the diffraction peaks after shelling with CdS. and CdSe or CdS can form mixed alloys, these STEM-HAADF results indicate that the resulting HgSe/CdE core/shell QDs have a sharp core/shell interface, even if some HgSe-CdE intermixing at the interface cannot be ruled out. The powder XRD pattern now indicates a shift to larger diffraction angles and a slight broadening of the diffraction peaks upon growth of CdS (see Figure 4f). Given the HgSe/CdS lattice mismatch of ≈ 4.3%, this indicates that opposite from HgSe/CdSe, HgSe/CdS core/shell QDs are strained. Similar results have been observed for strained CdSe/CdS and CdTe/CdS core/shell QDs. 45–47 From the absorption spectra and structural characterization analysis, we conclude that the c-ALD procedure developed here results in the formation of HgSe/CdSe, HgSe/CdS and HgSe/CdSe/CdS core/shell QDs in an additive, layer-by-layer manner.



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Photoluminescence of core/shell QDs HgSe QDs could be viable alternatives to Pb-based QDs for near-infrared light emission and amplification due to the lower degeneracy of the band edges, which yields an 8-fold instead of a 64-fold degenerate electron-hole pair. 16 However, core nanocrystals are often poor light emitters due to the presence of surface defects which create non-radiative pathways for electron-hole pair recombination. Shelling can overcome this issue, provided that a defect free core/shell interface is formed, Moreover, leakage of the electron and/or hole wavefunction to the outer surface should be suppressed, which can be accomplished in a so-called type I heterostructure. 47 Using a shell material that has a small lattice mismatch with the HgSe core such as CdSe or CdS is a first step to avoid defect-free interfaces. The second requirement is more difficult to assess a priori as little is known about the band offsets of HgSe with other semiconductors. Especially since size quantization and strain may further shift the energy levels of the small HgSe QDs studied here. However, the minor red shift of the band-edge exciton upon CdS(e) shelling indicates that size quantization in HgSe/CdE core/shell QDs is still mostly determined by the dimensions of the HgSe core, a finding that points towards a type I band alignment. To assess the influence of CdE shelling on the photoluminescence (PL), we started from 4.3 nm HgSe core QDs since such QDs still show a measurable band-edge PL (see Figure 5a). However, as they are partially doped, the PL quantum yield (PLQY) of the band-edge emission is low, amounting to a mere 0.3%. In addition, the emission band is rather narrow, which could reflect the different degree of doping in the QD ensemble. As the larger QDs are more heavily doped, their PL will be preferentially quenched and a suppression of the red side of the band-edge PL will result. Figure 5a and b show that a first effect of CdE shell growth is indeed a combined redshift and increase of the emission linewidth. In fact, comparing emission spectra normalized by the respective PLQY suggests that the apparent redshift consists at least in part of a recovery of the red side of the differentially quenched band-edge emission (see Supplementary Information, Figure S5). In line with the limited 14 ACS Paragon Plus Environment

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Figure 5: Photoluminescence of core (black) and full cycle (a) HgSe/nCdSe and (b) HgSe/nCdS, where the same core is chosen for all experiments with an absolute quantum yield (QY) of 0.3% (see Supporting Information). Color code: red (n=1), green (n=2) and blue (n=3). (c) Photoluminescence quantum yield as a function of c-ALD cycle for different shelling procedures as indicated. (d) Comparison of absolute photoluminescence quantum yield of HgSe/CdSe and HgSe/CdS core/shell quantum dots (full markers) before and (open markers) after annealing at 60o C in solution under inert atmosphere. effect of further shell growth on the band-edge absorption, we also observe little changes to the central wavelength of the band-edge emission after successive CdE c-ALD cycles (see Figure 5a-b). More interesting than the effects resulting from the undoing of the partial doping is the influence of further CdE shelling on the PLQY of the band-edge emission. Figure 5c shows the absolute quantum yield (QY) of different shelling approaches – CdSe, CdS and a CdSe/CdS combination – after excitation with 0.8 µm light, a wavelength corresponding to a photon energy well below the band gap of the CdE shell material. It can be seen that all approaches increase the emission efficiency of the original HgSe core QDs, yet the difference between CdSe and CdS shelling is large. Whereas 3 complete CdSe c-ALD cycles result in a modest 3-fold enhancement of the PLQY, 3 CdS c-ALD cycles under inert conditions lead to a pronounced 16-fold increase. Similar to PbS/CdS core/shell growth by c-ALD, CdS shelling under ambient conditions results in a somewhat poorer performance, which indicates that the chemical purity (oxidation) of the core/shell interface is a factor determining the eventual PLQY. 37 The beneficial role of CdS shelling is confirmed by the addition of CdS 15 ACS Paragon Plus Environment


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shells to HgSe/CdSe QDs, which pushes the PLQY to values comparable to HgS/CdS QDs. 34 Since CdSe and CdS have only a small conduction-band offset, the marked difference between CdSe and CdS shelling probably reflects the better confinement of the hole in the HgSe core in the latter case. Importantly, mild annealing at 60 ◦ C in solution and under an inert atmosphere hardly changes the absorption spectrum of HgSe/CdE core/shell QDs (see Experimental Section and Supplemental Information Figure S6). On the other hand, such a mild annealing treatment systematically increases the PLQY of HgSe/CdS(e) core/shell QDs, similar to what has been reported for HgS QDs. 34 Whereas annealed HgSe/CdSe reach a PLQY comparable to HgS/CdS prior to annealing, the latter exhibit an additional 3-fold increase. Probably, this further increase of the PLQY is due to the elimination of stacking faults or the formation of a more graded core/shell interface induced by the mild annealing. In absolute numbers, HgS/3CdS core/shell QDs thus attain a PLQY of 16.5% for an emission centered at around 1.8 µm. This value is comparable to state-of-the-art PbS cores that emit in the same spectral region. 48

Conclusion We have demonstrated the formation of HgSe/CdE (E=S,Se) core/shell QDs using a roomtemperature colloidal ALD (c-ALD) procedure that consists of successive half reactions. Transmission electron microscopy is used to confirm that shell growth occurs in an additive manner, where approximately one CdE monolayer is grown per cycle of two half reactions. We have found that shell growth removes the n-doping which is intrinsic to the native HgSe cores. A unique strain free core/shell QD system is obtained when CdSe is grown as a shell around HgSe core QDs. Especially in the case of thermally annealed HgSe/CdS core/shell QDs photoluminescence efficiencies are obtained that are comparable to that of PbS QDs emitting in the same wavelength range. These results further prime colloidal ALD


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as a shelling method well suited to form core/shell QDs in the case that high temperature treatments must be avoided and they open up avenues to study HgSe-based QDs through their photoluminescence and use them as IR light emitters.

Acknowledgement Z.H. acknowledges support by the European Commission via the Marie-Sklodowska Curie action Phonsi (H2020-MSCA-ITN-642656), the Belgian Science Policy Office (IAP 7.35, photonics@be), IWT-Vlaanderen (SBO-MIRIS), the Research Foundation Flanders (project 17006602) and Ghent University (GOA n◦ 01G01513) for funding. P.G. acknowledges the Flemish FWO for a postdoctoral fellowship (12K8216N).

Supporting Information Available The Supporting Information contains additional data on the HgSe quantum dots (QDs) and doping levels, the half-cycle reactions, shell growth on completely doped HgSe QDs, the formation of HgSe/CdSe/CdS core/shell QDs, and the photoluminescence of core/shell QDs. This material is available free of charge via the Internet at http://pubs.acs.org/.

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