Boosting the Photoluminescence Efficiency of InAs Nanocrystals Synthesized with Aminoarsine via a ZnSe Thick‐Shell Overgrowth

InAs‐based nanocrystals can enable restriction of hazardous substances (RoHS) compliant optoelectronic devices, but their photoluminescence efficiency needs improvement. We report an optimized synthesis of InAs@ZnSe core@shell nanocrystals allowing to tune the ZnSe shell thickness up to seven mono‐layers (ML) and to boost the emission, reaching a quantum yield of ≈70% at ≈900 nm. It is demonstrated that a high quantum yield can be attained when the shell thickness is at least ≈3ML. Notably, the photoluminescence lifetimeshows only a minor variation as a function of shell thickness, whereas the Auger recombination time (a limiting aspect in technological applications when fast) slows down from 11 to 38 ps when increasing the shell thickness from 1.5 to 7MLs. Chemical and structural analyses evidence that InAs@ZnSe nanocrystals do not exhibit any strain at the core‐shell interface, likely due to the formation of an InZnSe interlayer. This is supported by atomistic modeling, which indicates the interlayer as being composed of In, Zn, Se and cation vacancies, alike to the In2ZnSe4 crystal structure. The simulations reveal an electronic structure consistent with that of type‐I heterostructures, in which localized trap states can be passivated by a thick shell (>3ML) and excitons are confined in the core.


Introduction
Colloidal near infrared (NIR) emitting semiconductor nanocrystals (NCs) are being developed for next-generation optoelectronic devices, including photovoltaics, [1][2][3] light-emitting diodes (LEDs), [4] optical communication, [5,6] biological imaging, [7,8] night vision, [9] and lasers. [10,11]Aside from the well-developed, but toxic Hg-based (II-VI) and Pb-based (IV-VI) NCs, [12][13][14] the most promising NIR emissive NCs are those based on (III-V) InAs. [8,15,16]][26] To implement InAs NCs in optoelectronic devices, their optical properties require further optimization, not only in terms of their absorption peak position and linewidth, [24,25,27,28] but also of their photoluminescence (PL) quantum yield (QY) and Auger recombination rate. [26,29,30]In fact, as-synthesized colloidal InAs NCs typically exhibit a very weak PL emission (PLQY ≈1%), limited by surface defects, such as surface In vacancies, that act as non-radiative recombination centers for photoexcited carriers. [26,31,32]Boosting the PL efficiency of InAs NCs requires coating them with a wide bandgap semiconductor material that can form a type-I InAs@shell heterostructure.[34][35] The optimal choice for a shell material would be CdSe, given its negligible lattice mismatch with InAs, [33,36,37] and indeed InAs@CdSe core@shell based NCs can achieve high PLQY values.For example, the PLQY values of InAs@CdSe@CdS NCs with PL peak at 970 or 1425 nm were reported to be as high as 82% or 18%, respectively, [8] while that of InAs@CdSe@ZnSe NCs with PL peak at 1065 nm was 52%. [38]voiding the use of toxic elements, such as Cd, is of paramount importance to enable the application of core-shell NCs based on InAs in commercial devices.This constraint limits the choice of shell materials to basically ZnSe and ZnS.However, the presence of a relatively high lattice mismatch between these zinc chalcogenides and InAs (6.44% for ZnSe and 10.7% for ZnS) [33] results in a significant strain at the core/shell interface of InAs@ZnSe(S) core@shell NCs.][41] To address this issue, InAs@multishell NC systems have been synthesized by growing "buffer" shell layers between the InAs core and the ZnSe shell.These layers are made of InP, which has a relatively low lattice mismatch with InAs (3.13%) and/or GaP, further reducing the overall strain at the core/shell interface. [33,42]Examples of such systems include InAs@InP@ZnSe NCs emitting at 980 or ≈830 nm with high PLQY values of 76% and ≈50%, respectively, [43,44] InAs@InP@GaP@ZnSe NCs with PL emission at ≈1107 nm and PLQY of 23%, [42] and InAs@InP@ZnSe@ZnS with PL emission at 873 nm and PLQY of 25%. [45][48][49][50][51] These reasons have prompted various groups in recent years to develop synthesis approaches to InAs NCs and the corresponding core@shell structures based on alternative, less toxic and cheaper As precursors, with the most promising one being tris(dimethylamino)-arsine (amino-As). [27,28,39,41,52,53]The current challenge is not only to improve the control over the size and size distribution of InAs NCs made with amino-As, but also to prepare Cd-free core@shell heterostructures based on such NCs and having improved optical properties, such as high PL emission efficiency and suppressed Auger recombination rate, for the realization of efficient LEDs and lasers.
We recently demonstrated an effective amino-As and ZnCl 2based synthesis approach to grow InAs@ZnSe NCs with a record PLQY value of 42% (with PL emission peaking at ≈860 nm), which was ascribed to the formation of an In-Zn-Se interlayer following our synthesis conditions. [32]In this work, we have refined this reaction scheme to such a level of control that we could deliberately tune the thickness of the ZnSe shell from ≈1 to 7 mono layers (ML).Upon increasing the shell thickness, we observed the following trends: i) the PLQY of the core@shell NCs first increased from ≈40%-50% (1.5ML) up to ≈70% (3.5ML) and then reached a plateau from 3.5 to 7ML, setting a new record value for this system; ii) the PL lifetime (average  of 58 ns) did not vary as a function of the shell thickness; iii) the Auger recombination slowed down as a function of the shell thickness, with the associated lifetime increasing from 11 ps (1.5ML) to 38 ps (7ML).High-resolution scanning transmission electron microscopy (STEM) analysis indicated that no strain is present at the interface of our core@shell NCs (i.e., neither compressive nor tensile strain was observed in the core and shell regions, respectively).Moreover, X-ray photoelectron spectroscopy (XPS) evidenced the presence of indium atoms coordinated with Zn and Se in the whole shell region, with a much higher concentration in the first layer.Overall, these observations strongly suggest the presence of an In-Zn-Se interlayer between the InAs core and the ZnSe shell.
Our atomistic model, built by matching the experimental elemental ratios, NCs size and featuring an electronic structure compatible with the observed optical properties suggested that: i) our core@shell NCs have an interlayer composed of In, Zn, Se, and cation vacancies, alike to In 2 ZnSe 4 ; ii) such interlayer plays a key role in dampening the strain between the core and the shell; iii) the core-shell band alignment is compatible with that of a type-I heterostructure, although, given the elaborate interlayer structure and composition, the core and shell states are intermixed; iv) the core@shell system with two shell layers (i.e., one In-Zn-Se ML and one ZnSe ML) is characterized by an incomplete surface traps passivation, while a more effective passivation can be obtained by increasing the overall shell thickness to >3 ML.
These results are noteworthy not only for the record PLQY values achieved on such InAs@ZnSe core@shell structures, but also because it was previously believed that a thick ZnSe shell cannot be sustained on InAs NC cores without degradation of the optical properties.6][57] 2. Preparation and Characterization of InAs@ZnSe Core@Shell Nanocrystals InAs@ZnSe core@shell NCs with tunable shell thickness were prepared by following our recently reported synthesis scheme in which we systematically varied the shell growth parameters. [32]In detail, InAs core NCs were prepared by employing InCl 3 , amino-As, ZnCl 2 , oleylamine (OA), and DMEA-AlH 3 at 300°C with a fixed InCl 3 :amino-As:ZnCl 2 precursors ratio of 1:1:20 (see Experimental part for details).The ZnSe shell was grown in situ by adding TOP-Se and ZnCl 2 (dissolved in OA) at 90°C and heating the system to the desired temperature (Figure 1a).Optimization experiments were performed by fixing the size of the starting InAs NCs (namely ≈3 nm, Figure 1b, with absorption peaked at 853 nm, Figure 2 a) and systematically varying: i) the shell growth temperature (from 280 to 330°C); ii) the amount of TOP-Se (from a Se:In ratio of 5:1 to 37.5:1); iii) the amount of added ZnCl 2 (with a resulting Zn:In ratio ranging from 20:1 to 30:1); iv) the reaction time (up to 300 min) (see Section S1, Supporting Information for details).The best synthesis procedure in terms of shell thickness tunability and optical properties (which will be discussed later on) consisted of adding 7.5 ml of 1 M TOP-Se (total Se:In ratio of 37.5:1) and 2.5 ml of 0.8 Mm ZnCl 2 -OA (total Zn:In ratio of 30:1) solutions to the crude InAs core NCs reaction mixture at 90°C and performing the shell growth at 310°C.These conditions allowed the regulation of the ZnSe thickness from 1.5ML (0 min) to 7ML (300 min).The shell thickness in the systems presented in this manuscript was estimated by correlating the results from elemental analyses, performed by using Inductively Coupled Plasma Optical Emission spectroscopy (ICP-OES), and transmission electron microscopy (TEM), with a structural model consisting of a 3 nm tetrahedral InAs core surrounded by a shell with a variable thickness (see Section S2, Supporting Information for further details).
The low-resolution TEM images (Figure 1b-g) show that the overall size of the InAs@ZnSe NCs systematically increased together with the thickness of the ZnSe shell, without any apparent presence of newly homo-nucleated (small) ZnSe NCs.According to our measurements, the size of the NCs increased from 3 nm (InAs core) to 4.2 nm (1.5ML), 5.8 nm (3.5ML), 7.2 nm (5ML), 8.3 nm (6.5ML), and 9.2 nm (7ML) (Figure 1h and Table 1), consistently with our structural model (Table S2, Supporting Information).All samples had XRD patterns compatible with a cubic zinc-blende phase, with no indication of secondary phases (Figure 1i).The XRD pattern of the starting InAs core NCs matched that of bulk InAs (ICSD 98-002-4518), while the position of the XRD peaks gradually shifted toward higher two-theta angles with increasing ZnSe shell thickness (i.e., toward the position expected for bulk ZnSe, Figure 1i; Figure S6, Supporting Information).
As for what concerns the optical properties, the position of the exciton absorption peak of InAs NCs, initially located at ≈853 nm, featured a small red-shift (reaching ≈864 nm) upon the overgrowth of the first ZnSe ML, and then blue shifted when the shell thickness was increased from 3.5ML (≈841 nm) to 7ML (≈827 nm) (Figure 2a,b and Table 2).Similarly, the PL peak position of the InAs@ZnSe core@shell NCs was observed to red-Table 1.Shell thickness, composition of InAs@ZnSe core@shell samples measured via ICP-OES and corresponding size measured via TEM analyses.shift from ≈929 nm (InAs core) to ≈947 nm (InAs@1.5ML-ZnSe)and then to blue-shift (≈900 nm for InAs@7ML-ZnSe) (Figure 2a,b and Table 2).A possible explanation of this behavior is that the first monolayer, with composition and structure close to In 2 ZnSe 4 (see "First Principles Simulations" section below), on the one hand is able to passivate surface traps, explaining the increasing in PL, but on the other hand is not able to confine the carriers efficiently, thus explaining the PL red shift.With the growth of a thicker shell, the overall composition and structure of this first monolayer can be modified, hence its ability to better confine the carriers in the core can improve, and this would explain the PL blue-shift.Notably, the PLQY of the samples obtained with our optimized method increased from ≈1% (InAs cores), to ≈40%-50% after the overgrowth of the first shell layer and then up to ≈70% with 3.5ML of ZnSe (Figure 2c), setting a new record value for these systems.This was achieved only when employing large Zn:In and Se:In precursors ratios (>30:1), while instead working with lower Zn:In (20:1) and Se:In (5:1 or 10:1) ratios led to PLQY values similar to those reported by us in our previous work (≈40%-50%, Figures S1 and S2, Supporting Information).Another remarkable finding was that the increase in the ZnSe shell thickness did not lead to any drop in PLQY, which remained stable at ≈70% from 3.5 up to 7ML (Figure 2c).Time-resolved photoluminescence (TRPL) measurements (Figure 2d) indicated that core@shell NCs with a ZnSe shell thickness of 1.5ML featured a fast  1 component (Table S5, Supporting Information).This could be attributed either to an incomplete surface passivation of the InAs core or to the fact that a thin shell might only be able to induce a partial carrier confinement (which would explain also the redshift of the absorption and PL spectra observed when going from InAs core to the 1.5ML system, Figure 2a,b), similar to what reported for other III-V core@shell NC systems, such as InP@ZnSe. [58]Additionally, we only observed a lengthening of the PL lifetime for shell thicknesses >1.5 ML, which is in line with the increase in PLQY and can be ascribed to the improved surface passivation of thicker shells.From 3.5ML onward, no significant changes in the average PL lifetime were observed (Table 2), suggesting a type-I band alignment between InAs and ZnSe. [59]e then evaluated the exciton dynamics of these systems by measuring the Auger recombination rate as a function of the ZnSe shell thickness.The Auger rates were measured using transient-absorption spectroscopy (TAS) at a low fluence limit, close to the onset of biexciton formation, with an average exciton occupancy of <N> ≈ 0.4, assuming that the photogenerated exciton occupancy distribution follows a Poisson statistic.Our analysis (Figure 2e and Table 2) revealed a slight suppression of the Auger rate with varying the shell thickness, from 11 ps (1.5ML) to 38 ps (7ML) in line with the reported values in similar core@shell systems. [29,60]Since the InAs core and ZnSe shell feature a type-I band alignment, both electrons and holes are confined in the InAs core.Therefore, increasing the shell thickness should only lead to a limited suppression of the Auger rate. [61]We hypothesize that the three-fold suppression of the Auger rate observed here could be attributed to an increase of the exciton volume due to the smoothed interface potential in our type-I InAs@ZnSe heterostructures.In fact, it has been already reported that graded shell strategies and annealing procedures could soften the confinement potential, decreasing the rate of nonradiative Auger recombination. [29,61,62]

Chemical and Structural Characterization
][65] In InAs@InP@ZnSe core@shell@shell NCs, where InP was added to relieve the strain between the InAs core and the ZnSe outer shell, a similar trend of PLQY decline after the growth of two ZnSe ML was also observed. [43,44]Our results suggest that, under our reaction conditions, the ZnSe shell may grow while maintaining a reduced strain at the core@shell interface or by releasing such strain without defect formation.
The structure of our core@shell NCs was then investigated via an extensive X-ray photoelectron spectroscopy (XPS) analysis.The XPS data, consistent with ICP-OES data, revealed a sharp increase in the In/As ratio after the growth of 1.5 ML of ZnSe, moving from 1.2 to 1.44, suggesting that a non-negligible amount of In was incorporated in the shell.As the number of shell layers increased, XPS detected a further increase in the In/As ratio, reaching ≈3.1 for the 6.5 and 7 ML cases, which is much higher than the ratios estimated by ICP-OES (Table 1).Considering the XPS surface sensitivity, the comparison between the two analyses suggested that, during shell growth, the excess In was preferentially located in the outer part of the core@shell heterostructures (i.e., incorporated in the ZnSe shell).This was also reflected in the Zn/As and Zn/In ratios changes as a function of the shell thickness (Table 3 and Figure S7, Supporting Information).
The composition and structure of our core@shell structures were further investigated via high resolution XPS analysis of the In 3d 5/2 peak (Figure 3a).In the InAs core NCs, we observed a single In signal centered at (444.4 ± 0.2) eV, which is typical of InAs. [66]In contrast, the In peak of the 1.5ML InAs@ZnSe NC sample could be fitted with two components.One component was ascribed to InAs, and the other, centered at (444.9 ± 0.2) eV, was close in energy to that of In 2 Se 3 , which we previously associated with In coordinated to both Zn and Se.This suggests the formation of an In-Zn-Se interface layer. [32]This component increased in relative intensity when analyzing the 3.5-7ML samples (Figure 3a) indicating that In cations incorporated in the ZnSe region are distributed across the whole shell, in agreement with ICP elemental analyses that indicate an increase in the In/As ratio for thicker shells (Table 1).
We also excluded that the increase in the In content of the thicker shell samples emerging from both ICP and XPS analyses was due to an inefficient cleaning of the samples or to In acting as a Z-type ligand (in the form of InCl 3 for example).To do so, we performed a ligand stripping procedure based on Meerwein salts on the 6.5ML InAs@ZnSe NC sample and we performed both ICP and XPS analyses on the treated sample (see Table S6, Supporting Information). [67]Such procedure is commonly employed to replace original surfactants (e.g., organic molecules and/or Ztype ligands, such as ZnCl 2 bound to the surface of InAs NCs) with BF 4 − species. [32,67]Our results did not evidence any change in the In/As composition.This further supported the hypothesis that In was included in the whole ZnSe shell, since the removal of In as external impurity or Z-type ligand would have resulted in a lower In/As ratio in the stripped sample compared to the initial one.
To gain more insights into the local structure of our core@shell NCs, we conducted high-resolution (HR) STEM characterization of NC samples with a thick ZnSe shell.We first acquired compositional maps via STEM energy-dispersive Xray spectroscopy (EDX) on individual core@shell NCs and confirmed that they comprised a small (≈3 nm) InAs core and a thick (≈4 nm) ZnSe shell, resulting in a total NC size of ≈11 nm (Figure 3b).Quantification carried out by integration on the entirety of the NC yielded elemental ratios compatible with those acquired via XPS and ICP measurements (Figure S8a, Supporting Information).On the other hand, due to the limited EDX signalto-noise ratio that can be achieved on these NCs in realistic operating conditions, the quantification of low concentrations of In cations in the shell (as emerged from ICP and XPS analyses) was very challenging (Figure S8b,c, Supporting Information).
To investigate the local atomic structure of individual core@shell NCs, we performed a HRSTEM characterization.Figure 3c shows a NC imaged along the [111] zone axis with the corresponding fast Fourier transform (FFT) pattern displayed in the inset.The NC had a ≈3 nm core, a ≈4 nm thick shell and a total diameter of ≈11.5 nm.In the HRSTEM image it is possible to count ≈20 (20 2) ZnSe planes (belonging to the set of {220} planes of ZnSe, see Figure 3c) corresponding to a shell thickness of 10ML.Such estimate is in agreement with our atomistic model in which a size of 11.3 nm is expected for a 10ML core@shell NC (see Table S2, Supporting Information).Moreover, the pyrami-dal shape of the NC projected at a low-indexed zone axis orientation allowed us to visualize strain fields using Geometric Phase Analysis (GPA).Figure 3d and Figure S8d (Supporting Information) show the local variation in the lattice parameter within a single core@shell NC.The intensity of each point corresponds to the local relative variation of the interplanar distance for the {220} planes compared to the average value computed the overall average spacing used as reference.This map indicates that the lattice parameter is constant throughout the whole ZnSe shell (2.017 ± 0.015Å) and sharply increases in the InAs core region.The NC core had a larger (≈6%) lattice parameter (2.138 ± 0.015 Å) than that of the shell, which was consistent with the expected bulk values of the two materials.Our measurements  suggest that elastic strain is negligible in both the shell and core regions.
Based on the results of XPS, compositional and HRSTEM analyses, it appears that an In-Zn-Se interlayer is formed between the InAs core and the ZnSe shell in our core@shell systems. [32]This interlayer can act as a buffer (or spring) that "absorbs" the lattice strain, thereby explaining the absence of the expected strain resulting from the lattice mismatch between ZnSe and InAs.Additionally, the In-Zn-Se interlayer may contribute to the reduction in the Auger recombination rate by smoothing the interface potential and more effectively delocalizing the carriers in the ZnSe shell.

First Principles Simulations
We performed density functional theory (DFT) simulations to reveal the atomistic structure of InAs@ZnSe NCs and to investigate the effect of ZnSe coating layers on the PLQY.Extensive tests on atomistic models were performed to select, through a trial-and-error approach (see Section S7, Supporting Information for details), the NC model that displayed an electronic structure compatible with the observed optical properties while best matching the experimental size and elemental ratios.The models were limited, for computational reasons, to a 3 nm InAs core and 2ML shell thickness.The best model, with stoichiometry In 462 As 373 @Zn 722 Se 798 Cl 115 , had a smoothed tetrahedral shape, exposing the (111), (100), and (1̅ 1̅ 1) facets (Figure 4a ).Interestingly, the first shell layer displayed an In-Zn-Se composition and a structure similar to that of In 2 ZnSe 4 (Figure 4a).In 2 ZnSe 4 is the only stable compound comprising In, Zn and Se ions and crystalizes in a zinc blende-like structure with cation positions alternatively occupied by Zn, In, or by cation vacancies. [68,69]Its lattice parameters are intermediate between those of InAs and of ZnSe.This latter characteristic is believed to help in reducing the core-shell strain, preventing the formation of strain-induced defects, and could potentially explain the observed high PLQY values.Finally, the outermost ZnSe layer features cation vacancies, a type of reconstruction typically observed in the surfaces of zinc-blende semiconductors [70] that plays a role in reproducing the correct band gap for the ZnSe shell (Figure S9, Supporting Information).
In Figure 4b,c, we report the density of states (DOS) and representative molecular orbitals (MOs) of InAs@ZnSe NCs: i) the highest occupied (HO) and lowest unoccupied (LU) MOs are localized in the core and can thus be seen as the valence and conduction bands of InAs, respectively; ii) the shell contribution becomes relevant in the HOMO-5 and LUMO+1 MOs, hence they can be assigned to the valence and conduction bands of ZnSe, respectively.In the MOs farther from the band gap (e.g., HOMO-264), the shell contribution increases, although the core contribution remains significant.Overall, the alignment of the InAs core and ZnSe shell bands is in line with that expected for a type-I heterostructure. [71]However, our simulations reveal a more complex band alignment scenario than the one depicted by the energy diagrams typically employed to describe these systems, [38,71,72] in which the band gap of InAs (0.35 eV) lies roughly in the middle of that of ZnSe (2.82 eV). [73]In our NC model, the conduction bands minima of the core and the shell are well separated from each other in the DOS, and they lie 0.36 eV apart (LUMO and LUMO+1 in Figure 4c, respectively), while the valence bands of the core and the shell form a continuum in the DOS (see for example how HOMO, HOMO-1, and HOMO-5 are all close in energy).The shell contribution starts at only 87 meV below that of the core.We note that the adopted DFT functional underestimates the band gap, but it is expected to correctly reproduce the bands alignment (see Experimental Section).
Finally, our simulations provided a rationale for the increase in PLQY observed as the shell thickness increases (in line with the standard explanation for the PLQY increase generally observed upon shell growth): [71] i) the localization of both the valence and conduction band edge states within the InAs core effectively confines the exciton in the core region; ii) the presence of one In-Zn-Se ML and one ZnSe ML leads to an efficient, albeit not complete, surface traps passivation.In our 2ML model, indeed, we still observe a small contribution of localized, trap-like surface states near the valence band maximum (e.g., "HOMO-5" in Figure 4b) which originate from structural distortions on the surface.In detail, these trap states form in correspondence of points/sites where surface vacancies (in the ZnSe ML) are in contact with In, Zn and vacancies of the underlying In-Zn-Se layer.We speculate that when additional ZnSe layers are present, the surface vacancies (of the outer ZnSe layer) would not be any more in contact with the In-Zn-Se interlayer and these structural distortions (and thus the trap states) would not form.This would explain the PLQY increase measured when going from 1.5 to 3.5 ML systems (Figure 2c).

Conclusion
Our synthesis scheme for InAs@ZnSe core@shell NCs based on amino-As, DMEA-AlH 3 and ZnCl 2 was optimized to finely tune the ZnSe shell thickness up to 7 ML and boost the PLQY to a record value of 70%.We discovered that the shell thickness did not influence the PL lifetime of the core@shell NCs, but had an effect in reducing the Auger rates from 11 to 38 ps when it was increased from 1.5 to 7ML.Interestingly, the PLQY increased from 40%-50% to 70% when going from 1.5 to 3.5ML and then it remained constant up to 7ML.The core@shell NCs prepared with our synthesis approach do not exhibit strain at the interface, as revealed by HRSTEM analysis.This was ascribed to the formation of an In-Zn-Se interlayer, as revealed by our elemental and XPS analyses.We also proposed a sound model for the atomistic and electronic structures of InAs@ZnSe NCs: i) the interlayer, which plays a key role in dampening the core-shell strain, was composed by In, Zn, Se, and cation vacancies, alike to In 2 ZnSe 4 ; ii) the core-shell bands alignment was compatible with that of type-I heterostructures, but it was shown to differ qualitatively from the commonly adopted models derived from macroscopic interfaces.Overall, the core@shell NCs presented here represent a promising starting point for future full RoHS compliant InAsbased optoelectronic devices.
Preparation of the As Precursor: In a N 2 filled glovebox, 0.2 mmol of amino-As was dissolved in 0.5 mL of degassed oleylamine at 40 °C for 5 min until no bubbles further evolved.
Preparation of 1 M TOP-Se Precursor: In a N 2 filled glovebox, 10 mmol of Se powder was mixed with 10 mL of TOP in a 20 mL glass vial and heated at 250°C under constant stirring for ≈30 min to form a transparent solution, and then the mixture was cooled down to room temperature.
Preparation of 0.8 M ZnCl 2 -OA Precursor: In a N 2 filled glovebox, 8 mmol ZnCl 2 was mixed with 10 mL of OA in a 20 mL glass vial and heated at 250°C under constant stirring for ≈50 min.Because 0.8 M ZnCl 2 -OA Precursor solidified at room temperature, so, it had to be preheated before transferring into a syringe.
Synthesis of InAs Core NCs: InAs core NCs were synthesized following a former work of ours with minor modifications. [32]In a typical synthesis, 0.2 mmol of InCl 3 , 4 mmol of ZnCl 2 , and 5 mL of OA were loaded into a 100 mL three-necked flask under an inert atmosphere.The mixture was degassed at room temperature for 10 min and then at 120°C under vacuum for 40 min.Next, the flask was heated up to 180°C under N 2 to completely dissolve all the precursors, and then it was cooled down to 120°C and dried under vacuum for extra 30 min.The mixture was heated to 240°C under nitrogen, and the As precursor was injected into the flask, quickly followed by the injection of 1.2 mL of the DMEA-AlH 3 toluene solution.The temperature was quickly increased to 300°C (≈30°C min −1 ), the reaction was then allowed to run for 15 min and it was quenched by removing the heating mantle.The NCs were washed twice by the addition of toluene and ethanol and precipitated by centrifugation at 4000 rpm.The final product was dispersed in toluene for further characterization.
Synthesis of InAs@ZnSe Core@Shell NCs: After quenching the growth of the InAs NCs by cooling the reaction mixture to 90°C, 2.5 ml of 0.8 M ZnCl 2 -OA was injected into the flask followed by the injection of 7.5 ml TOP-Se.The mixture was heated up to 310°C (≈30°C min −1 ).Samples of different shell thickness were obtained by taking aliquots at different reaction times (after reaching 310°C), namely: 0 min (1.5ML), 5 min (3.5ML), 15 min (5ML), 180 min (6.5ML), and 300 min (7ML).The NCs were washed by the addition of toluene and ethanol and precipitated by centrifugation at 2000 rpm for two times.The final product was dispersed in toluene for further characterization.
Powder X-Ray Diffraction (XRD): XRD patterns were carried out on a PANanalytical Empyrean X-ray diffractometer, equipped with a 1.8 kW Cu K ceramic X-ray tube and a PIXcel3D 2 × 2 area detector, operating at 45 kV and 40 mA.Specimens for XRD were prepared by drop-casting the concentrated sample solution onto a quartz zero-diffraction single crystal substrate in the glovebox.The diffraction patterns were collected under ambient conditions at room temperature using a parallel beam geometry and the symmetric reflection mode.XRD data analysis was conducted on the HighScore 4.1 software from PANanalytical.
Transmission Electron Microscopy (TEM) Characterization: Diluted composite dispersions were drop-cast onto copper TEM grids with an ultrathin carbon film.Low-resolution TEM images were acquired on a JEOL JEM-1400Plus microscope with a thermionic gun (W filament) operated at an acceleration voltage of 120 kV.High resolution scanning transmission electron microscopy (STEM) images were acquired on a probe-corrected ThermoFisher Spectra 300 S/TEM operated at 300 kV.Images were acquired on a High-Angle annular Dark Field (HAADF) detector with a current of 50 pA.Compositional maps were acquired using Velox, with a probe current of ≈150 pA and rapid rastered scanning of the beam.The Energy-Dispersive X-Ray (EDX) signal was collected by a Dual-X system, comprising two detectors, one on either side of the sample, for a total acquisition solid angle of 1.76 Sr. Geometric Phase Analysis (GPA) of the HAADF-HRSTEM images was performed using the commercial Digital Micrograph software package and homemade scripts. [74]-Ray Photoelectron Spectroscopy (XPS): Specimens for XPS were prepared from concentrated NC solutions, dropped on freshly cleaved highly oriented pyrolytic graphite substrates in a glovebox.XPS measurements were carried out on a Kratos Axis Ultra DLD spectrometer using a monochromatic Al K source, operated at 20 mA and 15 kV.High resolution analyses were carried out at a pass energy of 10 eV.The Kratos charge neutralizer system was used on all specimens.Spectra were chargecorrected to the main line of the carbon 1s spectrum (adventitious carbon) set to 284.8 eV.Spectra were analyzed using CasaXPS software (version 2.3.24).
Inductively Coupled Plasma (ICP): The elemental analysis was also performed via inductively coupled plasma optical emission spectroscopy (ICP-OES) with an iCAP 6300 DUO ICP-OES spectrometer (ThermoScientific).The samples were dissolved in 1 mL of HNO 3 overnight and then diluted with 9 mL of Milli-Q water for measurements.The elemental analysis using ICP-OES was affected by a systematic error of ≈5%.
Optical Properties: The absorption spectra were recorded on a Varian Cary 5000 UV−vis−NIR spectrophotometer.The samples were prepared by diluting NC samples in 3 mL of toluene in 1 cm path length quartz cuvettes with airtight screw caps in a N 2 filled glovebox.The steady-state and time-resolved PL measurements were carried out on a Edinburgh Instruments FLS900 fluorescence spectrometer equipped with a Xe lamp and a monochromator for steady-state PL excitation and a time-correlated single photon counting unit coupled with a Edinburgh Instruments EPL-510 pulsed laser diode ( ex = 508.2nm, pulse width = 177.0ps) for timeresolved PL.The PLQY measurements were performed using the Edinburgh Instruments FLS900 fluorescence spectrometer equipped with an integrating sphere, exciting at 700 nm using the output of Xe lamp.All NC solutions were diluted to an optical density ≈0.1 at the excitation wavelength.
TA Characterization: Transient absorption measurements were carried out using a mode-locked titanium sapphire based ultrafast amplifier centered at 800 nm and generating 45 fs pulses at a repetition rate of 1 kHz.The optical setup utilized was a typical pump-probe non-collinear configuration.The fundamental mode from the amplifier was directed into a half wave plate and a thin film polarizer system to control the energy of the excitation energy.An optical parametric amplifier pumped with ≈1 mJ of the fundamental 800 nm energy was used to generate the probe beam at 870 nm.The probe beam optical path included a precise motorized translation stage to control the optical delay between the pump and the probe beam.The probe beam was directed on the sample where changes in transmission and reflection were recorded simultaneously using lock-in amplifiers.
First Principles Simulations: All simulations were performed with a density functional theory (DFT) Hamiltonian through the cp2k code. [75,76]he Perdew-Burke-Ernzerhof (PBE) exchange-correlation functional was adopted. [77]All core electrons were included in the pseudopotential, [78] while the valence electrons were described with the DZVP basis set included in the cp2k package and with plane waves having an energy cutoff of 320 and 400 Ry for geometry optimization and electronic structure calculation, respectively.Although the adopted PBE functional was known to strongly underestimate the band gap, [79] it was shown that this functional provides a good description of band alignments. [80]PBE was thus ideal for the purpose, i.e., studying the atomistic structure and the interplay between core and shell electronic states on real-dimension nanocrystals (≈2500 atoms).The periodic simulation box measured 75 × 75 × 75 Å.The atomistic details of the simulated system were described in Section S7 (Supporting Information).The DOS plot of Figure 4b was obtained by smearing the discrete energy levels of the nanocrystals MOs with a Gaussian function having a width of 0.03 eV.

Figure 1 .
Figure 1.Structural and morphological characterization of InAs@ZnSe core@shell NCs with different values of shell thickness.a) Schematic representation of the synthesis of InAs@ZnSe core@shell NCs.TEM images of b) InAs core and core@shell NCs with c) 1.5ML, d) 3.5L, e) 5ML, f) 6.5ML, or g) 7ML ZnSe shell thickness.h) Size distribution histograms calculated from the corresponding TEM images.The vertical dashed bars represent the expected size calculated in our structural/atomistic model.i) XRD patterns of InAs and InAs@ZnSe NCs together with the bulk reflections of InAs (ICSD 98-002-4518) and ZnSe (ICSD 98-007-7092).

Figure 2 .
Figure 2. Optical characterization of InAs core and InAs@ZnSe core@shell NCs.a) Optical absorption and PL spectra; b) position of the exciton absorption and PL peaks; c) PLQY values; d) PL decay traces; e) TA curves in which the carrier dynamics are governed by Auger recombination and biexciton.

Figure 3 .
Figure 3. XPS spectra of InAs and InAs@ZnSe NCs and HR STEM characterization of individual 5ML InAs@ZnSe NCs.a) XPS spectral decomposition of the In 3d 5/2 signal, collected on the core and core@shell samples.b) HR STEM-EDX elemental maps; c) atomic resolution HAADF-STEM image of a core@shell NC oriented along the [111] zone axis and d) the corresponding GPA analysis in which the intensity of each point corresponds to the local value of the interplanar distance for the {220} planes.

Figure 4 .
Figure 4. Atomistic and electronic structure of a NC model composed of a 3 nm InAs core and 2ML shell.a) Atomistic structure of the whole core@shell NC (left), the core and the In-Zn-Se interlayer (middle) and the core (right).b) DOS of the core@shell NC. c) Energy levels of MOs around the Fermi energy.The InAs and ZnSe contributions to each energy level are shown by red and green colors, respectively.Representative MOs are shown as blue/red ± 0.005 isosurfaces, with grey arrows indicating the corresponding energy in the plot.When cartesian axes are not shown, the orientation is the one displayed on the top of panel (c).The red (green) arrow highlights the assignment of the core(shell) band gap.

Table 3 .
XPS quantitative analysis of InAs and InAs@ZnSe NC samples.