Adrien Robin

Clément Livache

Sandrine Ithurria

Emmanuelle Lacaze

Benoit Dubertret
email: benoit.dubertret@espci.fr

Emmanuel Lhuillier

Surface Control of Doping in Self-Doped Nanocrystals

Keywords: HgSe, self-doping, Colloidal quantum dot, electrolyte gating, photoresponse, infrared

à la diffusion de documents scientifiques de niveau recherche, publiés ou non, émanant des établissements d'enseignement et de recherche français ou étrangers, des laboratoires publics ou privés.

Introduction

While nanocrystals are attracting a large interest for their tunable optical properties from UV to THz 1 , the control of their electronic properties remains more challenging, thus slowing down their integration into optoelectronic devices [START_REF] Talapin | Prospects of Colloidal Nanocrystals for Electronic and Optoelectronic Applications[END_REF] . The open question of the control of electrical doping is as of today of utmost interest. Indeed, in spite of large progresses over the recent years 3,[START_REF] Ott | Microscopic Theory of Cation Exchange in CdSe Nanocrystals[END_REF][START_REF] Schimpf | Electronic Doping and Redox Potential Tuning in Colloidal Semiconductor Nanocrystals[END_REF] , no systematic strategies have been proposed to tune the electrical doping [START_REF] Stavrinadis | Aliovalent Doping in Colloidal Quantum Dots and its Manifestation on their Optical Properties: Surface Attachment versus Structural Incorporation[END_REF] , [START_REF] Viswanatha | Copper-Doped Inverted Core/Shell Nanocrystals with "Permanent" Optically Active Holes[END_REF] of colloidal quantum dots (CQD). Partial cation exchange [START_REF] Ott | Microscopic Theory of Cation Exchange in CdSe Nanocrystals[END_REF][START_REF] De Trizio | Forging Colloidal Nanostructures via Cation Exchange Reactions[END_REF] is one of the most promising path to achieve the bulk doping of CQD. However, even in the low doping regime of 1 dopant per CQD, this actually corresponds to a large doping regime for bulk semiconductors (> 10 18 cm -3 ). Complementary methods are then needed to achieve a fine control of the carrier density in a range of 0.1 to 1 carrier per nanocrystal.

This question is especially interesting to finely tune the doping level of self-doped nanocrystals (i.e. doping is obtained without process to induce it). Recently, it has been shown that mercury chalcogenides nanocrystals of HgS [START_REF] Jeong | Air-Stable n-Doped Colloidal HgS Quantum Dots[END_REF][START_REF] Shen | and HgS/CdS Colloidal Quantum Dots with Infrared Intraband Transitions and Emergence of a Surface Plasmon[END_REF] and HgSe [START_REF] Deng | Colloidal Quantum dots Intraband Photodetectors[END_REF] are naturally n-doped under ambient conditions [START_REF] Lhuillier | Infrared Photodetection Based on Colloidal Quantum-Dot Films with High Mobility and Optical Absorption up to THz[END_REF] . They typically present 1 to 2 carriers per nanocrystal and consequently exhibit interband and intraband absorption [START_REF] Lhuillier | Infrared Photodetection Based on Colloidal Quantum-Dot Films with High Mobility and Optical Absorption up to THz[END_REF] in the mid or far IR. Thanks to this long wavelength absorption, HgSe CQD have been successfully integrated as mid and far infrared photodetectors [START_REF] Lhuillier | Infrared Photodetection Based on Colloidal Quantum-Dot Films with High Mobility and Optical Absorption up to THz[END_REF][START_REF] Deng | Colloidal Quantum dots Intraband Photodetectors[END_REF] . Nevertheless, the photocurrent modulation remains poor compared to the dark current, and the specific detectivity value (i.e. signal-to-noise ratio) remains smaller than the values reported for photodetectors based on epitaxial grown narrow band gap semiconductors (HgCdTe, InSb, type II superlattice…). Reducing the dark current is then mandatory to improve the photodetection performances. To fulfill this objective, controlling the doping level and the resulting carrier density is of utmost interest. Nevertheless, very little is known on the origin of this doping and on its control.

In this paper, we first describe some of the current limitations of devices based on these self-doped nanocrystals. In an effort to provide a better understanding of the doping, we evidence that the doping level can be finely tuned thanks to a change of surface chemistry. We propose a model where we ascribe the change in doping density to the magnitude of the surface dipoles originating from the ligands. This model is then used to quantify the value of the surface dipole directly from optical measurements. The resulting simple surface chemistry rules are then applied to design a near IR photodetector on soft and transparent electrodes.

Discussion

HgSe nanocrystals combine (i) a narrow band gap originating from the semi-metal character of the bulk material with (ii) a self-doping which leads to optical features in the infrared. For 5 nm CQD (see Figure 1a), the interband band edge is typically around 6000 cm -1 (0.8 eV). This ground state transition, from the 1s state of the valence band to the 1s state of the conduction band, is almost fully bleached. Since the nanocrystals are capped with dodecanethiol (DDT) ligands at the end of the synthesis, only the transition from the 1s state of the valence band to the 1p state of the conduction band appears at higher energy [START_REF] Wang | Electrochromic Nanocrystal Quantum Dots[END_REF] (σ ≈ 8500 cm -1 ), Figure 1b. This transition is better revealed after ligand exchange, see Figure 2. In addition, the CQD present a second optical feature, at 2500 cm -1 (310 meV), relative to a transition within the conduction bands [START_REF] Jeong | Air-Stable n-Doped Colloidal HgS Quantum Dots[END_REF][START_REF] Deng | Colloidal Quantum dots Intraband Photodetectors[END_REF] , see Figure 1b. The latter is an immediate proof of their doped character.

To design a photodetector from this material, a fine control of the doping is mandatory. Indeed, for midinfrared photodetectors, only the intraband contribution of the absorption is useful and an accurate doping value of two electrons per dot is optimal to combine the intraband absorption and prevents dark current through the 1s state. Indeed, a value of 2 carriers per dot fully fills the 1s states: the Pauli blockade thus prevents transport between the 1s states. In this case, the transport only occurs via the 1p states, which requests a thermal or photon activation. On the other hand, for a near infrared photodetector, the intraband transition energy is too low and generates more dark current. A lower doping value is thus more appropriated. In this sense, a careful tuning of the doping level is required. The doping of the HgSe CQD is also confirmed by transport measurements conducted in an electrolytic transistor configuration. The transfer curve of such transistor is presented in Figure 1c and presents an ntype character and a current modulation of roughly two orders of magnitude. Under illumination, the current modulation remains limited (≈ 100%). In narrow band gap materials, the large number of thermally activated carriers generates a high dark current which usually prevails over the photocurrent modulation 13 . The most common strategy to reduce this dark current is to cool the device. We consequently conducted transport measurements at low temperature and observed an Arrhenius dependence of the current, [START_REF] Lhuillier | Thermal Properties of mid Infrared Colloidal Quantum Dot Detectors[END_REF] with a low value of the activation energy, see Figure 1d. The latter ranges from 5 to 20 meV, depending on the size of the CQD. Indeed, the larger the nanocrystal is, the lower is the activation energy. These activation energies are one order of magnitude smaller than the optical intraband transition energy (≈ 300 meV for the small CQD). As a result, only a small improvement of the signal-todark ratio is achieved by cooling this material. Such behavior immediately raises the question of the driving parameter of the activation energy. Since the latter is related to the doping magnitude 13,[START_REF] Liu | -and p-Type HgTe Quantum Dot Films[END_REF] , the origin of the self-doping and the strategies to tune its magnitude need to be evidenced. Mercury chalcogenides are known to present electronic properties sensitive to ambient condition after ligand exchange [START_REF] Lhuillier | Mid-Infrared HgTe/As2S3 FETs and photodetectors[END_REF] , necessary to increase the inter-quantum dot coupling and in turn achieve an electrical conductivity. We thus investigated the effect of ligands in HgSe films on the electrical and optical properties. We conduct solid state ligand exchange from dodecanethiol to short organic ligands (ethanedithiol) or ionic ligands (SCN -, I -, S 2-, As2S3). We observe that not only the conductance of the film switches from an insulating behavior (R > 100MΩ) to a conductive state (R > 1-10 kΩ), but also the absorbance spectrum is affected, see Figure 2a andb. The relative interband contribution of the signal is indeed increased compared to the intraband contribution, which means that the doping level is reduced. This trend is observed for all sizes of CQD but is particularly striking for the smallest nanoparticles, see Figure 2c. These observations bring us to the question of how the presence of a self-assembled monolayer (i.e. ligands) can affect the electronic properties of a surface. While this question has been intensively discussed for bulk materials [START_REF] Deboer | Tuning of Metal Work Functions with Monolayers[END_REF] , the impact of surface dipole on confined nanoparticles remains far less discussed [START_REF] Jeong | Air-Stable n-Doped Colloidal HgS Quantum Dots[END_REF] . For nanoparticles, it has been evidenced using STM [START_REF] Soreni-Harari | Tuning Energetic Levels in Nanocrystal Quantum Dots through Surface Manipulations[END_REF] and X-ray photoemission spectroscopy [START_REF] Brown | Energy Level Modification in Lead Sulfide Quantum Dot Thin Films through Ligand Exchange[END_REF] , [START_REF] Miller | Revisiting the Valence and Conduction Band Size Dependence of PbS Quantum Dot Thin Films[END_REF] that the band energies are affected by the choice of the capping ligands. Indeed, the magnitude of the related dipole can tune the absolute energy position of the gap over almost 1 eV in the case of PbS [START_REF] Brown | Energy Level Modification in Lead Sulfide Quantum Dot Thin Films through Ligand Exchange[END_REF] . Such a shift could also explain the large ligand induced tuning of the doping level for self-doped mercury chalcogenides nanocrystals.

The band diagram of HgSe CQD is given in Figure 3a. Two characteristic energies are present: (i) the interband band edge (BE) energy and the (ii) intraband energy. In presence of ligands, the pristine band diagram of the CQD gets affected by the electric field resulting from the ligand induced dipoles. The electrostatic response of the CQD to a dipole pointing its negative charge toward its surface is to bring positive charges to the surface, which in turn bend the band, see Figure 3b. While the Fermi level of the ligands and of the nanocrystal equilibrates, a shift of the vacuum level occurs. This induces a similar shift of the band with respect to the Fermi level. In other words, a ligands decorated QD presents a shifted band diagram compared to the pristine CQD. The shift of the band is responsible for a change of the electronic population level within the CQD.

Figure 3 a. Band diagram of a self-doped nanocrystal in absence of surface dipole (left) and once surface dipoles are added (central and right part). The addition of dipoles on the nanocrystal surface leads to an electric field which shifts the vacuum level and the relative position of the band gap with respect to the Fermi level (grey dotted line). b. Plot of the conduction and valence band energy versus vacuum as a function of the surface dipole. c. Plot of the 1s state population as a function of the intensity of the surface dipole.

One can estimate the shift of the band thanks to the relation [START_REF] Deboer | Tuning of Metal Work Functions with Monolayers[END_REF][START_REF] Brown | Energy Level Modification in Lead Sulfide Quantum Dot Thin Films through Ligand Exchange[END_REF][START_REF] Yang | Tuning Semiconductor Band Edge Energies for Solar Photocatalysis via Surface Ligand Passivation[END_REF] ( )

        + - =  - ⊥ HgX ligand Hg ligand vac µ ligand µ N E     0 0
where N is the surface density of dipole,

(
) )

ligand
ligand µ ⊥
is the averaged dipole magnitude over the whole nanocrystal volume.

Bulk mercury chalcogenides present large work function values [START_REF] Nethercot | Prediction of Fermi Energies and Photoelectric Thresholds Based on Electronegativity Concepts[END_REF] (around 6 eV for II-VI compounds), which, together with their narrow band gap nature, implies that their conduction band can be lower than the O2/H2O redox couple potential (-1.23 V H + /H2 or 5.7 eV below vacuum level), see Figure 3b. This actually explains why HgS and HgSe CQD can be stable under a negatively charged form and why HgTe, whose band structure is shifted toward the vacuum level, is not. In those large work function conditions, the water is a reducing agent for the QD according to the following redox equation

+ + + ⎯→ ⎯ + 2H 2 / 1 2QD H 2QD 2 - 2 O O
. This hypothesis is further confirmed by the fact that conducting the ligand exchange in presence of water is not affecting the absorption spectrum, see Figure S2. This model also explains why the smallest HgSe CQD (3 nm and less) do not present intraband absorption. Indeed, they are so confined that their conduction band remains above the redox potential of the O2/H2O couple. Similarly, when dipoles are added on the CQD surface, the conduction band of the CQD shifts above the O2/H2O redox couple potential and the previous reduction can no longer occur. The stable form of the CQD becomes the neutral form, see Figure 3b. This allows us to answer a key question regarding the doping origin, which is to know if it rather comes from bulk or surface effect. Our experimental data pledges for a bulk doping, since that the larger the CQD, the smaller is change of the optical signature (ratio of intraband on interband) and in turn the doping is affected by a change of the surface dipole, see Figure 2c.

In the framework of this model, we can quantify the change of carrier density within the 1s state of the conduction band as a function of the surface dipole magnitude. We assume that the dipole density is 3 nm -2 , which corresponds to a full coverage of the surface (1 dipole per cation), consistent with experimental values [START_REF] Anderson | Ligand Exchange and the Stoichiometry of Metal Chalcogenide Nanocrystals: Spectroscopic Observation of Facile Metal-Carboxylate Displacement and Binding[END_REF] reported for chalcogenides nanocrystals. This value is actually an effective value and is averaged over the different exposed facets of the CQD, which present some variations of the capping density.

According to the optical absorption, in presence of DDT ligands the first interband transition is fully bleached which pledges for a Fermi level at the top of the 1s state. This corresponds to 2 electrons per nanocrystal.

The carrier density can be calculated by

 -  = dE E E f E DOS n F FD s ) ( ) ( 2 1
where the factor 2 results from the spin degeneracy, the density of state is given by

2 1 2 ) ( 2 1 2 ) ( S E E E DOS - +         = 
where gamma is the energy level lorentzian broadening and is fitted with the experimental spectrum given in Figure 2a (  = 800 cm -1 ≈ 100 meV) and finally the Fermi Dirac population factor is given by

        - + = - T k E E E E f b F F FD exp 1 1 ) ( . Taking F E E E - = '
and considering the difference ( ) According to the population of the 1s state that we have measured thanks to optical spectroscopy (Figure 2c), we can now estimate the change of the dipole value for each ligand. In other word, this model allows us to estimate the surface dipole related to a ligand using only optical spectroscopy. For the short ligands, this value ranges from 0.7 to 1.1 D, which is consistent with the typical value expected from DFT simulations [START_REF] Arefi | Role of the Head and/or Tail Groups of Adsorbed[END_REF][START_REF] Arefi | Chemical Trends in the Work Function of Modified Si(111) Surfaces: A DFT Study[END_REF] . In the case of S 2-, we can assimilate the dipole length to the Hg-S bond length (2.53 Å) and estimate the gating charge as 0.08 electrons per dipole.

F S Vac E E E -  =  1 , we can write          +   - +         = ' ' ' exp 1 1 ) ' ( 2 1 2 2 2 2 1 dE T k E E E n b Vac s  . Considering that ( ) ( )         -  -  =  ⊥ ligand F S Vac ligand µ N E E E  
We finally take advantage of this surface tunability of the doping to design a near IR photodetector on a soft and transparent substrate. Since we previously demonstrated that S 2-ligands lead to a full disappearing of the intraband transition [START_REF] Jeong | Air-Stable n-Doped Colloidal HgS Quantum Dots[END_REF] , we use this strategy as a way to reduce the dark current in the HgSe based photodetector while the sample is illuminated by a 1.55 µm laser source.

As organic electronics, CQD based devices offer the significant advantage of being processable on a large range of substrates without the constrain of an epitaxial growth. However, organic materials can hardly access the telecom wavelength range. Furthermore, very little work has been dedicated to the integration of narrow band gap CQD on such plastic substrates. The Kagan's group 29 has demonstrated some complementary electronic building block from PbSe CQD, but they did not report their use as photodetector. The Kim's group also pioneered the design of such devices with the use of HgTe CQD [START_REF] Kim | Bottom-and Top-Gate Field-Effect Thin-Film Transistors with p Channels of Sintered HgTe Nanocrystals[END_REF][START_REF] Kim | HgTe Nanocrystal-based Thin-Film Transistors Fabricated on Glass Substrates[END_REF] . Nevertheless, the achieved level of performances remains quite limited, with photoresponsivity of the order of few nA.W -1 for illumination at a telecom wavelength [START_REF] Jang | Transparent and Flexible Thin-Film Transistors with Channel Layers Composed of Sintered HgTe Nanocrystals Nanotechnology[END_REF][START_REF] Kim | Electrical Characteristics of HgTe Nanocrystal-based Thin Film Transistors Fabricated on Flexible Plastic Substrates Thin Solid Films[END_REF][START_REF] Seong | Photocurrent Characteristics of Solution-Processed HgTe nanoparticle Thin Films Under the Illumination of 1.3 µm Wavelength Light[END_REF] . We designed ITO electrodes on a PET substrate, both of them being transparent in the telecom wavelength range, see figure S3. A film of HgSe CQD capped with S 2-ligand has been prepared on these electrodes. The film then presents only the interband absorption with an excitonic structure perfectly optimized for 1.55 µm (6450 cm -1 , see Figure 2a) signal. In a transistor configuration, we observe a large current modulation which can go above two orders of magnitude, see Figure 4a-b. The material remains n-type. Under illumination, the device presents a positive photoresponse and the current modulation can be as large as a factor 2.5, see Figure 4c. The related photoresponsivity is in the 10-20 mAW -1 range, see Figure 4d, which is far larger than the previously reported values for CQD based photodetectors on plastic substrates operating at the same wavelength. However, one key current limitation remains their slow response (> 1s, see Figure 4e andS4), which promptly drops while the signal frequency increases, see Figure 4f.

Future works are needed to find strategies to speed up this response time at least in the 50 ms range to be compatible with video frame rates. The design of shorter devices with a smaller inter-electrodes spacing reducing the transit time will certainly be helpful.

Conclusion

We elucidate some questions regarding the origin of the self-doping in mercury chalcogenides CQD. We propose that the doping results from the bulk reduction of the CQD by water because of their large work function and the narrow band gap of the HgS and HgSe compounds. We further demonstrate that the doping can be tuned in the 0.1 to 1 carrier per dot range thanks to ligand exchange. The related tuning of the surface dipole changes indeed the relative energy of the conduction band with respect to the O2/H2O couple. We propose a quantitative model to determine the magnitude of the dipole related to each ligand and evidence that a change of 1D of the surface dipole correspond to a decrease of the doping level by a factor 10. Eventually, we use this method to design a near IR photodetector. We choose a ligand which fully suppresses the intraband absorption and optimizes the absorption signal in the telecom wavelength range. Compared with previous reports, the photoresponsivity of our detector is enhanced by almost six orders of magnitude.

Methods

Chemicals

Selenium powder (Sigma-Aldrich, 99.99%), selenium disulfide (SeS2, Sigma-Aldrich), mercury acetate (Hg(OAc)2, Sigma-Aldrich), trioctylphosphine (TOP, Cytek, 90%), dodecanethiol (DDT, Sigma-Aldrich), oleic acid (Sigma-Aldrich, 90%), oleylamine (Acros, 80-90%), arsenic trisulfide (As2S3 ,Strem, 99.9%), propylamine (Sigma-Aldrich, 98%), n-methylformamide (NMFA, VWR, 98%), lithium perchlorate (LiClO4, Sigma-Aldrich, 98%), polyethylene glycol 6k (PEG 6k, MW=6kg.mol -1 , Fluka), ethanol absolute anhydrous (99.9%, Carlo Erba) acetone (Carlo-Erba, 99.8%), ammonium iodide (NH4I, Sigma-Aldrich, 99%), ammonium thiocyanate (NH4SCN, Sigma-Aldrich, 99%), 1,2-ethanedithiol (EDT, Fluka, 98.0%), sodium sulfide nonahydrate (Na2S,9H2O, Sigma-Aldrich 98.0%).

All chemicals are used as received, except oleylamine which is centrifuged before use. Precursors 1.54 g of Se powder is mixed in 20 mL of TOP. After sonication, a transparent solution is obtained.

Nanocrystal synthesis

In a 100 mL three-neck flask, 1 g of mercury acetate is mixed with 20 mL of oleic acid and 50 mL of oleylamine. The solution is degassed at room temperature, then at 80°C for 30 min. The atmosphere is then switched to Ar. At 90°C, 3.2 mL of TOPSe (AM) is quickly injected and the solution turns dark. After 5 min of reaction, 1 mL of dodecanthiol is added to quench the reaction while the heating mantle is removed. The content of the flask is split into two 50 mL Falcons and equivalent volume of ethanol is added to precipitate the solution. After centrifugation, the clear supernatant is discarded and the formed pellet redispersed in fresh toluene. The cleaning procedure is repeated two other times.

Mercury compounds are highly toxic. Handle them with special care! Film deposition and solid state ligand exchange procedure

The cleaned HgSe QDs suspension in toluene is typically diluted between 20 to 100 times in a 9:1 volume ratio hexane:octane mixture. The resulting suspension is light brown and transparent, and has a target optical density at 15 000 cm -1 ranging between 0.5 and 1. Meanwhile, a ligand exchange solution in ethanol is prepared. For NH4SCN and NH4I, a 1% in weight solution of the compound is prepared; for Na2S,9H2O, a saturated solution is prepared, then filtered at 0.2 µm and subsequently de-saturated by doubling the solvent volume; for As2S3, 1 mg of the latter is partially dissolved in 1 mL of propylamine, a 1% in volume of the latter solution in ethanol is then prepared; for EDT, a 1%v solution in ethanol is prepared. A 50 to 100 nm thick and homogeneous film is then prepared by repeating the following procedure 10 times: a 5 µL droplet of the QDs suspension is drop-casted in air onto the electrodes, then dried in open air. The sample is dipped 1 min in the ligand exchange solution, then rinsed in pure ethanol for 1 min, then dried in open air. This monolayer-by-monolayer procedure produces a clean crack-free film whose roughness is on the order of the nanocrystal size.

Electrodes on Si/SiO2 wafer

The surface of a Si/SiO2 wafer (400 nm as thickness) is cleaned by sonication in acetone. The wafer is rinsed with isopropanol and finally cleaned using an O2 plasma. AZ 5214E resist is spin-coated and baked at 110°C for 90 s. The substrate is exposed under UV though a pattern mask for 2 s. The film is further baked at 125°C for 2 min to invert the resist. Then a 40 s flood exposure is performed. The resist is developed using a bath of AZ 326 for 32 s, before being rinsed in pure water. We then deposit a 3 nm chromium layer and a 40 nm gold layer using a thermal evaporator. The lift-off is performed by dipping the film for in acetone for 1h. The electrodes are finally rinsed using isopropanol and dried by an air flow. The electrodes are 2 mm long and spaced by 20 µm.

Electrodes based on ITO on PET

Indium tin oxide (ITO) coated on polyethylene terephthalate (PET) (≈ 80 nm coating with a 60 Ω/cm² resistance) sheet are purchased from Sigma-Aldrich. The film is rinsed using acetone and then isopropanol before being dried. AZ 5214E resist is spin-coated and then baked for 90 s on a 110°C hot-plate. The film is then exposed to UV for 4 s though a shadow mask. The resist is then developed for 45 s in AZ 726 and rinsed in pure water. The naked ITO is then etched using 25% HCl solution for 15 s and then quickly rinsed in pure water. The lift-off of the resist is made by dipping the substrate in acetone for 5 min and then rinsing the film with isopropanol. The designed electrodes are interdigitated electrodes with 50 µm spacing. Each electrode is itself 50 µm large and 1 mm long. The total active area is 1 mm 2 . A side gate electrode is also present for electrolyte gating.

Electrolyte preparation

50 mg of LiClO4 are mixed with 230 mg of PEG on a hot plate in an Ar filled glove box at 170°C for 2 h.

Transistor fabrication

Electrical measurements are conducted in air and made using a Keithley 2634b as source-meter. Transistor measurements are conducted by biasing the drain source under low bias (10 or 50 mV) and applying a varying gate bias. The gate bias step is 1 mV and the sweep rate 2 mV.s -1 for electrolyte gating.

Material characterization

For TEM imaging, the solution of CQD is drop-casted on a copper TEM grid and degassed overnight under vacuum. The TEM images are then obtained from a JEOL 2010 microscope. IR optical absorption is measured in a transmission configuration using a Bruker Vertex 70 spectrometer. A global source equivalent to a 700°C blackbody is used as source and DTGS detector. Spectra are typically acquired between 10 000 cm -1 and 380 cm -1 with a 4 cm -1 resolution and averaging over 64 spectra.

Figure 1 a

 1 Figure 1 a. TEM image of HgSe colloidal quantum dots. b. Infrared spectrum of the same HgSe CQD. c. Transfer curve (drain and gate current as a function of gate voltage) for an electrolytic transistor which channel is made from HgSe CQD capped with sulfide ligands. The drain bias is set to 10 mV. The gate bias step is 1 mV. d. Current as a function of temperature for a thin film made from HgSe CQD. The experimental points are round (large CQD > 10 nm) and square (small CQD of 5 nm). The solid lines are Arrhenius fits of the experimental data.

Figure 2 .

 2 Figure 2. Absorption spectra of a. small sized (4.7 nm) and b. medium sized (5.7 nm) HgSe nanocrystals capped with different ligands. c. Plot of the ratio of the intraband signal divided by the fitted magnitude of 1s interband signal for two sizes of nanocrystals. The related population of the CQD is given on the right scale for the small sized CQD (4.7 nm diameter).We can relate the relative absorption of the interband and intraband signal to the population of the 1s state of the conduction band. Assuming a density of state of two electrons on the 1s states[START_REF] Lhuillier | Optical Properties of HgTe Colloidal Quantum Dots[END_REF] due to spin degeneracy, the interband absorption, inter  relates to the population of this level, ) 1 ( s n 18 , though the



  µ ⊥ is the component of the intrinsic ligand dipole normal to the surface, 0 is the vacuum permittivity, ligand  the dielectric constant of the ligand, ligand Hg µ - the dipole relative to the interface between the crystal surface and the ligand and HgX  the dielectric constant of the nanocrystal. For the sake of simplicity, we assume that this surface band shift is translated to a band shift in volume

Figure 4 a

 4 Figure 4 a. Transfer curve and b. IV curves under different gate bias of an ion-gel transistor made from a film of HgSe CQD capped with S 2-and deposited on ITO electrodes on a PET substrate. The inset of a. is a picture of the film and the electrodes. c. Photoresponse under dark conditions and under illumination by a 1.55 µm laser at different laser intensities, d. related power dependence of the photoresponsivity, e. change of the current signal under light illumination with a 808 nm laser as a function of time and f. dependence of the photocurrent versus illumination modulation frequency of a film of HgSe CDQ capped with S 2-on ITO electrodes on a PET substrate.

Acknowledgments

We thank Olivier Pluchery for fruitful discussions. We thank Agence Nationale de la Recherche for funding through grant Nanodose and H2DH. This work has been supported by the Region Ile-de-France in the framework of DIM Nano-K. This work was supported by French state funds managed by the ANR within the Investissements d'Avenir programme under reference ANR-11-IDEX-0004-02, and more specifically within the framework of the Cluster of Excellence MATISSE.

Supplementary information

The Supporting Information is available free of charge on the ACS Publications website at DOI: Additional data including material characterization and setup for measurement are given in the supplementary information.


Marion Dufour

Violette Steinmetz

Eva Izquierdo

Thomas Pons

Nicolas Lequeux

Emmanuel Lhuillier

Laurent Legrand

Maria Chamarro

Thierry Barisien

Sandrine Ithurria
email: sandrine.ithurria@espci.fr

Engineering Bicolor Emission in 2D Core/Crown CdSe/CdSe 1-x Te x Nanoplatelet Heterostructures Using Band-Offset Tuning

Keywords: nanoplatelets, cadmium chalcogenides, bi color emission, colloidal heterostructure, exciton transport

à la diffusion de documents scientifiques de niveau recherche, publiés ou non, émanant des établissements d'enseignement et de recherche français ou étrangers, des laboratoires publics ou privés.

INTRODUCTION

Nanocrystals (NCs) have attracted a huge synthetic effort because of their tunable band gap from the UV to the THz range 1 . Their use as phosphor for display has recently become the first mass market application of these nanomaterials. In such devices, a blue GaN/InGaN diode is used to pump red and green quantum dots (QDs) instead of the Ce:YAG phosphors in order to expand the color gamut. The use of two populations of QDs is nevertheless constraining since their ratio has to be finely adjusted to take into account their different photoluminescence (PL) efficiency. In this sense, having a single NC with a double emission will be highly desirable.

Bicolor emitting NCs have already been reported. 2,3 They have been proposed as dual source for classic (ie bunched photon) and quantum light (ie antibunched photon) 4 or, to cover both visible and infrared emissions. 5 To obtain such bicolor emission two main strategies have been explored. They rely on QDs heterostructure 6,7 and on doped QDs. 8,9 In particular, the observation of a phosphorescence signal from Mn doped NCs have been extensively studied. 10 However, in most of these reports, the two PL efficiencies are very different which prevent their use as phosphor for display.

For the display application the NCs, in addition of a high PL efficiency, need to present a narrow PL signal in order to obtain a large gamut. In this sense, the 2D nanoplatelets (NPLs) are the NCs exhibiting the narrowest PL signal reported with a linewidth barely above kbT at room temperature. 11 Since their introduction in the late 2000, 12,13 a huge synthetic effort 14 has been devoted to grow 2D heterostructures which could preserve the lack of inhomogeneous broadening 15 and the associated PL narrowness or controlled cut-off wavelength for photodetection 16,17 . Two growth approach can be studied to obtain those properties. The platelets can be grown in the thickness or 18,19 , they can be grown laterally. In that case, the thickness and the confinement remain constant. With a type I CdSe/CdS band alignment, the core/crown geometry has been demonstrated to be an efficient light concentrator. 20,21 However, this is the type II CdSe/CdTe band alignment 22,23,24 which generates the largest interest because of the possibility to obtain an emission though a charge transfer state. 25 In these NPLs, the emission appears strongly redshifted with respect to the absorption of both CdSe and CdTe. This has so far been attributed to the formation of an exciton at the interface of the core and the crown where the hole stays in the CdTe crown and the electron in the CdSe core. Here, we want to use this core/crown geometry to achieve bicolor emission. Indeed, these heterostructures combine several potential advantages such as (i) a large cross section for each material which can be tuned independently from the other and (ii) three recombination paths for the exciton. It can recombine within the core, the crown or at the interface. This can, in theory, lead to up to three emission wavelengths. 26 However, the practice is much more complex, and a bicolor emission already requires a fine tuning of the recombination state and charge transfer between the different materials.

In CdSe/CdTe core/crown NPLs, only interface emission has been observed so far. This results from the fastErreur ! Signet non défini. (<1ps) / immediate 27 (<100 fs) delocalization at the interface of the electrons and of the holes respectively in CdSe and in CdTe. These time ranges have to be compared with the ns photoluminescence (PL) lifetime of cadmium chalcogenide NPLs at room temperature. 28,29 Thus, recombinations in the core or in the crown are unlikely to occur. In this paper we grow CdSe/CdSe1-xTex core/crown NPLs and demonstrate that by carefully choosing the crown composition and lateral extension, a dual color emission can be obtained.

DISCUSSION

To grow CdSe/CdSe1-xTex core/crown NPLs, we first synthetize CdSe core NPLs using the procedure of the literature. 22 These NPLs have a lateral extension of 9nmx22nm and a 4 monolayers (MLs) thickness ( 5planes of cadmium and 4 planes of selenium) which lead to a PL peak at 510nm. The growth of the crown is ensured by mixing Cd precursors and CdSe core NPL at high temperature (215°C) and by slowly injecting the Se and Te precursors (TOPSe and TOPTe respectively). The thickness of the NPLs is unchanged (1.3nm) while their lateral extension now reaches up to 20 nm x 100 nm, see the electron microscopy image in Figure 1c, S1 and S2. A series of NPLs with an increasing ratio of Te in the crown have been synthetized from pure CdSe up to pure CdTe with a 10% step of Te content. More details about the synthesis and characterization of the materials are given in the methods and supporting information, see Figure S1-S4 and table S1. The room temperature (RT) absorption spectra of the 11 core/crown NPLs are presented in figure 1a and show two types of optical features. First, the peaks relative to the CdSe core (510nm associated to the heavy hole exciton transition and 478 nm associated to the light hole exciton transition) can be observed for all the NPLs. A second feature resulting from the crown appears at redder wavelengths. The energy of this red peak is clearly non monotonic with the Te content. This is the signature of a bowing effect and it will be further discussed later. The RT PL spectra excited at 420 nm and associated with CdSe/CdSe1-xTex core/crown NPLs are presented in Figure 1(b) and Figure 2. From sample to sample the spectrum drastically changes. Four main behaviors can be observed as a function of the Te content. (i) When 0<x<0.2, we observe a single large PL peak with a typical full width half maximum (FWHM) of 160 meV, see Figure 2(c.). PreviouslyErreur ! Signet non défini. , 30 , it has been shown that even very little Te content leads to a significant broadening of the peak compared with pure CdSe core as well as an asymmetric shape. This is the signature of the exciton trapping due to the inhomogeneous content of the alloy. 31 (ii) For 0.3<x<0.5, we obtain a single peak but it is narrower. The peak is also asymmetric and presents a long tail at lower energy. When x=0.5 the FWHM is equal to 19 nm, it is a typical signature of the NPLs type I emission. The peak observed in this range comes from a recombination inside the crown. The attribution of these two kinds of peaks to the alloyed crown emission is further confirmed by the fact that the PL peak is also subject to the bowing effect similar to the one observed in absorption. (iii) When 0.6<x<0.9, the spectra present two emission peaks with very different line widths. The high energy peak presents a limited Stokes shift with the absorption and a FWHM below 60 meV. It corresponds to the crown alloy emission. A contrario, the second peak at lower energy is clearly red-shifted compared to the absorption and presents a FWHM of more than 130 meV. This large peak is similar to the one observe for (iv) x=1 which has been studied and attributed to either the signature of the type II band alignment 26 or at least a recombination through interface statesErreur ! Signet non défini.. ) and (f.) emission and scheme of the four types of emission depending on the crown in the core/crown heterostruture.

In the third region (iii) and, for intermediate contents of Te and Se, we observe a dual color emission made of two overlapping peaks, see Figure 2(e). These two peaks can be fitted by two Gaussian centered at 588 nm and 635 nm respectively and a full width at half maximum of 17 nm and 57 nm. We estimate that ≈25% of the signal magnitude at 588 nm (maximum of the blue peak) is actually resulting from the overlap with the red peak. To ensure that these two peaks indeed come from the same NPLs, excitation of the photoluminescence spectra centered (PLE) on the two emissions have been performed (see supporting information, figure S5) and show a perfect overlapping.

We then investigate the effect of the lateral extension of the crown on the PL emission, see Figure 2(a) and 2(b) (and supporting information figure S2 andS3). The PL peak is already strongly redshifted after the injection of 250µL of precursors. This corresponds to an increase of the NPLs surface of approximately 100%. This first large peak corresponds to the recombination through the interface. The second peak related to the crown emission only appears when the crown gets larger (typically an increase of 200% of the surface). This peak becomes more and more predominant as the crown size further increase. (inset of figure 2(b)). When we synthesize a crown with a spatial extension of only few nm the interface core/crown play an important role and the unique contribution to the emission is the interface contribution. The 2D recombination of the crown exciton becomes possible when the crown size is large enough: in NPLs, theoretical studies have indeed shown that the exciton oscillator strength is an increasing function of the lateral size before it reaches a plateau. 32 This may partly explain the progressive emergence of the crown emission with growing extension. The key role of the crown size on the relative magnitude of the two emissions is certainly the reason why Demir's group did not report about this dual color emission within CdSe/CdSe1-xTex core/crown NPLs. They, indeed, focused on smaller crown size and lower Te content to ensure a fine tuning of the emission from green to red with an efficient emission quantum yield.Erreur ! Signet non défini.

Figure 3: (a.) Low temperature (T~5 K) luminescence (blue-grey) and absorption (blue) of thin solid films prepared from initial solution of CdSe/CdSe1-xTex core/crown NPLs with x=0.6. For comparison the luminescence of NPLs in solution is also sketched (red-grey). Inset (b.): absorption of thin films for x=0.6 at ambient temperature (red) and low temperature (blue). (c.) Evolution of CdSe/CdSe1-xTex NPLs luminescence (x=0.6) with increasing dispersion on glass slides. The degree of dilution of used solutions (impregnation) compared to original solution is indicated above each spectrum. (d.) Typical luminescence spectrum, at ~5 K, of a single CdSe/CdSe1-xTex core/crown NPL (x=0.6) (dark blue). Dark green curve is a focus on the emission of the interface state (x 10). Inset (e.): Luminescence dynamics associated to the Xcrown line (562.7 nm) at ~5K and monoexponential fit of the decay (green solid line).

In order to get insights into the intra-platelets nature of the bicolor emission, micro-PL experiments are conducted on CdSe/CdSe0.4Te0.6 NPLs for which the two PL peaks have the closest magnitude.. Spectroscopy on diluted NPLs films are carried out to ensure identification of individual optical responses. The experiments are performed at low temperature (T5K) to reduce homogeneous broadening of the emission line that results from exciton-phonon coupling 33,34 . Figure 3(a.) compares the absorption and luminescence spectra of a low optical density film. Ambient absorption spectrums of the film as well as the luminescence of the solution are also shown. At ambient temperature, the solution and thin film have nearly identical absorption. When lowering temperature to 5 K, the film spectrum experiences a rigid blue shift (~100 meV) (see Figure 3(b.)). Such shift is expected in bulk semiconductors 35 and is primarily explained by band shifting resulting from lattice contraction and varying electron-lattice interaction. For comparison a shift of ~75 meV was measured in CdSe NPLs within the same temperature interval. 29 Slight narrowing of the absorption linewidth is also observed and is sufficient to reveal the light hole-electron (lh-e) and heavy hole-electrons (hh-e) transitions respectively from the crown and the core around 495 nm. 24 In the experiments the excitation is tuned at 515 nm (c.w. argon laser) so that only higher-lying states of the crown exciton progression get populated. In this configuration the interface state is exclusively populated from the Xcrown exciton reservoir by exciton transfer to the core. 36 Both emission of dense films and solution present the same characteristic structure with two bands Xint and Xcrown. The energy difference between the two bands (Xcrown and Xint) passes from ~140 meV for NPLs in solution to 158 meV in spread NPLs at T=5 K showing that, again, temperature induced effects do not considerably modify the energy landscape of the hetero-system. Figure 3(c.) shows the evolution of low temperature PL for varying NPLs coverage rates. As the NPLs density decreases the inhomogeneous Xcrown line splits into sharp peaks while the Xint emission remains identical (see Figure 3(c.)). By diluting the sample, we measure the signal from a single NPL characterized by an extremely narrow emission. Several observations evidence the intrinsic nature of the measured PL. The Xcrown linewidths, 0, are found in the range 380-450 eV which wells compare to the values of emission widths reported in pure CdSe NPLs. 29,37 Second the emission is subject to blinking and spectral diffusion that is responsible for well resolved shifts,  of the emission. In the brightest NPLs we estimate  0 within the acquisition time hence the reported 0 values should be considered as upper values of the actual homogeneous widths, the latter being very likely slightly less. Measurements of the PL intensity decay for the Xcrown line are also performed at the single NPL level. The decays are fitted using single mono-exponential functions and measurements on ~20 NPLs provide a lifetime range from 70 ps to 190 ps (a typical trace is shown in figure 3(e.)). Note that in our PL traces there is no signature of longer constants. However due to the relatively poor intensity dynamic of the detector (Streak Camera) It is not possible to exclude their presence. Our results indicate that such long time contributions, if present, have a much weaker contribution (at least one order of magnitude) at low temperature compared to the ambient situation. The measurement of sub-nanosecond decay times is a strong evidence that the Xcrown line corresponds to the radiative recombination of a direct exciton generated in the crown. Similar decay times (~200 ps) are reported at 20K in single CdSe NPLs where the emission is only determined by the lowest band edge exciton. 29 Single object studies also directly inform about the intrinsic character of the bicolor emission. It is indeed remarkable that the interface state band Xint is always present in the single NPL emission as shown in figure 3(d.). The associated signal is observable in the c.w. luminescence spectrum but not in the time resolved emission (using the synchroscan camera) whereas intensity would have been expected well above the measurable threshold for sub-nanosecond decays, which confirms the long-lived nature of the state. A shared feature of single NPLs spectra is the broad spectral extension, E, of the interface state emission (FWHM of ~100 meV) as well as its relative weak amplitude compared to the crown emission. We also note a more or less apparent structure in the spectrum which is found NPL dependent. In a very simple scheme, if intrinsic responses are measured, the ratio, , between area under the crown line and interface state line in the single NPL spectrum should necessary roughly equals the ratio of their counterpart contributions in the inhomogeneous spectrum. This is actually verified: we find very little variations of  within the observed NPLs and find ~0.6-0.65 in a series of ~5 isolated NPLs exhibiting particularly well resolved emissions, directly comparable to the one extracted from the film PL following spectrum deconvolution. At low temperature single platelet emission is thus composed of a narrow exciton line similar to a QWs free exciton line, accompanied by a broad low energy band with much longer dynamics. In epitaxial double QWs, low temperature inhomogeneously broadened indirect exciton luminescence has width in the meV range 38,39 , which makes a difference of nearly two orders of magnitude. One possible explanation is that the emission originates from the recombination of an electron-hole pair with one of the carriers or both in deep traps. E would characterize the distribution of the energy levels associated to the traps. As investigation is at the single NPL level the following important conclusion can be drawn: the observation of narrow Xcrown lines is a strong indication that neither hole or electron trapping due to alloy potential fluctuations, arises in the crown volume at low temperature (hence also at higher temperature). 40 As narrow free excitonic lines are also observed at low temperature in pure CdSe NPLs 29 , we also come to the conclusion that trapping prevails at the interface between the core and the crown material.

In the following, we conduct time-resolved PL spectroscopy at RT to further quantify the emission dynamic and characterize the dual emission. Figure 4(a) follows the luminescence spectrum at early time (<1.5 ns) after excitation for CdSe/CdSe0.4Te0.6 core/crown NPLs. High excitation energy is used (λ=407 nm->E≈3 eV) so that both the core and the crown are excited. The density of excitation is about 4µJ/cm -2 . Over the first few ns, only the PL peak relative to the crown is observed while the red broad peak is extremely weak and hardly observable (see figure 4(a)). This strongly differs from the CdSe/CdTe core/crown objects for which the transfer of the electron is expected to occur over a sub ps time scale. 25,27 Time dependence of the blue peak emission at short time scale is plotted in Figure 4(b) for three samples, x=0.5 when only the crown emits, x=0.6 when both the crown and the interface emits and x=0.8 when the crown barely emits. We underline very distinct behaviors. At high Te content (x=0.8), we observe a very fast decay with a constant below the instrumental resolution (~15 ps), which already corresponds to ~30% of the integrated time resolved PL signal. When the Te content is reduced (x=0.5, 0.6) the dynamic considerably slows down and the decay can be fitted with two exponentials with characteristic times of 35 ps and 160 ps respectively. Moreover, the weight of those fast components in the total integrated time becomes much weaker, only a few percent. We then investigate the dynamics at longer time scale (>ns) for the crown emission (Figure 4(c)) and the interface emission (Figure 4(d)). 41 Decays show in general a multi-exponential dynamics that depends on x, the Te content. In particular, crown and interface emission decays are very similar for the sample CdSe/CdSe0.4Te0.6. They correspond to a multi-exponential dynamic including four components. Decay times (and respective weights) of 6 ns (0.627) , 60 ns (0.290), 250ns (0.079) and 2.0s (0.004) are measured for the crown peak whereas the values of 18 ns (0.38), 80ns (0.44) , 350ns (0.15) and 1.7s (0.03) characterize the interface peak dynamics. Note that the first three constants as well as their weight are determined with a rather low uncertainty (of a few percents) whereas it is higher but keeps moderate (~10%) for the last one, due to the limited extension of the observed temporal window. In a general manner the presence of a distribution of decay times may be related to exciton transport via hopping. 42 It thus would be tempting to ascribe the lengthening of the decays to the introduction of traps on the percolation path of excitons through alloying. However RT observation of multi-exponential dynamics with constants in the same range in the relaxation of single homo-NPLs of CdSe 29 precludes the assignment of decay lengthening to the introduction of the alloy alone. Hence the complex structure of the dynamics at longer scale will not be discussed further here. The main point of the study is the observation of a concomitant shortening of the fast component in the decay (that also gain weight) and strong decrease of the integrated Xcrown PL with increasing x. This is clearly the indication that non radiative processes become important when the crown composition gets closer to pure CdTe. At that point it is impossible to explicitly identify the process as corresponding to the fast electron transfer towards the core. However we emphasis that by contrast the reduced effect for intermediary compositions (x~0.6) makes favorable the populations balance towards Xcrown compared to Xint.

8). A narrow band pass filtered is used to collect photon around 590nm (mainly crown emission). (c.) PL decay at long time scale

for CdSe/CdSe1-xTex for three different x ratio of Te (x=0.5, 0.6, 0.8). The same filter as for (b.) is used. (d.) PL decay at long time scale for CdSe/CdSe1-xTex for two different x ratio of Te (x=0.6, 0.8) (mainly associated to the interface emission). For x=0.5, the core/crown NPLs don't exhibit emission from the interface. A narrow band pass filtered is used to collect photon around 630nm. In all cases, measurements are acquired at room temperature and the excitation wavelength is 407nm.

We have a clear dependence of the PL spectra as a function of the Te content in the crown and want to further elucidate its origin. From the absorbance and PL spectra, we extract the valence (ΔVv) and conduction (ΔVc) band offsets between the CdSe core and the CdSe1-xTex crown when the dual emission happens (which means for 1<x<0.5).Erreur ! Signet non défini. The procedure is further detailed in the figure S6 of the supporting information. The value for ΔVv is always large (≈0.5eV) and poorly depends on the Te content, see Figure 5(a). As a result, the holes are always promptly transferred into the crown. This explains in particular why there is no PL signature from the CdSe core. The situation is very different for the electron where ΔVc spans from 0.15eV to 0.35eV depending on the Te content. This value has to be compared with the exciton binding energy since their relative magnitude will drive the interface charge transfer 43 . The exciton binding energy is much larger in 2D NPLs than in QDs, and is typically around 250meV. 44,45 Thus, depending on the Te content, the electron transfer driving force (ΔVc) can be smaller or larger than the strength of the exciton Coulomb attraction (ie the exciton binding energy Eexc). When ΔVc> Eexc which occurs at high Te content, the electron promptly gets the core/crown interface and the interface emission is only observed. On the other hand, when x is around 0.6 (ΔVc< Eexc) the electron remains coulombically retained in the crown. These band offsets come as a consequence of the band bowing effect occurring in alloy CdSe1-xTex highlighted in figure 5(a). The band gap of the crown can be fitted with the equation ECdSe1-xTex=xECdTe+(1-x)ECdSe-bx(1-x), where b the band bowing parameter is equal to 0.78 (see figure 5 Thus positioning of the conduction band offset in the range of the exciton binding energy make the emission of the interface and of the crown having similar weight in the integrated time resolved luminescence and leads to the observation of bicolor emission. The evolution of the dynamics with crown composition corroborates this picture.

CONCLUSION

As a summary, we can rationalize the observation of bicolor emission by the fact that adjustment of the crown composition allows fine tuning of the conduction band offset compared to the excitonic binding energy. The latter is only made possible by the presence of bowing effect. In a consistent way short time dynamics show that when Coulomb attraction is stronger and binding states energetically more favorable, non radiative quenching of the crown exciton is also drastically reduced, very likely due to less efficient charge transfer towards the core. In addition unambiguous demonstration of the bicolor emission by individual nanoplatelets is done. This work unveils that model 2D geometries allows new possibilities at the single object level, for the design of multicolor emission compared with 0D quantum dots. Future works will be dedicated to the design of dual emitter with a larger energy splitting.

METHODS

CHEMICALS

Octadecene (ODE) (Aldrich, 90 %), Cadmium acetate dihydrate (Sigma Aldrich, 98 %), Cadmium oxide (Strem, 99.99 %), myristic acid (Aldrich, >99%), Oleic Acid (OA) (Aldrich 90 %), Trioctylphosphine (TOP) (Aldrich, 97 %), Selenium (Strem Chemicals 99.99 %), Tellurium (Aldrich, 30 mesh, 99.997 %), Hexane (Carlo Erba, 95 %), Ethanol (Carlo Erba, 99.9 %), Methanol (Carlo Erba, 99,9%) PRECURSOR PREPARATION TOP Se 1M In a glove box, 50 ml of TOP are stirred with 3.95 g of selenium, overnight up to a complete dissolution of the powder.

TOP Te 1M

In a three neck flask, 2.54 g of Te powder are mixed with 20 mL of trioctylphosphine and degassed under vacuum at room temperature until the pressure gets reduced below 1 mBar. Then the flask is heated under Ar, at 275°C until the black Te powder gets fully dissolved. The orange solution is then cooled and turns yellow. Once the temperature is below 40°C, the solution is further degassed under vacuum for 10 min and finally stored in an air free glove box.

Cadmium myristate

In a 50 ml three neck flask, 1.28 g of CdO (10 mmol) and 5.5g of myristic acid (24 mmol) are degassed for 30 minutes at 70 °C. Then, under argon flow, the mixture is heated up to 200 °C for 30 minutes until it becomes colorless. At the end of the reaction 25 ml of methanol are added between 60 °C and 70 °C in order to solubilize the excess of myristic acid. The cadmium myristate is then precipitated in a centrifuge tube with addition of excess methanol. The washing procedure is repeated at least three times and the cadmium myristate is finally dried overnight under vacuum.

NANOCRSYTAL SYNTHESIS Core CdSe 5 ML synthesis

The procedure for the synthesis of the NPLs is similar to the one described by Pedetti et al.Erreur ! Signet non défini. 24 mg of Se powder, 340 mg of Cd(Myr)2 and 25 mL of ODE are mixed in a 50 mL three-neck flask. The solution is degassed under vacuum for 10 minutes at room temperature. Then, under an argon flow, the temperature is set to 240°C. 110 mg of Cd(Ac)2 are added when the color becomes orange (200°C approximately). After 10 minutes at 240°C, the reaction is quenched with 200 µL of oleic acid and then quickly cooled to room temperature. As obtained, the solution is a mix of dots and platelets. After adding 25 mL of hexane and 25 mL of ethanol, the platelets are selectively precipitated by centrifugation at 6000 rpm for 5 minutes. The orange supernatant is removed and the precipitate is dispersed in hexane. The amount of hexane is chosen so that 100 µL of the NPLs solution diluted in 3mL of hexane have an optical density of 0.6 at 512 nm in a 1 cm cuvette. CdSe/CdSeTe core/crown heteronanoplatelets 1.5 mL of the CdSe NPLs solution are centrifuged, the platelets are then dried and dispersed in 2 mL of ODE. In a 25 mL three-neck flask, 24 mg of Cd(Ac)2, the NPLs in ODE and 45 µL of oleic acid are mixed and vacuumed at room temperature for 10 minutes. Under an argon flow the temperature is set to 215°C. In parallel, a mix a variable ratio of TOPSe (1 M) and TOPTe (1 M) in ODE is prepared (total anion concentration of 0.025 M). When the temperature is reached, this second solution is injected at a rate of 2 mL/h into the reaction media. When the desired injection volume is reached (between 0.25 mL and 2 mL), the mixture is cooled to room temperature and washed once with ethanol. The core/crown CdSe/CdSexTe1-x NPLs are suspended in hexane and stored in the dark to prevent them from photodegradation. Aliquots are taken every 0.25 ml of the second mixture injected in the three neck flask.

MATERIAL CHARACTERIZATION

UV-visible spectrum have been acquired with a Shimadzu UV-3600 spectrometer. Photomuminescence (excitation at 420 nm) and excitation spectra are obtained from an Edinburgh Instrument spectrometer. The transmission electron microscope is a JEOL 2010. The grids are prepared by a dropcast of a solution of NPLs dispersed in hexane and degassed overnight under secondary vacuum. X ray diffraction is conducted by dropcasting a solution of NPLs on a Si wafer. The diffractometer is a Philips X'Pert based on the emission of the Cu Kα line operated at 40 kV and 40 mA. Energy Dispersive X ray spectroscopy (EDS) is performed using a Oxford EDS spectrometer in a FEI Magellan scanning electron microscope. Micro-photoluminescence Experiments are performed with a confocal-like setup using a 0.6 numerical aperture microscope objective to focus light on the nano-object and collect light in reflection configuration. Scattering light from the laser is suppressed placing a sharp edge dielectric filter from Semrock company on the collecting path. Imaging of the sample plane is made on the entrance slits plane of an imaging Acton SP2760 spectrometer (from Princeton Instruments) and luminescence detected with a nitrogen-cooled CCD camera (SPEC10 -PI). For low temperature studies core/crown NPLs are dispersed on glass slides (thickness ~100 m) and glued on the cold finger of a micro-photoluminescence cryostat from Oxford Instrument company. Time resolved spectroscopy Time resolved spectroscopy was performed resorting to two distinct methods. Sub-nanosecond decays of the photoluminescence were measured using a streak-camera (C5680 model from Hamamatsu incorporating a M5675 synchroscan unit) coupled to the Acton SP2760 spectrometer. In this configuration excitation lightpulses of ~2 ps duration -is the second harmonic of a titane-sapphire laser operating at 82 MHz ; a temporal resolution of ~15 ps is then typically obtained following dispersion of the emitted light through the spectrometer. Classical gated photon counting provided access to long range decays. A gate scanning unit including a discriminator (SR400 from Stanford Research Systems) was associated to a single photon module with high efficiency in the visible spectrum (SPCM-AQRH-13-FC from Excelitas Technologies') allowing studies in the nanosecond to millisecond range with a minimal resolution of 5 ns. In this latter case a picosecond laser photodiode (from Alphalas) delivering pulses of 60 ps at =407 nm was used. In order to improve collection efficiency spectral selection was achieved using highly selective dichroïc filters from Semrock company. All measurement were carried out in a confocal geometry in extremely low power regimes of excitation where non linear effects were not observed.

Figure 1 :

 1 Figure 1: (a.) RT absorption and (b.) photoluminescence spectra of solution of CdSe/CdSe1-xTex core/crown NPLs with 0 ≤ x ≤ 1. The emission was excited at 420 nm. The absorption and luminescence background have been shifted for clarity. On figure 1(a.), the blue line (green line) corresponds to the evolution of the wavelength of the maximum of the heavy hole exciton absorption coming to the crown (core) as a function of the Te content (c.) TEM image of the CdSe/CdSe0.3Te0.7 core/crown NPLs

Figure 2 :

 2 Figure 2: (a.) RT absorption spectra of 9 aliquots taken during the synthesis of CdSe/CdSe0.4Te0.6 core/crown NPLs. The first grey spectrum is the NPLs before injection. Between each aliquots, the surface of the NPLs increase by about the surface of the core NPLs (see supporting information). The spectra have been normalized at 420 nm. (b.) RT photoluminescence spectra of the corresponding aliquots taken during the synthesis of CdSe/CdSe0.4Te0.6 core/crown. The exciton wavelength was 420 nm. The absorption and luminescence background lines have been shifted upward for clarity Inset: blue emission contribution versus injected volume. (c.), (d.), (e.) and (f.) emission and scheme of the four types of emission depending on the crown in the core/crown heterostruture.

Figure 4 :

 4 Figure 4: (a.) PL signal for a solution of CdSe/CdSe1-xTex with x=0.6 at early times (t<1.5ns) after the light. For comparison, the steady signal with the two emissions peaks is plotted in dark red. (b.) PL decay at early times (t<1.5ns) for CdSe/CdSe1-xTex for three different x ratios of Te (x=0.5, 0.6, 0.8). A narrow band pass filtered is used to collect photon around 590nm (mainly crown emission). (c.) PL decay at long time scale for CdSe/CdSe1-xTex for three different x ratio of Te (x=0.5, 0.6, 0.8). The same filter as for (b.) is used. (d.) PL decay at long time scale for CdSe/CdSe1-xTex for two different x ratio of Te (x=0.6, 0.8) (mainly associated to the interface emission). For x=0.5, the core/crown NPLs don't exhibit emission from the interface. A narrow band pass filtered is used to collect photon around 630nm. In all cases, measurements are acquired at room temperature and the excitation wavelength is 407nm.

  (a)).

Figure :

 : Figure : (a.) Band gap of the CdSe1-xTex crown as a function of x between 0 and 1. The experimental points are fitted with the equation of the gap with a bowing parameter of 0.78. On the right side, the valence band ΔVv and conduction band ΔVc offsets between the CdSe core and the CdSe1-xTex crown and the Stokes shift between the absorption and the emission of the crown are given as a function of x. The blue zone corresponds to 250meV, a typical value of the exciton binding energy in NPLs. When the ΔVc< Eexc the two emission are observable. (b.) Representation of the relative band gap of the CdSe core with respect to a CdTe and CdSe0.4Te0.6 crown.

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ACKNOWLEDEGMENTS

We thank Agence Nationale de la Recherche for funding through grant Nanodose and H2DH. This work has been supported by the Region Ile-de-France in the framework of DIM Nano-K. This work was supported by French state funds managed by the ANR within the Investissements d'Avenir programme under reference ANR-11-IDEX-0004-02, and more specifically within the framework of the Cluster of Excellence MATISSE.