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Influence of Interfacial Strain on Optical Properties of PbSe/PbS Colloidal Quantum Dots Anna Rubin-Brusilovski, Youngjin Jang, Arthur Shapiro, Aron Safran, Aldona Sashchiuk, and Efrat Lifshitz Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.6b04098 • Publication Date (Web): 21 Nov 2016 Downloaded from http://pubs.acs.org on December 2, 2016

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Chemistry of Materials

Influence of Interfacial Strain on Optical Properties of PbSe/PbS Colloidal Quantum Dots Anna Rubin-Brusilovski, Youngjin Jang, Arthur Shapiro, Aron Safran, Aldona Sashchiuk, and Efrat Lifshitz* Schulich Faculty of Chemistry, Solid State Institute, Russell Berrie Nanotechnology Institute, Nancy and Stephen Grand Technion Energy Program, Technion–Israel Institute of Technology, Haifa 3200003, Israel

ABSTRACT: The interface in PbSe/PbS core/shell colloidal quantum dots (CQDs) is subject to strain forces due to a 3% crystallographic mismatch between the constituents. The strain profile in PbSe/PbS CQDs was simulated using the classical linear elasticity model, under the assumption of spherical-symmetric dot and isotropic materials. The derived strain profile was incorporated into a band structure calculation to evaluate the influence on the electronic bandedges of the core/shell CQDs. The electronic energy states evaluated were in close agreement with the absorption-edges of various core/shell CQDs with different core diameters and shell thicknesses. Furthermore, the synthesized CQDs underwent thermal annealing at various temperatures, thereby creating the alloying interface; consequently, their absorption and photoluminescence spectra exhibited spectral red-shift compared with the untreated samples. The band gap energy red-shift was simulated by the theoretical model, including smoothing potential at the interface. Measurements of the photoluminescence decays indicated an extension of the 1 ACS Paragon Plus Environment


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radiative lifetime after a controlled annealing process, denoting removal of defect quenchers around the core-shell interface. Thus, the study suggests practical means for mitigating interface strain to leverage the quality of core/shell structures.

Introduction: IV-VI colloidal quantum dots (CQDs) have been of great scientific interest during the past two decades, due to their tunable optical properties over a wide spectral range in the IR spectral regime, and also due to the relatively small effective masses of the carriers and high dielectric constant. These characteristics make them good candidates for various optoelectronic applications1, 2 such as solar cells3 and photosensors,4 and also useful for biological tagging.5 However, most IV-VI CQDs (e.g., PbSe) show high sensitivity to oxidation in ambient conditions,6 a process that might degrade their optical properties and make device fabrication challenging. The oxidation is related to the existence of surface dangling bonds or vacancies, which can be mitigated by using various kinds of surface passivation treatments. Various surface coverage alternatives had been proposed through the years, including coating by organic or inorganic molecular ligands7, 8 as well as atomistic ions,9 each may bound selectively to a certain facet (e.g., PbSe(100)).10 Alternatively, inorganic epitaxial coverage of internal cores by another semiconductor, in the form of core/shell heterostructures, was developed, for example, PbSe/PbS CQDs11-14 or PbSe/CdSe CQDs.15-18 The core/shell CQDs have chemical stability and higher photoluminescence (PL) quantum yield (QY) compared to corresponding cores, and they deliver improved performance as optical switches19 and solar cells.20 The physical properties of core/shell structures are controlled by the mutual alignment of the core and shell energy gap band-edges with respect to the vacuum, changing from wrapping of the core by the shell edges 2 ACS Paragon Plus Environment

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(called type-I) to a staggered arrangement (called type-II). The crystallographic match between the core and the shell is important in avoiding interface defects. A PbS shell exhibits crystallographic parameters close to those of a PbSe core, (6.12 Å and 5.94 Å for bulk PbSe and PbS, respectively),14 but even a small deviation of ~ 3% is a point of concern. Generally, shell growth is pseudomorphic, leading to a strain at the core-shell interface, due to a crystallographic mismatch between the constituents. The core experiences compressive strain in all directions, while the shell experiences compressive strain in the radial direction, and a tensile strain in the direction tangential to the interface when considering a spherical dot.21 When there is a significant lattice mismatch, a core-shell interface results in substantial lattice strain energy, affecting the shell morphology.22 When the shell is deposited at a low temperature, high strain energy causes the CQDs to move into a metastable phase, compared to alternative creation of a surface defect. Beside the metastability, the core/shell structure still maintains high PL QYs since the exterior surfaces are less pruned for oxidation.23 Fortunately, annealing of core/shell heterostructures partially or fully relieves the compressive strain in the core, with consequent pronunciation of a spectral energy red-shift of optical transitions.23 Mild increase of the shell thickness leads to relaxation and net reduction of the tensile strain within the shell. However, on formation of a thick shell, the continuity across the interface often involves creation of dislocations at the interface, a process that also relaxes some of the strain.24 Core-shell interfacial strain influences the bulk energy band offset25, 26 and the energy band structure of the CQDs.27 Core/shell CQDs with type-I behavior can be converted into type-II nanostructures under strain influence, leading to spatial separation of electrons and holes, extending the excited state lifetimes, all accompanied by a spectral shift.25 The energy bandgap temperature coefficient

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(dE/dT) is strongly influenced by strain. The negative contribution to dE/dT from the core-shell interfacial strain becomes important in the relatively larger core/shell CQDs.12, 28 Interfacial alloying has been shown to be a good method for decreasing interfacial strain with immediate influence on the optical properties of the CQDs,21, 26 mitigating interfacial defects and accordingly, increasing the optical QY. Alloying was implemented to generate high quality PbSe/PbS CQDs more than a decade ago, and at a later stage, was also implemented in the PbSe/CdSe CQDs.12, 29 It has been shown theoretically and experimentally that alloying reduces blinking and Auger recombination.30 Strain was pronounced in the Raman spectra of the core/shell compounds as a resonance shift, and the alloying composition appears as a new feature in the spectra, in accordance with some theoretical predictions.21, 29 While IV-VI core/shell CQDs have been studied extensively, there has been little research on the effect of core-shell interfacial strain on the optical properties nor much research about the influence of alloying on the strain relaxation. This work discusses the effect of the core-shell interfacial strain on the electronic and optical properties in PbSe/PbS CQDs with a 3% lattice mismatch. The strain profile in PbSe/PbS CQDs is calculated using the classical linear elasticity model,31 under the assumption of isotropic materials. The discussion is limited to a relatively thin shell with compressive/tensile interface distortion, excluding a threshold case in which a release of a strain occurs via formation of interface dislocations.24 The formation of interface dislocations could occur with involvement of a thick shell, when the formation of dislocation creates an interfacial carrier trapping site, and a reduction of the PL QY.24 The influence of strain on the energy band structure32 was investigated by incorporating the strain profile into a k●p method,33-35 to ascertain a realistic picture for the core/shell energy levels in the vicinity of the L point of a Brillouin zone. The results of the calculations were 4 ACS Paragon Plus Environment

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compared to the experimentally measured absorption of the core/shell CQDs before and after annealing. The results discussed below show that strain relaxation was pronounced as an energy red-shift of the absorption edge and as a reduction of decay rate in respect to the non-annealed samples. Presumably, the annealing process has relaxed the interfacial strain by an ionic interdiffusion and formation of graded alloying composition.

Results and Discussion: The core-shell interface strain in PbSe/PbS CQDs was calculated numerically using the stationary linear elasticity model31, 36 in the weak limit, as given in eq. (1): ∂σ ij ∂x j

+ Fi = 0 (1)

where σ ij is the stress tensor, and Fi are the body forces. The interfacial strain was calculated by using the Solid Mechanics module in Comsol, assuming a spherical shape of the CQDs and isotropy of the constituent materials. Important to note that the performed CQDs often show deviation from a perfect spherical shape with characteristic faceting,37 that could have an influence on the global strain forces. However, investigation of ensemble of CQDs with random orientation of the individual dots blurs the influence of anisotropic faceting, and therefore the symmetric shape modelling proposed here should provide a close agreement to the experimental observations (see below). Also worth noting that the piezoelectricity effect is neglected in the current study. This effect can be emanated either from intrinsic or extrinsic properties:38,

39

intrinsic is related to shape anisotropy such as wurtzite crystal structure; extrinsic effect is due to faceting or surface polarization. The PbSe/PbS core/shell CQDs possess rock-salt crystallographic structure with an inversion of symmetry, they show shape uniformity way 5 ACS Paragon Plus Environment


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beyond that of the corresponding cores11, 12 and the exterior surface polarization is further away from the core constituent, thus the contribution of an external piezoelectric effect can be neglected. The material parameters40, 41 used in the simulation for PbSe and PbS - the lattice constant, Young modulus and Poisson’s Ratio (the ratio of transverse to axial strain) are presented in Table 1.

Table 1. Material parameters used for strain simulations Lattice constant [nm] Young modulus [GPa] Poisson’s ratio PbSe 0.612 66 0.4 PbS 0.593 75 0.24

A plot of the calculated strain profile along the radial axis is shown in Figure 1a. The black line represents the strain along the radial direction, and the red line is the strain in the direction tangential to the core-shell interface. Positive strain values indicate tensile strain, while negative values indicate compressive strain. A schematic representation of the forces created at the cross section of the CQD along the equatorial due to the interfacial strain is shown in Figure 1b. The core is compressively strained in all directions (blue arrows); the shell is stretched in the direction tangential to the core-shell interface (red arrows) and compressively strained in the radial direction (black arrows). The x direction represents the radial coordinate in Figure 1a.


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(a)

(b)

Figure 1. (a) Plot of the calculated strain in radial (black line) and tangential (red line) directions, relatively to x axis for a PbSe/PbS CQD, with core radius of 1.85 nm and total radius of 2.7 nm. Negative (positive) strain values indicate compressive (tensile) strain. (b) Schematic drawing of the strain distribution across the equatorial of a CQD. The core is compressed in all directions (blue arrows), while the shell is compressed in the radial direction (black arrows) but stretched in the direction tangential to the interface (red arrows).

The following discussion evaluates the influence of an interfacial strain in PbSe/PbS core/shell CQDs on the electronic and optical properties. The energy band structure of the PbSe/PbS CQDs studied was calculated using the k●p isotropic four band envelope function method33, 34 for PbSe and PbS,35 with the Hamiltonian read as (eq. (2)):  h 2 kQr − k  Ec + 2   0   H0 =  h  Pk z m0    h  m P ( k x + ik y )  0

0

h Pk z m0

h 2 kQr − k Ec + 2

h P ( k x + ik y ) m0

h P ( k x − ik y ) m0

Ev −

−

h Pk z m0

h 2 kQr + k 2 0

 h P ( k x − ik y )  m0   h − Pk z  m0    0   h 2 kQr + k  Ev −  2 

(2)

Ec and Ev are the energies of the bulk conduction and valence band, respectively. k = −i∇ , ki (i=x,y,z) are the partial derivatives ki = −i ∂ ∂i , P is the momentum matrix element, m0 is the free electron mass and Q-r(Q+r) are radial functions describing the behavior of the electron (hole) effective mass. The variation of the effective masses from core to shell are described in the form of a step function, while the step sharpness gets smoother after annealing due to inter-diffusion of Se and S between core and shell, which leads to alloyed composition that minimizes 7 ACS Paragon Plus Environment


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crystallographic mismatch. The zero energy at the bottom of the conduction band was arbitrarily chosen. Strain modifies the crystal periodicity, causing lattice points to move from their equilibrium positions. The small effect on a single lattice point accumulates to a pronounced effect when considering the entire dot with hundreds of atoms. Therefore, static homogeneous strain cannot be considered as a perturbation. As crystal atoms move to their new location under a strain, the electronic potential of the crystal changes, and it is expressed in the nonrelativistic strained Hamiltronian32 as in eq. (3):

H (ε ) =

p2 + Vε 2m

(3)

where Vε is the strained crystal potential, m is the effective mass of a charge carrier, p is the momentum and ε is the strain, which is related to a stress by ε ij = 1 / 2 ( ∂ui / ∂j + ∂u j / ∂i ) . A transformation of coordinates given by eq. (4) describes a lattice point (x') in the strained crystal in terms of a lattice point (x) in the unstrained crystal. x = (1 + ε ) x '; p = (1 − ε ) p '; p 'i = −ih

∂ ∂x 'i

(4)

p ≈ p ' − 2∑ p 'i ε ij p ' j 2

2

ij

The momentum p is transformed, as described in eq. (4), as well as the derivatives and squares. Under this transformation, one can use the same Bloch functions and boundary conditions used for the unstrained crystal. As a result, the strained potential Vε(x)=Vε[(1+ε)x'] and the unstrained potential V0(x) have the same periodicity, when the deviation between them is defined by eq. (5):

Vε (1 + ε ) x ' − V0 ( x ') = ∑ Vij ( x ')ε ij

(5)

ij


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To determine the wavefunctions of the electron and hole in the strained crystal, we begin from ψεk' which is the eigenfunction of H(ε), corresponding to eigenvalue E(ε, k') and wavevector k'. Under the transformation as in eq. (4) and k = (1 + ε )k ' , we get a modified wavefunction ψk', with the same periodicity as the wavefunction ψ0, that corresponds to the unstrained Hamiltonian H0, and the same k vector, leading to eq. (6) :

ψ ε k ' = uε k ' ( x )eik ' x → ψ 'k = u 'k [(1 + ε ) x ']eikx (6) Calculating ψ 'k

H '(ε ) ψ 'k

gives a system of equations (eq. (7)):

( Eα ∑  α

k0

)

− E δαα ' + H 'α 'α cα = 0 

(7)

where E is the energy level, α and α’ are matrix indexes. The new Hamiltonian H' is defined by (eq. (8) and (9)): H ' = H 0 + Hε

(8)

where

∑ pε

i ij

Hε = −

ij

m

pj

+ ∑Vij ( x ')ε ij

(9)

ij

Eq. (9) is due to the ∇ → (1 − ε )∇ transformation and H' has the same periodicity as the unstrained Hamiltonian H0. Neglecting inter-band terms H'αβ, in the nonrelativistic approximation, the matrix elements Hβ'β includes the expression as in eq. (10) and (11): H β0 ' β =

h2k 2 h h2 δ ββ ' + kpβ ' β + 2 2m m m

∑γ

( kpβ 'γ )( kpβγ ) E β − Eγ

(10)



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 (p p )  H βε ' β = ∑  − i j β ' β + Vβij' β  ε ij = ∑ Dβij ' β ε ij m ij  ij 

(11)

where Dβij ' β the material’s deformation potential given by (eq. (12)): Dβij ' β = −

( pi p j ) β ' β m

+ Vβij' β

(12)

Perturbation theory is used to find the change in band energy caused by the strain, using the relation (eq. (13)):

δ E (ε ) = E (ε , (1 − ε )k ) − E0 (k ) (13) where E(ε,(1-ε)k) is the electronic energy level in the strained crystal at point k'. For a Hamiltonian of the form shown by eq. (8), containing a small correction given by eq. (13), first order correction on the perturbed bandgap energy42 is given by eq. (14):     E gε =  Ec + ∑ Dijcε ij  −  Ev + ∑ Dijvε ij  ij ij    

(14)

where Egε is the bandgap energy of a strained crystal, and Dc,v are the deformation potentials of conduction (c) and valance (v) bands. Eq. (14) was implemented into the energy calculations described above. The deformation potentials were taken from reported literature values,42 as presented in Table 2.

Table 2. Deformation potentials used for the calculations42 Decore = 1.06; Deshell = 1.38;

Dhcore = 3.14; Dhshell = 3.42;


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Inserting eq. (14) into eq. (2) enables calculating the electron and hole energy levels for a strained bulk and CQD, when the strain profile is taken from a numerical simulation. As said above, centro-symmetric structures were assumed in the present case. Thus, a change of volume is isotropic in all directions. Annealing was found to be a good method to release interfacial strain by atomic diffusion and formation of alloyed composition at the interface of PbSe/PbS CQDs. Se and S atoms diffuse between core and a shell, creating a smooth defect free interface, minimizing crystallographic mismatch and interfacial strain. The diffusion lengths for Se and S are in Table 3 for annealing at various temperatures for 15 minutes. D0 is the pre-exponential factor used for the diffusion length calculations.43, 44

Table 3. Pre-exponential factors and diffusion lengths of Se and S- ions after annealing during 15 minutes at various temperatures Annealing Temperature

D0 for Se

D0 for S

Diffusion Diffusion length for length for S [nm] Se [nm]

[nm2/s]

[nm2/s]

80 0C

5.99×10-6

1.86×10-8

0.07

0.004

100 0C

4.97×10-5

1.59×10-7

0.21

0.01

110 0C

1.31×10-4

4.29×10-7

0.34

0.02

120 0C

3.32×10-4

1.09×10-6

0.54

0.03

130 0C

8.01×10-4

2.68×10-6

0.84

0.05

150 0C

0.0041

1.41×10-5

1.92

0.11



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The diffusion length of Se is similar to a typical thickness of the shell nearly at all temperatures; however, it is consistently larger by an order of magnitude from the diffusion length of the S- ion. Hence, the short diffusion length of S insures the alloyed layer remains in proximity to the core-shell interface and a complete alloying is avoided under the annealing time duration and temperatures chosen for this experiment. At higher annealing temperatures, as thermal energy increases, the longer diffusion lengths of Se may induce a crystalline disorder that would be pronounced as a quenching of the luminescence process (see below). Figure 2 shows the PbSe and PbS bulk energy band offset (black lines) and first energy levels (blue lines) of a PbSe/PbS CQD with a representative core radius of 1.95 nm and overall radius of 2.3 nm with interfacial strain between the core and the shell (a), or after annealing treatment (b).

(a)

(b)

Figure 2. Calculated PbSe and PbS bulk band offset (black) and energy levels (blue) for PbSe/PbS CQDs with representative core radius of 1.95 nm and overall radius of 2.3 nm (the red line represents the boundary between the core and the shell), for (a) as synthesized CQDs with strain; (b) after annealing process. The green dotted lines represent the bulk energy band offset without strain or annealing.


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The red dotted line shows the limit between core and shell, and the green dotted line shows the bulk energy band offset without strain or annealing. The band offset profile and energy levels for the CQDs with strain were calculated using the simulated strain profile (Figure 1) and the corrections to the unstrained Hamiltonian were described above. The band offset profile and energy levels for the annealed CQDs were calculated using the Hamiltonian in eq. (2). A smooth error function (erf(r)) was incorporated into the potential V and the effective mass function Qr, in order to describe the gradual change of composition along the core-shell interface after annealing. As seen in Figure 2a, a strain bends the bulk energy band offset of the shell, leading to a larger energy offset of the valance band near the interface, with reduced influence away from the interface. But alloying is expressed as smooth transitions from core to shell band offset, as seen in Figure 2b. The curvature of the error function at Figure 2b is suited to simulated alloyed region thickness generated at 110 °C, based on the variables given in Table 3. Figure 2 reveals an increase of energy by 12 meV in the valence band edge of annealed CQDs with respect to the valence band edge of strained CQDs.

The theoretical predictions elaborated so far were examined experimentally using various core/shell CQDs with a fixed core radius and variable shell thickness. The samples were synthesized at a low temperature of 80 °C, and then annealed at higher temperatures between 80 °

C–150 °C (see methods). Low temperature synthesis enables strain formation at the core-shell

interface, while the strain relaxation occurs during the annealing process, where there is sufficient energy for ionic diffusion. Representative experimental results are shown in Figure 3.



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Figure 3. (a) TEM images of PbSe/PbS with core radius of 1.95 nm and an overall radius of 2.3 nm. The inset includes TEM image of the same CQDs after annealing at 110 °C. (b) Absorption (dashed) and emission (solid) lines of the same CQDs presented in (a) before annealing (black) and after annealing for 15 minutes at 80 °C (red), 100 °C (green), 110 °C (blue), 120 °C (cyan), 130 °C (magenta) and 150 °C (orange). (c) PL and absorption peak position as a function of annealing temperature of the samples in (b). (d) Full width at half maxima (FWHM) of the PL peaks as a function of annealing temperature of the samples in (b).

Figure 3 shows TEM images of PbSe/PbS CQDs before (a) and after annealing at 110 °C (inset of (a)). Both particles have core radius of 1.95 nm and an overall radius of 2.3 nm, indicating that the annealing process does not affect the size or shape of the particles. Figure 3b shows the absorption and continuous-wave (CW) PL spectra of the same CQDs as in (a), dissolved in hexane and drop casted on glass, before (black) and after annealing at six different temperatures (using the color code indicated in the figure caption). The red-shift in the absorption and PL spectra shown in Figure 3b is probably due to the atomic diffusion at the core-shell interface. There is a clear energy red-shift of the absorption and emission after annealing with respect to the core/shell before annealing, as shown in Figure 3c. This red-shift is associated with the strain release and diffusion at the core-shell interface. The annealing process at 110 °C for 15 minutes led to an energy red-shift of 14 meV, which is in close agreement with the theoretical value of 12 14 ACS Paragon Plus Environment

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meV, as shown in Figure 2. Figure 3d shows a plot of full width at half maxima (FWHM) of PL peaks of CQDs as in panel (b) as a function of annealing temperature. The FWHM shows only a mild change up to 110 °C, and grows dramatically at higher annealing temperatures. The change of trend in the FWHM above 110 °C may be related to the deterioration of the crystal quality and also may be related to an intensive dominant inward diffusion of Se atoms.

Figure 4. (a) Plot of the PL intensity (log scale) versus time after the excitation pulse, for the samples shown in Figure 3, with the following color code: before annealing (black) and after annealing for 15 minutes at 80 °C (red), 110 °C (blue), 120 °C (green) and 130 °C (violet). Inset: Plot of the decay rate values (extracted from the decay curves) before (red) and after annealing (black) versus the annealing temperature. (b) Plots of the band gap energy red-shift Eg, before (b.a) and after (a.a) annealing at 110 °C for 15 minutes, versus the shell width. The plots compare the energy shift of experimental first exciton absorption band (solid symbols) with the calculated values (hollowed symbols) for two different core radii (as indicated in the legend). 15 ACS Paragon Plus Environment


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The transient PL measurements of the CQDs in film are plotted as a function of time in Figure 4a before (black) and after annealing at different temperatures (using a color code indicated at the figure caption). The decay rate is plotted versus the annealing temperature in the inset of Figure 4a. It should be note that decay measurement was achieved by nearly resonance excitation 13, 45, 46

by generating electron-hole pair across the band-gap which avoid occupation of higher

electronic states. So, loss of population by carrier trapping shows direct relation between the PL QY and the lifetime, showing different trend of II-VI CQDs.47 For samples annealed at temperatures between 80 °C to 110 °C, the decay rate is reduced compared to that found before the annealing process, presumably associated with the improvement of interface quality. However, at higher annealing temperatures, a gradual increase of the decay rate takes place. It is likely that an annealing process at elevated temperatures leads to a loss of ligating at the exterior surface with consequence agglomeration or formation of crystal disorder, globally inducing nonradiative routes for the photo-generated carriers or excitons. Thus, the broadening of the exciton emission band (shown in Figure 3) and the partial quenching of the luminescence (Figure 4) could be correlated with the same adverse deterioration of the CQDs quality, when vigorous thermally treatment was applied. Indeed, TEM image of the CQDs annealed at 150 °C presented certain percent of small aggregates.48 This outcome should lead to the conclusion (see below) that the annealing process is extremely effective, if it is restricted to temperatures that efficiently recover the core-shell interface, without exceeding a threshold limit. According to Vegard’s law, it is anticipated that full body annealing of the CQDs would lead to a blue-shift in the absorption and emission. Contrary to the classical Vegard’s law, an energy red-shift of absorption and emission spectra has been observed upon annealing at and below the


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optimal temperature of 110 °C, supporting the creation of partial annealing at a restricted regime around the core-shell interface, as predicted by the theoretical calculations. The experimental and calculated band-edge energy red-shifts obtained after annealing, compared to those for stained core/shell CQDs, are plotted in Figure 4b, versus the shell thickness, for two different CQDs with core sizes as indicated in the legend. The unfilled symbols correspond to the calculated values, and the filled symbols correspond to the experimental observations (energy of the first exciton absorption band). The figure shows that a red-shift increases substantially up to a certain thickness with the increase of the shell width, and then the change becomes more gradual. In any event, the influence on a small core is larger, due to the larger fractional contribution of the shell. It was previously shown that thicker shells impose stronger strain on CQDs, to the point at which the strain is released by dislocation formation.24 This is consistent with the trend observed in Figure 4b. To conclude, in this work, the absorption red-shift caused by core-shell interfacial strain relaxation in PbSe/PbS CQDs was calculated and measured. Annealing is a good method to relax interfacial strain by introducing alloying at the interface, without changing the CQDs’ shape or size. Relatively low annealing temperatures (110 °C) are sufficient for strain relaxation and improvement of the optical properties, while higher annealing temperatures lead to degradation of the optical properties due to crystal defect formation.

Methods Materials. Lead(II) oxide (PbO; 99.999%), selenium (99.99%), sulfur (99.99%), bis(trimethylsilyl) sulfide (TMS2S; synthesis grade), n-hexadecane (HDC; 99%), 1-octadecene (ODE; tech. grade, 90%), and oleic acid (OA; 90%) were purchased from Sigma-Aldrich. 17 ACS Paragon Plus Environment


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Ammonium

chloride

(A.C.S.

grade)

was

purchased

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from

Spectrum

Chemicals.

Trioctylphosphine (TOP; 97%) and diphenylphosphine (DPP; 99%) were purchased from Strem. Tetrachloroethylene (TCE; spectroscopic grade) was purchased from Merck. Toluene (analytical), hexane (analytical), and ethanol (absolute) were purchased from Bio-Lab Ltd. Acetone (absolute) was purchased from Gadot. These chemicals were used without further purification.

Synthesis. The PbSe and PbSe/PbS CQDs were synthesized by modifying the procedure developed by Yanover et al49 and Shapiro et al.50 For the synthesis of PbSe CQDs, 4 mmol of PbO, 12 mmol of OA and 10 ml of HDC were mixed in a three-neck flask and the mixture was heated to 100 °C under vacuum for 1 hour to form Pb(OA)2. Then the mixture solution was heated to 90 °C under nitrogen. Next, the selenium precursor solution containing 2 ml of 2 M TOPSe and 4 mmol DPP was rapidly injected into the Pb(OA)2 solution under nitrogen. The temperature was reduced to 80 °C and the reaction mixture was maintained at the same temperature until the required CQDs size was reached. Afterward, the reaction mixture cooled down and the CQDs were isolated by centrifuging using ethanol, repeated twice. Finally, the precipitate was dried and then re-dissolved in hexane for further use. The amount of the S precursors was calculated by considering the concentration of PbSe CQDs and the desired PbS shell layer. The required mass of Pb precursors was calculated using 8:1 of Pb:S molar ratio. In a typical synthesis, PbO, OA (PbO:OA=1:8) and ODE (5.84 g) were mixed in a three-neck flask. The solution was heated to 100 °C under vacuum for 1 hour and then cooled down to room temperature under nitrogen. A hexane solution of 3.7×10-7 mol PbSe core CQDs was injected into the Pb oleate solution. The solution was degassed until hexane was evaporated. Afterward, the reaction mixture was heated to 80 °C and the diluted TMS2S solution 18 ACS Paragon Plus Environment

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prepared by mixing 0.1 mL of TMS2S and 4 mL of ODE was injected dropwise into the reaction vessel using a syringe pump (0.1 mL/min). After attaining the desired shell thickness, the reaction was cooled down to room temperature. The purification of the product was carried out by adding ethanol and centrifuging, repeated twice. The precipitate was re-dissolved in hexane for further use and characterization.

Annealing. The annealing of PbSe/PbS CQDs was carried out under inert atmosphere. A certain amount of PbSe/PbS CQD solution in hexane was dissolved in HDC after removing hexane. The solution was heated to 80 °C and maintained for 15 minutes. The temperature of the solution was increased with increments of 10 °C to 150 °C and the solution was maintained at each annealing temperatures for 15 minutes. After annealing at each temperature, aliquots were taken for optical measurement. The CQD solution was washed by adding ethanol and dissolved in hexane.

Characterization. Transmission electron microscopy (TEM) was conducted using a FEI Titan operated at 300 keV. The absorption spectra of the CQDs were recorded by using a JASCO V570 UV-VIS-NIR spectrometer. CW-PL and transient PL decay measurements were performed at room temperature. The CQDs (in a film on a glass substrate) were excited either by a 532 nm diode laser or Coherent 890 nm laser or by a pulsed Nd:YAG laser (see also reference 45). The emission was detected by an Acton Spectrapro 2300i monochromator equipped with a photomultiplier tube (Hamamatsu NIR-PMT H10330-75) operating in the NIR spectral region.



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ASSOCIATED CONTENT AUTHOR INFORMATION

Corresponding Author *E-mail: [email protected]

Notes The authors declare no competing financial interests.

ACKNOWLEDGMENTS The work was supported by the Israel Council for Higher Education-Focal Area Technology (Project No. 872967), the Volkswagen Stiftung (Project No.88116), the Israel Ministry of Defense (Project No. 4440665406), the Israel Ministry of Trade (Maymad Project No. 54662), the Israel Science Foundation Bikura (Project No. 1508/14), the Israel Science Foundation (Project No. 985/11 and 914/15), the Niedersachsen-Deutsche Technion Gesellschaft E.V (Project No. ZN2916), and supported by the European Commission via the Marie-Sklodowska Curie action Phonsi (H2020-MSCA-ITN-642656)


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TOC graphic:


PbS

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