Nanocrystals into Tunable and Efficient Light Emitters by Aliovalent Allo

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Turning Weakly Luminescent SnO Nanocrystals into Tunable and Efficient Light Emitters by Aliovalent Alloying Vahid Ghodsi, and Pavle V. Radovanovic Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.8b01574 • Publication Date (Web): 20 Apr 2018 Downloaded from http://pubs.acs.org on April 20, 2018

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Chemistry of Materials

Turning Weakly Luminescent SnO2 Nanocrystals into Tunable and Efficient Light Emitters by Aliovalent Alloying Vahid Ghodsi and Pavle V. Radovanovic* Department of Chemistry, University of Waterloo, 200 University Avenue West, Waterloo, Ontario N2L 3G1, Canada

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ABSTRACT

We report the synthesis of ternary gallium tin oxide nanocrystals, and demonstrate that their photoluminescence can be tuned through the visible region by changing Ga:Sn ratio. By substitutional doping with Ga3+, the PL intensity of SnO2 nanocrystals is enhanced by nearly three orders of magnitude, reaching photoluminescence quantum yield of > 40 %. Increase in PL intensity is attributed to the formation of donor and acceptor pairs, and the increase in emission energy is discussed in the context of band gap expansion and stronger Coulomb interaction between charged defect sites. Time-resolved and steady-state photoluminescence spectroscopies reveal that the interaction of extrinsic and native defects is driven by the nature of the dopant ion. By adjusting various reaction conditions, we prepared the nanocrystals with nearly ideal scotopic-to-photopic ratio and a quantum yield of ca. 34 %, attesting to the potential of these nanocrystals for general lighting applications. The results of this work provide new insight into the role of defect chemistry in tailoring the optoelectronic properties of transparent metal oxide nanocrystals, and pave the way for the rational design of light sources and photonic devices with high photoluminescence efficiency, minimum toxicity, and optimal lighting characteristics.


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INTRODUCTION Ever since the discovery of the potential of semiconductor nanocrystals (NCs) as photoluminescent materials, they have been a focus of an intense research to develop new energy-efficient and cost-effective sources of lighting.1-9 II-VI NCs have attracted particular interest as a tunable light source, owing to their size dependent electronic structure and optical properties.10,11 Much effort has consequently been directed toward maximizing the photoluminescence quantum yield (PLQY) or increasing the emission color tunability.8,12-14 However, CdSe NCs, as the most widely used nanophosphor in light emitting diodes (LEDs), may suffer from instability and toxicity. Transparent metal oxide (TMO) NCs have more recently been explored as an alternative due to their strong defect-based emission, chemical stability, environmentally benign nature, and transparency to visible light.15-18 In addition, functionally-relevant point defects, size and phase dependent properties, and structural versatility have motivated researchers to study these materials to establish deeper understanding of their fundamental properties and expand their functionality.15,19-27 Ga2O3 has the widest bandgap among TMOs (4.9 eV) and crystallizes in five different polymorphs (α, β, γ, δ, and ε phases). Optoelectronic properties of the monoclinic gallium oxide (β phase), as the stable polymorph at room temperature, have been most extensively studied.28-31 We have recently shown that γ-Ga2O3 can also be stabilized in the nanocrystalline form through the colloidal synthesis, and its optical properties can be altered by tuning the size and composition of the NCs.15,16,32 γ-Ga2O3 has a defective spinel structure, with two Ga3+ vacancies per each 18 gallium sites. Oxygen vacancies ( VO• ), which are readily formed in Ga2O3, act as electron donors, while a complex consisting of gallium-oxygen vacancy pair ( (VO , VGa )′ ) likely gives rise to the acceptor state.28,30 Upon exciton formation, the conduction band (CB) electron is


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trapped in the donor state, while the hole in the valence band (VB) is trapped in the acceptor state. The photoluminescence (PL) properties arise from the electron-hole recombination upon transfer of the electron from the donor to the acceptor site (tunnel transfer). This donor-acceptor pair (DAP) mechanism is a hallmark of Ga2O3, and the resulting PL can be manipulated by controlling defect interactions.32,33 We subsequently showed that luminescent dopants or conjugated molecular adsorbates can interact with native defects via energy transfer, and could be utilized to manipulate the PL properties of these NCs to generate white light emission.5,34-37 Another TMO that has found various applications ranging from catalysis and sensing to solar cells, spintronics, and lithium ion batteries is tin(IV) oxide.38-41 SnO2 is an oxygen-deficient ntype semiconductor with the rutile crystal structure and the band gap of 3.6 eV. The presence of oxygen vacancies is the origin of its high conductivity and visible emission.42 Majority of the studies on this material have centered around its remarkably high carrier concentration (1020 cm3

), while PL properties remain underexplored relative to Ga2O3. Visible emission of SnO2 NCs is

generally thought to involve surface defect states.38,43,44 Although quantum yield has rarely been reported, relative increase in the PL efficiency for SnO2 NCs has been achieved by doping44 while spectral modulation has been achieved by changing the coordinating ligands or solvent.38 A recent study reports a quantum yield of ca. 60 % for deep blue emission of SnO2 NCs synthesized by thermal decomposition of Sn(IV) precursor in a high boiling point solvent.45 The emission of SnO2 nanostructures originates from the transition of an electron trapped in a shallow donor state formed by oxygen vacancies to native surface states within the band gap.46,47 The depth of the donor and the surface states were estimated to be about 0.03-0.15 eV below the conduction band minimum and 0.9 eV above the valence band maximum, respectively.46,48,49 Despite having a direct band gap, the band-to-band transition in SnO2 is dipole-forbidden which


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hampers its functionality in optoelectronic devices. Nevertheless, doping and reducing dimensionality of tin oxide could lead to breaking of the selection rule and recovering its optical activity.47,50 Doping and alloying have proven to be efficient methods for altering the electronic structure of semiconductor

NCs,

thereby

introducing

new

magnetic,

optical,

and

electrical

properties.16,22,23,26,41,51-54 Rational incorporation of aliovalent impurities into a host lattice generates extrinsic defects with energy levels usually located within the band gap. Substitutional doping of external cation with higher oxidation state than the host lattice cation introduces energy levels below the conduction band minimum acting as localized donor states. On the contrary, substitution with a cation possessing lower oxidation state than the host lattice cation results in introducing the acceptor states above the valence band maximum.53 To compensate for the charge imbalance, other point defects, such as oxygen vacancies, may concurrently form in the lattice.55 As an example, the incorporation of Ga3+ in SnO2 and Sn4+ in Ga2O3 can be depicted by eqs 1 and 2, respectively, using Kröger-Vink notation:

SnO 2 Ga 2 O3  → 2Ga ′Sn +3O×O +2VO• (1) Ga 2 O3 SnO 2  → Sn •Ga + 2O×O +e′

(2)

The negatively and positively charged states, created as a result of doping, attract owing to the Coulomb interaction, forming donor-acceptor pairs. The ability to induce such extrinsic defects and control their concentration opens up the possibilities for tailoring the functional properties of materials at nanoscale.


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Motivated by these opportunities and considering the fact that octahedrally-coordinated Ga3+ and Sn4+ have similar ionic radii (Ga3+ 0.62 Å and Sn4+ 0.69 Å), several research groups have attempted to grow Sn4+-doped Ga2O3 or Ga3+-doped SnO2 in order to enhance the charge carrier concentration and mobility, or tune the band gap and optical properties of the respective host lattices.53,56-59 Since the adopted synthetic methods were largely based on solid state reactions, the solubilities of Ga in SnO2 and Sn in Ga2O3 have been relatively low (up to ca. 5 mol %).60-62 However, in certain nanostructures the achievable doping concentrations were reported to be significantly higher (> 20 %).63-65 Herein, we report for the first time the hydrothermal synthesis of ternary gallium tin oxide (Sn1-xGaxO2-0.5x, 0 ≤ x ≤ 1), further denoted by GTO NCs. Doping of Ga2O3 NCs with Sn4+ or SnO2 NCs with Ga3+, together with the presence of appropriate native defects, leads to the formation of donors and acceptors in both NC lattices, allowing for a modulation of the PL throughout the visible part of the spectrum. The effects of the reaction duration, synthesis temperature, and NC composition on the PLQY and lifetime of the NCs are discussed. We also show how the interaction among intrinsic and dopant-induced defects can lead to phosphors with high PLQY (> 40 %) even in SnO2 NCs with intrinsically low PLQY (ca. 0.5 %). The ability to simultaneously control the emission efficiency and fine-tune the PL of these broadly-emitting NCs allows for the design of new energy-efficient lighting sources with optimal photopic-to-scotopic ratio and other lighting characteristics.

EXPERIMENTAL SECTION Materials. All reagents and solvents are commercially available, and were used as received. Gallium(III) nitrate hydrate (Ga(NO3)3·xH2O, 99.99 %), and tin(IV) chloride pentahydrate (SnCl4·5H2O, 98 %) were purchased from STREM Chemicals. Oleylamine (OAm, 70 %), oleic


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acid (OA, 90 %), trioctylphosphine oxide (TOPO, 90 %), ethanol (EtOH, HPLC grade) and hexane (HPLC grade) were purchased from Sigma-Aldrich. Methanol (MeOH, 99.9 %) was purchased from Fisher Scientific. Synthesis and Processing of GTO NCs. In a typical synthesis of SnO2 NCs, 1 g of SnCl4·5H2O, 7 mL MeOH, 20 mL OA, and 4 mL OAm were placed into a 45 mL teflon-lined stainless-steel autoclave. The autoclave was transferred to a preheated oven (190 oC). The reaction duration varied from 4 to 32 hours, after which the autoclave was left to cool down to room temperature naturally. The obtained xerogel was washed with ethanol and centrifuged three times, followed by TOPO treatment at 90 oC for 30 minutes.66 The obtained NCs were then washed with ethanol, centrifuged, and finally dispersed in hexane for optical measurements. For the preparation of alloyed GTO NCs with variable composition (including pure Ga2O3 NCs), a portion of tin chloride was replaced with Ga(NO3)3·xH2O. This nominal Ga content was used throughout the manuscript to distinguish the samples, unless stated otherwise. The synthesis and post-synthesis treatment were performed following the procedure described above. For the preparation of white-light emitting nanoconjugate, a suspension of GTO NCs in hexane with the absorbance of 0.5 at 240 nm was prepared. 4 mL of this suspension was added to 1 mL aqueous solution of Atto-590 (5.9 µM), forming a bilayer. The bilayer emulsion was gently stirred and the non-polar (hexane) portion was extracted every 10 minutes to monitor the emission color under UV lamp and collect a PL spectrum. After 40 minutes of stirring and allowing the dye to bind to the surface of GTO NCs, the emission color turned white. Characterization and Measurements. The NCs were structurally characterized by powder Xray diffraction (XRD), and transmission electron microscopy (TEM). The electronic structure


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and optical properties of NCs were studied by optical absorption, and steady-state and timeresolved PL measurements. XRD patterns were collected with an INEL powder diffractometer featuring a positionsensitive detector and monochromatic Cu Kα radiation (λ = 1.5418 Å). JOEL 2010F microscope operating at 200 kV was utilized for TEM imaging and energy dispersive X-ray (EDX) spectroscopy. TEM samples were prepared by dropping a hexane suspension of NCs on a copper grid containing carbon support film (Ted Pella, Inc). Absorption

spectra

were

measured

at

room

temperature

with

Cary

5000

ultraviolet−visible−near-infrared (UV-vis-NIR) spectrophotometer. Steady-state PL spectra were recorded at room temperature with a Varian Cary Eclipse fluorescence spectrometer. The samples were excited at 240 nm with the excitation and emission slits set to 5 nm. Time-resolved PL measurements were performed with a Horiba Jobin-Yvon IBH Ltd. spectrometer using timecorrelated single photon counting (TCSPC) method. Ludox solution (Sigma Aldrich), was used to collect the instrument response function (IRF) and the data were fit with multi-exponential functions. The samples were excited using a 250 nm nano-LED (IBH Ltd.), and emission signal was recorded at a right angle with respect to the excitation source at the corresponding emission band maxima. A solution of quinine bisulfate (QBS) in 1N sulfuric acid and 1,4-bis(5-phenyloxazole-2yl)benzene (POPOP) in cyclohexane were used to obtain the relative PLQY of the NCs. The PLQY was calculated using the following expression:

QR AR EX nX2 QX = AX ER nR2

(3)


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where the subscripts R and X denote the quantities of the reference and unknown compound, respectively, Q is the photoluminescence quantum yield, A is the absorbance of the solution, E is the integrated fluorescence intensity, and n is the refractive index of the solution. The PLQYs of QBS and POPOP were reported to be 0.55 and 0.97 under these conditions. We reported the average values for the NC PLQY with respect to QBS and POPOP.

RESULTS AND DISCUSSION We first examined the structure and composition of GTO NCs synthesized under different conditions. XRD patterns of GTO NCs prepared with different starting concentrations of Ga3+ precursor (nominal Ga content) for the reaction duration of 8 h are presented in Figure 1a. The patterns of pure SnO2 and Ga2O3 NCs agree well with the pattern of bulk rutile tin oxide (JCPDS 088-0287, red sticks) and γ-phase gallium oxide (JCPDS 020-0426, blue sticks), respectively. With increasing content of substitutional Ga3+ in SnO2 NCs, the rutile XRD peaks shift to higher angles (Figure 1b), which is consistent with the smaller ionic radius of octahedrally coordinated Ga3+ (0.62 Å) compared to Sn4+ (0.69 Å). Until the molar Ga:Sn precursor ratio of 50:50, the XRD patterns reveal only the formation of rutile SnO2 NCs. Although a concurrent formation of small gallium-oxide-based clusters (either amorphous or crystalline) cannot be completely excluded, high resolution XRD patterns in Figure 1b suggest that these solvothermallysynthesized SnO2 NCs have a higher capacity to dissolve Ga relative to bulk SnO2, due to lower synthesis temperature which favors kinetic doping and prevents expulsion of the dopant impurities. This observation is consistent with the reports of rather high Ga3+ doping concentrations in nanostructured SnO2 synthesized under non-equilibrium conditions.63-65 The NC transformation to γ-Ga2O3 structure becomes evident when Ga:Sn starting precursor ratio


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exceeds 75:25. The broadening of the XRD peaks for undoped SnO2 and Ga2O3 NCs is associated with their small size. An increase in the concentration of dopant impurity (Ga3+ in SnO2 and Sn4+ in Ga2O3) causes an increase in the local disorder of the NC host lattice, leading to further increase in full width at half maximum (fwhm) for both NC host lattices. On the other hand, the XRD peaks become narrower by increasing the reaction duration (Figure 1c) or temperature (Figure 1d), for the same ratio of Ga:Sn (50:50), suggesting the improvement in the crystalline order of GTO NCs. It is also evident that the peak at ca. 35 degrees for rutile GTO NCs has the smallest fwhm, suggesting the anisotropic morphology of SnO2 NCs, and the longest dimension along {101} plane.38

Figure 1. (a) XRD patterns of GTO NCs synthesized with different starting Ga3+ concentration (in atom %), as indicated in the graph. (b) High-resolution XRD patterns of GTO NCs having nominal Ga content up to 50 % (as shown in the graph). Dashed red line designates the reflection for (101) plane of bulk rutile SnO2. (c) XRD patterns of GTO NCs with 50 % nominal Ga content, synthesized for different reaction duration (4, 8, 16, and 32 hours). (d) XRD patterns of


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GTO NCs with 50 % nominal Ga content, synthesized for 8 h at different temperatures, as shown the graph. Red and blue sticks in all panels represent the patterns of bulk rutile SnO2 and γGa2O3, respectively. TEM images of GTO NCs, synthesized with different Ga3+ precursor concentrations for 8 h reaction duration, are shown in Figure 2. The histograms of NC size distribution are shown in Figure S1 (Supporting Information), and the corresponding values of the average NC sizes are summarized in Table S1 (Supporting Information). At low Ga3+ starting concentration (≤ 25 %), the NCs have elongated morphology and high crystallinity, as evident in Figure 2a and b, respectively. Increasing Ga content induces some lattice disorder and adversely affects the crystallinity of the SnO2 NCs (Figure 2c). The measured average lattice spacing for {110} planes in GTO NCs for overall Ga content of ≤ 50 % is smaller than that for SnO2 NCs (3.35 Å) owing to the substitutional incorporation of Ga3+ in the NC lattice, and is consistent with XRD patterns in Figure 1b. The morphology of the NCs becomes spherical for high Ga content (90 %), and the measured lattice spacing reflects the transformation to γ-Ga2O3 crystal structure (Figure 2d). The actual Sn and Ga concentrations, determined by EDX spectroscopy, reveals that Sn tends to be incorporated in the lattice more readily than Ga for all starting precursor ratios (Table S2 in Supporting Information). The spatial distribution of Ga and Sn was assessed by EDX elemental mapping (Figure 2e-g), and is consistent with the coexistence of Ga and Sn in the NC samples (Figure 2f and g, respectively). To understand the origin of the elongated morphology of rutile GTO NCs, we examined the NC shape evolution with respect to the reaction time. Overview and high resolution TEM images of GTO NCs prepared for varied reaction times are shown in Figure S2 (Supporting Information). At shorter reaction times the isolated NCs are largely spherical. As the reaction continues, these NCs undergo oriented attachment, assembling into nanowires with


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orientation,38,41 consistent with narrowing of the fwhm of the corresponding XRD peak (Figure 1c). No significant change in the Ga content was observed at longer reaction times (Table S2 in Supporting Information). Performing the reaction at higher temperature also elongates the particles (Figure S3 in Supporting Information), although it is accompanied by the exclusion of Ga, as indicated in Table S2 (Supporting Information).

Figure 2. (a-d) TEM images of GTO NCs synthesized under the same conditions, but with different nominal Ga contents: (a, b) 10 % Ga, (c) 50 % Ga, and (d) 90 % Ga. The measured average lattice spacing corresponds to rutile SnO2 (b, c), and γ-Ga2O3 (d). (e-g) Scanning transmission electron microscopy (STEM) image (e), and the corresponding elemental maps for Ga (f) and Sn (g) of GTO NCs with nominal 50 % Ga content synthesized for 32 h.


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We have previously shown that colloidal γ-Ga2O3 NCs do not exhibit quantum confinement in the size range achievable by colloidal synthesis,15,33 rendering their band gap energy independent of the NC size. However, because of the considerable difference between the band gap of SnO2 and Ga2O3, the optical band gap of GTO NCs is strongly dependent on their composition. As evident from Figure 3a, the onset of the band edge absorption of GTO NCs shifts continuously to lower energy with decreasing Ga content up to band edge absorption for pure SnO2 NCs. In contrast to Ga2O3, the Bohr radius of SnO2 (2.7 nm) is comparable to an average size of hydrothermally-synthesized SnO2 NCs, at least in one dimension, causing a blue shift of their band edge absorption relative to bulk SnO2. It is widely accepted that oxygen vacancies are common structural defects in different metal oxides. In SnO2 these point defects form donor states ca. 0.03-0.15 eV below the conduction band minimum, delivering both high n-type electrical conductivity and transparency in the visible region of the spectrum.49 On the other hand, aliovalent doping of SnO2 with cations having lower oxidation state than 4+, induces acceptor centers within the bandgap. Theoretical calculations have suggested low formation energy of acceptors and high binding energy between acceptors and donors (oxygen vacancies), resulting in the formation of donor-acceptor complexes.50,67,68 Following the excitation of rutilephase GTO NCs into the band gap, the electrons and holes are trapped in the donor and acceptor states, respectively, which then act as radiative recombination centers. Similar mechanism has been suggested for γ-Ga2O3 NCs, involving oxygen vacancy and gallium-oxygen vacancy pair as donor and acceptor sites, respectively.15,28,30,32 Large energy separation between excitation (Figure 3b, blue trace) and emission (Figure 3b, orange trace) bands for all GTO NCs is consistent with PL originating from localized intra-band-gap states. Neglecting the lattice


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phonon effects, the energy of the DAP emission (EDAP) can be described by the following equation.

EDAP = Eg − ( ED + EA ) +

e2 4πε r

(4)

where Eg is the band gap energy, ED and EA are donor and acceptor binding energies, respectively, and r is the donor-acceptor separation. The last term in eq 4 represents the Coulomb interaction between singly-charged donor and acceptor sites, which has been shown to be the determining parameter of the DAP emission energy dependence on the NC size.32 Increase in the average donor-acceptor separation with increasing NC size leads to a red shift of the DAP PL band. The range of separations between donor and acceptor sites within a NC is also responsible for a red shift of the delayed emission relative to the steady-state PL spectrum (vide infra), as shown in Figure 3b (green trace).15,33 However, the dependence of the DAP emission energy on the NC composition is mostly due to the change in the electronic structure. The DAP band follows the blue shift of the band edge absorption with increasing Ga content (Figure 3c). The increase in Eg has a larger effect on EDAP than the decreased average separation between donor and acceptor sites caused by Ga doping (Figure 3e). The ability to manipulate the difference in energy between donor and acceptor states in GTO NCs by composition control allows for an expansion of the PL tuning range throughout the entire visible part of the spectrum. Photograph of the light emission of colloidal GTO NCs having different Ga content is shown in Figure 3d, attesting to their potential for photonic, optoelectronic, and biomedical applications.


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Figure 3. (a) Absorption spectra of GTO NCs with different Ga content, as indicated in the graph. The numbers in the legend represent nominal Ga concentration (in atom %). (b) Excitation (blue trace) and PL (steady-state, orange trace; delayed, green trace) spectra of GTO NCs with 50 % Ga content. Absorption spectrum (red trace) is shown for comparison. (c) PL spectra of GTO NCs with different Ga content, as indicated in the graph. (d) Photograph of the emission of colloidal GTO NCs with various Ga content. (e) Schematic illustration of the alteration of sub-band-gap trap states due to the change in the band gap width. The photoexcited electrons and holes are trapped in donor and acceptor states, located below CB and above VB, respectively. Radiative recombination of the carriers in GTO NCs with smaller band gap (low Ga content) is responsible for yellow-orange emission (right). Increased Ga content results in band


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gap widening and increased energy separation between donor and acceptor levels, causing the blue emission (left). The separation between donor and acceptor sites also determines the decay dynamics of the DAP emission.69,70 The rate determining step in the DAP recombination is a tunnel transfer of the electron trapped in the localized donor state to the associated acceptor site where the charge carriers recombine. The electron and hole trapped in the nearest donor and acceptor site, respectively, have the highest probability of recombining first.70 More distant donor-acceptor pairs recombine subsequently, causing a red shift and a delay in the observed emission (see Figure 3b). Typical time-resolved PL data for GTO NCs containing different amount of gallium are shown in Figure 4a. The average lifetimes and the time-decay fitting parameters for samples synthesized under different conditions are shown in Table S3 (Supporting Information). The average DAP recombination lifetime shortens with increasing Ga:Sn ratio in rutile-phase GTO NCs (i.e., nominal Ga % < 50), which is consistent with the reduction in the DAP separation owing to an increased density of donor-acceptor pairs. Any Ga2O3 clusters that may be formed as an undetected secondary phase would have a comparatively narrow-band DAP PL at substantially higher energy and decay rate eliminating the gallium secondary phase as a source of the observed decrease in the measured average lifetime. In principle, the highest density of the defect-based recombination sites is formed for the maximum concentration of aliovalent dopants. The closest associated donor-acceptor pairs will recombine with the highest efficiency, suggesting the possibility of increasing PLQY by incorporation of aliovalent dopants. Figure 4b compares the PLQY of GTO NCs having different composition. The quantum efficiency increases with increasing Ga content, and reaches the maximum value at ca. 50 % nominal Ga content. Further increase in Ga:Sn ratio leads to diminished PLQY, concomitantly with a


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structural transformation to cubic γ-Ga2O3 phase and decreasing doping level of Sn4+. The results of Figure 4b confirmed our hypothesis that higher concentration of aliovalent dopants in both NC lattices induces smaller average separation of donor-acceptor pairs resulting in more efficient radiative recombination. It should be noted that the effect of doping concentration on PLQY might be even more pronounced than shown in Figure 4b given the possibility that the actual doping concentrations are lower than the measured compositions due to conceivable dopant segregation (see discussion above). This is in contrast to isovalent doping (i.e., In3+ in Ga2O3 NCs), which leads to significant reduction in PL efficiency.16 The highest PLQY we measured for GTO NCs is ca. 45 %, which is one of the highest values reported for defect-based emission in oxides.

Figure 4. (a) Time-resolved PL decay of selected GTO NCs synthesized for 8 h. The Ga content corresponding to each sample is shown in the graph. The PL lifetime data for other samples were left out for clarity. (b) PLQY of GTO NCs as a function of Ga content.


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To probe the impact of the reaction duration on the optical properties of GTO NCs, we synthesized NCs with 50 % nominal Ga content by varying the reaction time between 4 and 32 hours. The EDX measurements confirmed that the composition did not change with prolonging the reaction time (Table S2 in Supporting Information). However, as the reaction time increases, the NCs continue to grow and their band edge shifts to lower energy due to reduced quantum confinement of SnO2 NC host lattice (Figure 5a). The prolonged growth of NCs also leads to a small decrease in donor/acceptor defect density and an increase in their average separation, causing minor but detectable red shift of the DAP emission band (Figure 5b). Surprisingly, the PLQY enhances from ca. 25 % to 43 % by increasing the reaction duration from 4 h to 32 h (Figure 5d, green symbols). A decrease in the PL decay rate with increasing reaction time (Figure 5c) is consistent with a larger average donor-acceptor separation, although it is evident that the change in the PL decay rate is much more significant than the shift in the steady-state emission band. The overall lattice disorder, stemming from heavy doping, concurrently leads to the formation of non-radiative recombination centers. The improvement in the long-range lattice order, caused by increasing the reaction time, stimulates the radiative recombination pathway and therefore enhances the PL efficiency. Radiative (kr) and non-radiative (knr) recombination rates were estimated utilizing the following equations:

τ=

1 ( kr + k nr )

Qx =

kr ( kr + k nr )

(5)

(6)


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where τ and Qx are the average PL lifetime and quantum yield, respectively. Using the average lifetimes estimated form the fitting of the data in Figure 5c (see Table S3 in Supporting Information) and the corresponding relative PLQYs, we calculated kr and knr for different synthesis times. Figure 5d plots the non-radiative recombination rate (red circles) and PLQY (green circles) with respect to the reaction duration. With extended reaction time radiative contribution to the PL decay increases, while non-radiative contribution decreases. A decrease in knr is accompanied by a symmetric increase in PLQY, highlighting the utility of hydrothermal synthesis and aliovalent alloying in tailoring the electronic structure and optical properties of metal oxide nanostructures.

Figure 5. (a) Absorption spectra of GTO NCs with 50 % Ga content, synthesized for varied lengths of time, as indicated in the inset (inset shows Tauc plots for the spectra in the main panel). (b) PL spectra of GTO NCs in (a). (c) Time-resolved PL decay of GTO NCs in (a). (d) PLQY and non-radiative recombination rate constant for GTO NCs as a function of the reaction duration.


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To investigate the effect of the synthesis temperature on the efficiency of the GTO NC emission, we performed the NC synthesis at different temperatures between 140 oC and 220 oC, keeping the other parameters unchanged. The increase in the synthesis temperature leads to the improvement in the sample crystallinity (see Figure 1c and the associated discussion). However, the Ga content drops from 43 % (for 140 oC synthesized NCs) to 28 % (for 220 oC synthesized NCs). Lower Ga concentration and larger NC sizes are both responsible for the red shift of the band gap absorption (Figure 6a) and DAP PL band (Figure 6b) with increasing temperature. Similar to reaction time dependence in Figure 5, the influence of synthesis temperature on the PL lifetime and quantum yield must be discussed in the context of the competition between the crystallinity of and DAP concentration in the NCs. On one hand, the NC treatment at elevated temperatures in solution diminishes the fast non-radiative recombination pathway, leading to an increase in the DAP PL lifetime and PLQY. On the other hand, it reduces the concentration of radiative centers by removing Ga3+-induced defects, leading to a decrease in the PL efficiency. As depicted in Figure 6d, while kr experiences relatively little change over the investigated temperature range, knr drops sharply when the temperature is raised from 140 oC to 190 oC, followed by a slight upward turn at higher temperatures. PLQY at first increases with increasing reaction temperature and then significantly declines (Figure 6d, green spheres). The maximum PLQY is achieved for the reaction temperature of ca 190 °C due to the optimal balance between DAP concentration and long-range structural order.


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Figure 6. (a) Absorption spectra of GTO NCs with 50 % nominal Ga content, synthesized in ethanol at different temperatures, as indicated in the inset (inset shows Tauc plots for the spectra in the main panel). (b) PL spectra of GTO NCs in (a). (c) Time-resolved PL decay of GTO NCs in (a). (d) PLQY and electron-hole recombination rate constants (radiative and non-radiative) for GTO NCs as a function of reaction duration. The ever increasing global demand for lighting and the necessity to lower energy consumption have, in recent years, placed an emphasis on finding new ways of improving the efficiency and characteristics of light emitting phosphors and devices. One of the emerging ways to maximize the efficiency of a light source is to design a light emitting material that can most effectively stimulate human eye under particular lighting conditions. Human vision occurs by two primary modes: photopic and scotopic vision. Photopic vision is the vision under well-lit conditions and is primarily enabled by cone cells in the eye. It is responsible for color perception, visual acuity and temporal resolution. Scotopic vision, on the other hand, is vision in dark, which is monochromatic and functions due to rod cells in the eye. A combination of photopic and


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scotopic vision at low lighting levels is known as the mesopic vision, and involves a function of both rod and cone cells. The mesopic luminous efficiency is defined as a linear combination of photopic and scotopic luminosity functions:

Vm (λ ) = (1 − x)V ′(λ ) + xV (λ )

(7)

where V (λ ) and V ′(λ ) are is the standard photopic and scotopic luminosity functions, respectively (Figure 7a), and x is the function of photopic adaptation luminance (Lp). Therefore, the objects are perceived brighter in dark when the light source maximum emission is at ca. 505 nm (scotopic function), while in well-lit environment the maximum emission at ca. 555 nm (photopic function) allows the objects to be identified brighter. Light sources with scotopic to photopic ratio (S/P) of ca. 2.5 are recognized to be visually most efficient. Such light sources can be operated at lower power relative to those with low S/P ratio.71,72 The ability to simultaneously manipulate PL properties of GTO NCs, including emission energy, spectral band width, and quantum efficiency, renders GTO NCs promising for designing light sources with optimal S/P ratio. We prepared a series of GTO NCs under diverse reaction conditions (Table S4 in Supporting Information), allowing us to generate samples that have maximum emission wavelengths matching those of the maxima of both scotopic (Figure 7a) and photopic (Figure S4 in Supporting Information) luminosity functions. We have measured the PLQY of ca. 34 % for S/P = 2.60 (see Figure S5 in Supporting Information for sample characterization), which is approximately 2 times greater efficiency than that reported for CdSebased QDs with the similar S/P ratio (Figure 7a).71,72 The photograph showing the yellowishgreen emission of this sample under UV excitation is displayed in the inset of Figure 7a. To


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generate white light for lighting applications we prepared nanoconjugates by functionalizing these NCs with dye molecules emitting in the complementary spectral range (Figure 7b).5,73 These hybrid organic-inorganic nanostructures act as single nanophosphors owing to the distance-dependent energy transfer between the two moieties (Figure S4 in Supporting Information).5,35,73 A photograph of a typical white light emitting nanoconjugate sample is also shown in Figure 7b (inset). These results serve as an instructive example how control of the electronic structure and defect interactions by aliovalent doping/alloying, and the choice of the synthesis method can result improved PL efficiency and characteristics of metal oxide NCs for lighting and other photonic applications.

Figure 7. (a) PL spectrum of GTO NCs (green trace) covering both scotopic (blue trace) and photopic (red trace) luminosity functions. The function corresponding to S/P ratio of 2.5 is shown as a brown trace. The photograph of the emission of the colloidal NCs excited into the band gap is shown in the inset. (b) PL spectrum of GTO NCs in (a) conjugated with Atto-590.


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The photograph of the emission of the nanoconjugates excited into the band gap is shown in the inset.

CONCLUSIONS In summary, we synthesized new ternary gallium tin oxide NCs, and demonstrated that their PL color can be tuned from blue to orange-red by composition modulation. Substitutional incorporation of Ga3+ into SnO2 lattice creates acceptor levels within the band gap, while simultaneously increasing the concentration of intrinsic donors ( VO• ) to compensate for charge imbalance. Using SnO2 NCs in which majority of the photogenerated carriers tend to recombine non-radiatively, we showed that induced donor-acceptor pairs introduce a new radiative recombination pathway, resulting in the enhancement of the PLQY to above 40 %. The observed blue shift of the DAP emission of GTO NCs with increasing Ga content originates from widening of the band gap and stronger Coulomb interaction between the donor and acceptor sites. Systematic increase in the DAP concentration, associated with an increase in Ga3+ doping level, leads to smaller separation between electron and hole trap sites, resulting in a faster PL decay rate. Non-radiative recombination mechanisms can be further suppressed by improving the long range order of NCs through increased reaction duration and temperature. The interaction of the dopant impurity with the NC host lattice native defects is a function of its chemistry, and serves as another degree of freedom to control the electronic structure and recombination mechanism in both TMO NCs examined in this work. Broad PL tunability, chemical stability, and efficiency of these TMO NCs are compelling features for various photonic applications, including the generation of light sources having ideal scotopic/photopic ratio and appreciable PLQY. The results of this work demonstrate the possibility of converting non-luminescent TMO


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NCs into tunable and efficient light emitters by simultaneous control of the NC composition and defect interactions. This process is enabled by judicious selection of the synthesis method and reaction conditions, and allows for a design of new light emitting devices with optimal characteristics.

ASSOCIATED CONTENT Supporting Information. GTO NC size distribution histograms, TEM images, XRD patters, elemental analysis data, PL spectra and lifetime measurements, S/P and PLQY data. This material is available free of charge via the Internet at http://pubs.acs.org.

AUTHOR INFORMATION Corresponding Author *[email protected] Notes The authors declare no competing financial interest. ACKNOWLEDGMENT This work was supported by the Natural Sciences and Engineering Research Council of Canada (RGPIN-2015-67304032) and Canada Foundation for Innovation (Project number: 204782). P.V.R. acknowledges the support from the Canada Research Chairs program (NSERC).


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Nanocrystals into Tunable and Efficient Light Emitters by Aliovalent Allo

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