Nanoparticles of I–III–VI semiconductors are promising candidates for novel non-toxic fluorescent materials. However, removal of defect levels responsible for their broad-band emission has not been successful to date. The present study demonstrates, for the first time, the coating of core AgInS2 nanoparticles—one of the I–III–VI group semiconductors with a bandgap in the visible region—with III–VI group semiconductors. The AgInS2/InSx and AgInS2/GaSx (x = 0.8–1.5) core/shell structures generate intense narrow-band photoluminescence originating from a band-edge transition at a wavelength shorter than that of the original defect emission. Microscopic analyses reveal that the GaSx shell has an amorphous nature, which is unexpected for typical shell materials such as crystalline lattice-matching ZnS. Single-particle spectroscopy shows that the average linewidth of the band-edge photoluminescence is as small as 80.0 meV (or 24 nm), which is comparable with that of industry-standard II–VI semiconductor quantum dots. In terms of photoluminescence quantum yield, a value of 56% with nearly single-band emission has been achieved as a result of several modifications to the reaction conditions and post-treatment to the core/shell nanoparticles. This work indicates the increasing potential of AgInS2 nanoparticles for use as practical cadmium-free quantum dots.
Semiconductor nanoparticles of sufficiently small size have their optical properties determined by quantum confinement, which causes a shift in the energy of an optical transition with crystal size. Such semiconductor nanoparticles (NPs), or quantum dots (QDs), have attracted considerable attention following the development of methods for synthesizing highly monodisperse, well-stabilized, and intensely emitting NPs1,2. Cadmium selenide and its alloys with other II–VI semiconductors have been at the center of research activities due to their narrow band-edge emission deriving from their characteristics as direct semiconductors3. In addition to being the subject of academic and industrial research, CdSe and other II–VI semiconductor QDs have been increasingly finding commercial applications in liquid-crystal displays and television sets to optimize the color composition of backlights, in lighting to improve the quality of light for specific applications and in biological research as labels for bioimaging and drug-delivery components4. With the rapid spread of applications, however, the high toxicity of cadmium is becoming a matter of increasing concern, and CdSe and related QD materials are likely to face a ban in many countries. The search for a replacement for toxic CdSe QDs with non-toxic alternatives with comparable properties, such as stability, size monodispersity, high emission quantum yield (QY), narrow-band emission spectrum and chemical ‘wet’ preparation, is therefore an utmost priority for many research groups. Recent attempts to substitute other materials for toxic cadmium have successfully utilized III–V semiconductors such as InP, InAs, and GaAs5,6,7 and I–III–VI ternary semiconductors represented by CuInS2, AgInS2, and their selenide derivatives8,9,10,11. Although the quantum yield of the latter has been improved to exceed 50%, the lack of monochromaticity of their photoluminescence (PL), represented by broad PL spectra, has been recognized as a major problem12,13,14. There are at least two possible origins of the spectrally broad PL from ternary semiconductor QDs: defects on the particle surface and/or defects inside the crystals. The structural diversity characteristic of ternary materials is considered to be the origin of various donor–acceptor levels, which have recently been identified by PL decay and single-particle analyses15. Efforts to generate intense, defect-free semiconductor band-edge PL from I–III–VI semiconductor nanomaterials have not been successful to date. Some researchers have even concluded from time-resolved and X-ray measurements that emission originating from several sets of donor–acceptor pairs is inevitably generated in I–III–VI-based and I–III–VI-based semiconductor nanomaterials, giving rise to PL spectra with a full width at half maximum (FWHM) larger than 100 nm, or 300 meV13,14,16. Although recent studies by our group on AgInTe2 nanorods have led to a relatively narrow emission (FWHM of 100–150 nm or 110–180 meV) in the infrared region (with PL peak of 1000–1100 nm), the PL origin was attributed to donor–acceptor transitions located close to the band edge17.
In the present study, we developed a reproducible method for obtaining spectrally narrow band-edge PL from AgInS2 NPs in the visible region. The band-edge emission was achieved by coating AgInS2 NP cores, synthesized by a conventional heating method, with III–VI group semiconductor shells. As a result, we observed PL that was nearly an order of magnitude narrower than that of uncoated AgInS2 NPs (FWHM of 28.6 nm, or 103 meV, compared with FWHM of 220 nm, or 388 meV, respectively). Single-particle analysis revealed even higher monochromaticity at the single-NP level, with an average PL spectral width FWHM of 24.0 nm (or 80.0 meV) at room temperature. This value is now fully comparable with that of industry-standard II–VI semiconductor QDs, such as CdSe and CdTe3, and less than half of that of other alternative cadmium-free QDs, such as InP7. The final section discusses the great potential of these cadmium-free QDs by demonstrating a several-fold improvement of PL QY as a result of modifications to the synthetic conditions and post-treatment of the as-synthesized core/shell NPs. At present, QDs with a PL QY as high as 56% have been obtained reproducibly.
Materials and methods
Silver acetate (Ag(OAc), Wako), In(OAc)3 (Aldrich), Ga(acac)3 (Aldrich), 1-dodecanethiol (Wako), tri-n-octylphosphine (TOP, TCI), thiourea (Kishida Chemical Co., Ltd), 1,3-dimethylthiourea (TCI), and elemental sulfur (5N, Mitsuwa Chemical Co., Ltd) were used without further purification. Oleylamine (Aldrich) was used after vacuum distillation in the presence of calcium hydride.
In the synthesis procedures, samples were heated in a two-neck flask using a heating mantle. The solution temperature was directly measured, and the heater was controlled to fix the solution temperature. PL and UV–vis spectra were recorded with a multichannel PL analyzer (Hamamatsu, PMA-12) and a UV–vis spectrometer (JASCO, V-670). PL QYs were measured by attaching an integrating sphere to the PL spectrometer. PL lifetimes were recorded with a time-correlated single-photon counting setup (Hamamatsu, Quantaurus-Tau). Electron microscopy observations were performed with a Hitachi H-7650 TEM (an acceleration voltage: 100 kV) for TEM and Cs-corrected HR-STEM (JEM-ARM-200F, JEOL Co. Ltd., an acceleration voltage: 200 kV) for high-resolution TEM (HRTEM) and high-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM). XRD patterns were recorded with a powder X-ray analysis apparatus (Rigaku, SmartLab) equipped with a parallel-beam/parallel-slit analyzer. XPS spectra were recorded with an X-ray photoelectron spectrometer (Shimadzu, KRATOS AXIS-165x) equipped with an aluminum target. The apparatuses used for single-particle measurements are described in a separate section below.
Synthesis of AgInS2 NPs at constant heating rate
In a two-neck flask, Ag(OAc) (0.4 mmol) and In(OAc)3 (0.4 mmol) were mixed with thiourea (0.8 mmol) in 12 mL of an oleylamine/1-dodecanethiol mixed solvent, typically in a ratio of 59/1 (vol/vol). The mixture was degassed and filled with Ar, heated to 200 °C at a rate of 10 °C/min and then cooled to room temperature. After removing large precipitates by centrifugation, a clear red solution was obtained. The smallest possible amount of methanol was added to precipitate the NPs. After centrifugation, AgInS2 NPs capped with oleylamine were obtained as a solid.
Synthesis of III–VI semiconductor shells from elemental sulfur
Oleylamine-capped AgInS2 NPs (30 nmol in terms of particle) were dissolved in 12 mL oleylamine, and then either In(OAc)3 or Ga(acac)3 (0.1 mmol) and elemental sulfur (0.15 mmol) were mixed at room temperature. The mixture was degassed and filled with argon, then heated to 200 °C (InSx shell) or 260 °C (GaSx shell) at a rate of 10 °C/min. The temperature was held constant for 60 min, and aliquots of the mixture were removed from the flask intermittently for PL measurements. The resulting AgInS2/GaSx core/shell NPs were purified by precipitating with methanol followed by centrifugation, and the solid was dissolved in hexane.
Synthesis of small AgInS2 NPs at low temperature
For AgInS2 core NPs, a slight modification to the aforementioned method improved the monodispersity. Ag(OAc) (0.4 mmol) and In(OAc)3 (0.4 mmol) were mixed with oleylamine (8 mL) and 1-dodecanethiol (1.25 mmol), and the temperature was then increased to approximately 80 °C under an Ar atmosphere. Thiourea, recrystallized from water, was added to the flask, and the temperature was increased to 135 °C at a rate of 30 °C/min with stirring. The temperature was maintained for 10 min and then cooled to room temperature immediately by dipping the flask into a water bath (25 °C). The NPs were purified as described in the previous section and dispersed in oleylamine (2 mL).
Synthesis of GaSx semiconductor shells from 1,3-dimethylthiourea
Approximately 30 nmol (in terms of NPs) of AgInS2 core NPs were mixed with an oleylamine (7 mL) suspension containing Ga(acac)3 (0.1 mmol) and 1,3-dimethylthiourea (0.1 mmol). The reaction mixture was heated rapidly to 230 °C, and the temperature was then increased to 280 °C at 2 °C/min under an Ar atmosphere. The solution was then cooled to room temperature in air, during which the solution was degassed for several minutes to remove volatile, unreacted sulfur species. The core/shell NPs were purified as previously described and then dissolved in chloroform (4 mL). Finally, the same volume of TOP was added to the chloroform solution and incubated for several hours at room temperature.
Single-particle imaging and spectroscopy were performed on an epifluorescence microscope setup. Briefly, an inverted optical microscope (Olympus, IX71) equipped with a high-numerical-aperture objective (Olympus, UPlan FLN 100×/1.3 NA, Oil) was used for PL imaging. Samples were excited with a 375 nm (~13 W cm−2) diode laser (PicoQuant, LDH-PC-375), and PL was detected using electron-multiplying CCD camera (Andor Technology, iXon) after passing through an imaging spectrograph (Bunkou Keiki, CLP-50, 0.5 nm resolution). A 400 nm longpass filter (Thorlabs, FELH0400) was used to suppress unwanted excitation and scattered light in the detection path. Data were recorded as image sequences at 0.05 s (20 Hz) and 1 s (1 Hz) acquisition times for imaging and spectroscopy, respectively. All single-particle measurements were performed at 295 K under a N2 environment. For sample preparation, AgInS2/Ga2S3 core/shell NPs were dispersed in toluene and spun-cast onto precleaned quartz coverslips (1 × 0.2, Technical Glass Products) at 2000 rpm. The concentration of NPs was optimized to obtain well-separated (typically with 1 µm interval) and randomly distributed diffraction-limited emission spots (Fig. S1a–c, supplementary information). Well-separated spots showing characteristic single-step blinking and photobleaching (Fig. S1d, supplementary information) were considered to have originated from a single emitter.
Results and discussion
Choice of shell materials
Preparation of a type I core/shell structure, meaning that the bandgap of the shell covers the bandgap of the core, is a common strategy for obtaining strong and stable emission from semiconductor NPs18. Zinc sulfide (ZnS) has been widely used as a shell material for II–VI and III–V semiconductor NPs because of its chemical stability, high crystallinity, and wide bandgap (3.7 eV), which is crucial for type I band alignment. Coating with ZnS has also been attempted for CuInS2 ternary NPs by reacting them with a mixed solution of zinc and sulfur sources19. The treatment resulted in blueshifts in PL, in contrast to the redshift generally observed for semiconductor NPs after overcoating with larger-bandgap semiconductors. The expected redshift is usually explained by the leakage of excitons from the core into the shell, whose bandgap is essentially smaller than the HOMO-LUMO gap of insulating organic ligands on its surface20. The blueshift, in contrast, is often explained as the result of changes in surface defect levels by surface passivation with ZnS and/or the result of partial alloying by cation exchange from copper and indium into zinc21,22,23. In our experiments, we also observed a blueshift in PL when AgInS2 NPs were reacted with zinc acetate (Zn(OAc)2) and elemental sulfur at 200 °C, as shown in the optical spectra in Fig. 1a, b.
According to the abovementioned considerations, the semiconductor used for the shell should be selected from materials that do not easily form alloys, or solid solutions, with the core NPs. However, the formation of defect levels between the core and shell must be avoided. Considering the band-edge criteria for the type I core/shell structure, indium sulfide, which shares two of the three elements composing the core AgInS2 and has a wider bandgap (2.2 eV) than that of the core (1.8 eV for bulk), emerged as a candidate material. The phase diagram for the Ag2S–In2S3 quasi-binary system shows that AgInS2 forms continuous solid solution series with In2S3 only within a limited region of high indium content (81–96% In2S3)24, which is less miscible than with ZnS25. Figure 1c, d shows UV–vis absorption and PL spectra obtained after the AgInS2 core NPs were reacted with indium acetate (In(OAc)3) and elemental sulfur. The indium sulfide (InSx) shell was synthesized by slowly increasing the temperature of the oleylamine solution of the AgInS2 core in the presence of these precursors (10 °C/min up to 200 °C) and maintaining this temperature for 1 h. The resulting PL spectrum was apparently different from that of the ZnS coating, with the emergence of a new narrow PL peak on the short-wavelength side of the original defect emission. In contrast, InSx NPs synthesized by a method similar to that used for the shell coating was completely non-luminescent, confirming that the narrow PL peak originated from the AgInS2/InSx core/shell structure. However, due to insufficient passivation, substantial broad emission persisted, with a blueshift of 70 nm, suggesting that part of the indium source participated in changing the core composition to an indium-rich form (e.g., AgIn5S8) during the reaction26.
The monochromaticity of the emission was further greatly improved by coating the AgInS2 core NPs with gallium sulfide (GaSx) instead of with InSx. First, a structure composed of several shoulders appeared in the UV–vis spectrum, resembling the absorption of size-quantized, well-stabilized semiconductor QDs (Fig. 2a)27. In Fig. 2b, PL spectra are shown for different elapsed times (0–60 min) after the reaction temperature reached 260 °C. The broad-band defect emission (>600 nm) of the uncoated AgInS2 NPs decreased with the formation of the gallium sulfide shell, followed by the appearance of a new narrow peak on the short-wavelength side of the defect emission, as demonstrated in the imaged obtained under UV illumination (Fig. 2c). Similarly to the indium sulfide of the same group, no single phase has been reported between AgInS2 and Ga2S3 according to the quasi-ternary phase diagram of Ag2S, In2S3, and Ga2S328. The great improvement of monochromaticity might be attributed to the wider bandgap of gallium sulfide (2.8 eV) compared with that of indium sulfide (2.2 eV) (Fig. 2d). The similarity in wavelength of the newly generated PL peaks between InSx and GaSx-coated AgInS2 NPs ruled out the possibility of cation exchange between indium in the core and gallium in the shell.
To the best of our knowledge, such drastic changes in the PL and UV–vis spectra only appeared when the AgInS2 NPs were reacted with indium (III) or gallium (III) species in the presence of sulfur and were not been observed when the mixture was heated in the absence of one or both precursor materials. When the core NPs were heated at 270 °C for 1 h in oleylamine without adding any precursor materials, the NPs partially agglomerated and the quantum yield decreased by approximately 80%, possibly due to Ostwald ripening (Figure S2, supplementary information). Therefore, the emergence of narrowband emission, which reflects a diminishing contribution of the defect sites to the PL, likely originated from the NPs’ surface modification, rather than from the removal of defect levels inside the crystals or the transformation of the crystalline phase by the applied heat. Several studies have investigated the mechanism of broad defect emission with a large Stokes shift13,14,16,26. The defect emission is often explained in relation to the existence of localized donor defect or acceptor defect states caused by vacancies, interstitials or anti-sites within the crystal structure or by structural defects on the surface. The defect states are often proposed to explain the fact that the actual composition of I–III–VI semiconductors often deviates from the stoichiometric composition. However, the present results indicate the prevailing importance of surface conditions in generating band-edge emission.
Origin of narrow-band emission
Figure 3a shows PL spectra of the original bare AgInS2 NPs and of those coated with gallium sulfide in chloroform after isolation of the NPs via the precipitation/dissolution method. The PL QY was 12.4% for the sharp band emission (530–630 nm) and 21.7% for the entire PL range (530–900 nm). The narrowband emission of the core/shell NPs had a FWHM as small as 28.6 nm (103 meV), much smaller than that of the uncoated AgInS2 NPs (220 nm, 388 meV). The corresponding photoluminescence excitation (PLE) spectra recorded at the respective peak wavelengths are shown in Fig. 3b, along with a UV–vis spectrum of the core/shell NPs. A distinct excitonic feature at 530 nm was observed for the excitation spectrum of the core/shell NPs only, whereas the uncoated AgInS2 NPs showed a monotonic increase in PL intensity with decreasing excitation wavelength. The PL decay curves in Fig. 3c show a significant decrease in the lifetime components for the core/shell NPs compared with those of the uncoated AgInS2 NPs. The decay curves for the core/shell NPs were fitted to a three-exponential equation:
where I(t) is the PL intensity at time t; A1, A2, and A3 are amplitudes; and τ1, τ2, and τ3 are lifetimes. In contrast, the decay curve of the uncoated AgInS2 NPs was sufficiently fit with a single-component exponential. The fitting results are summarized in Table 1. For the core/shell NPs, the second component (24.8 ns) contributed to 45% of the total emission intensity and the third component (96.6 ns) contributed to 50% of the total emission intensity compared with the components described in terms of products of amplitude and lifetime (A × τ). These lifetimes are comparable with the lifetime of the band-edge emission from CdSe QDs29, indicating that the new short-wavelength peak that appeared after coating with gallium sulfide may originate from the band-edge transition as well.
Plots of (αhν)2 as a function of hν, generally known as Tauc plots, support this interpretation (Figure S3, supplementary information). Flat regions in these plots, which appeared both for AgInS2 core NPs (Fig. S3a) and for those coated with ZnS (Fig. S3b), InSx (Fig. S3c), and GaSx (Fig. S3d) shells, are spectroscopic features of direct allowed transitions and indicate their presence in these samples. The estimated bandgap was very similar between these four samples, indicating that alloying was limited under the experimental conditions used for ZnS coating. However, differences between these bandgap values and the PL peak energies, i.e., Stokes shifts, were significantly larger for the AgInS2 core NPs (0.59 eV) and for those coated with ZnS (0.37 eV) and much smaller for those coated with InSx (0.05 eV) and GaSx (0.06 eV), indicating a band-edge or near band-edge emission for the latter two samples.
Structural characterization of core/shell nanoparticles
The synthesized core-only and core/shell NPs were characterized by TEM. Figure 4a, c shows TEM images of AgInS2 NPs before and after the formation of gallium sulfide shells, respectively, as well as size histograms of the two types of NPs (Fig. 4b, d). The AgInS2 NPs show irregular, angular shapes that transition to rounded ones after the formation of the shell with a slight increase in the average particle diameter (0.4 nm) corresponding to the formation of 0.2 nm thick shell. However, the actual effect of overcoating on the average NP diameter is difficult to quantify at this stage due to the polydispersity of the original uncoated AgInS2 NPs. HRTEM (Fig. 4e) and HAADF-STEM (Fig. 4f) images indicate the presence of an amorphous layer surrounding the crystalline AgInS2 NPs, suggesting that the GaSx shell is amorphous. Elemental analyses for the ensemble showed an increase in gallium and sulfur after shell formation. Compared with the composition of the cores, that of the amorphous GaSx was closer to that of Ga2S3 rather than to that of another possible candidate, GaS (Table S1 and Figure S4, supplementary information). The amorphous nature of the shell was also supported by X-ray diffraction (XRD) analysis; the XRD patterns showed no changes in the NP crystal structure before and after the formation of the Ga2S3 shells (Fig. 5a). In the X-ray photoelectron spectroscopy (XPS) spectra, the positions of the peaks for silver, indium, and sulfur coincided with those of a metal sulfide, and they remained nearly unchanged during the formation of Ga2S3 shells, indicating that the NP cores were chemically unaffected by the formation of the shells (Fig. 5b, c, e, f). Although we could not obtain reference data directly for gallium sulfide, the measured peaks for gallium (Fig. 5g, h) were singular, and their position was different from that of the two possible byproducts, metal gallium and gallium oxide30,31. In terms of chalcogens, no additional peaks were detected for the core/shell NPs, indicating that the sulfur in the shell occurred in the form of a metal sulfide (Fig. 5i).
The concept of an amorphous shell is unusual compared with that of commonly used II–VI shells, such as ZnS and CdS, coated onto II–VI or III–V semiconductor NPs. Generally, II–VI shells have been preferred because of their good crystallinity and expected lattice matching at the interface between the core and shell materials used. Indium and gallium sulfides possess crystal structures that are different from those of both chalcopyrite and wurtzite AgInS2 and are known to contain many crystalline defects32,33,34. Therefore, these III–VI semiconductors have never been used as shell materials. However, we consider that the amorphous nature of the shell may sometimes be advantageous in stabilizing the complex surface elements on I–III–VI semiconductors, as this amorphous nature allows for small, continuous displacement of elements into energetically stable conformations. Similar effects have been observed for solar cells based on copper indium gallium sulfide (CIGS); in particular, the introduction of an In2S3 buffer layer between CIGS and n-type ZnO layers improved cell performance35,36. These buffer layers deposited by chemical vapor deposition (CVD) or atomic layer deposition (ALD) were reportedly amorphous in nature, although a definite correlation between the amorphous nature and improved photovoltaic performance has not been proven.
We next used single-particle spectroscopy to gain insight into the extent of inhomogeneity in the band-edge emission. This method has been previously used to successfully uncover the multiple-defect nature of individual ZnS–AgInS2 solid solution NPs and to demonstrate suppression of blinking by two-color excitation37. In the context of the current study, the spectra of single core/shell NPs can directly answer the important question regarding whether the residual defect emission originates from the same NPs as the bend-edge emission or whether the two types of emission derive from physically distinct particles. Figure 6 shows nine examples of PL spectra from individual NPs immobilized on a quartz substrate. The single-particle nature of the emission was confirmed by measuring the concentration dependence of the sample and by observing the PL intermittency (blinking) phenomenon (Figure S1 in the supplementary information) from individual NPs. The examples in Fig. 6 were chosen from a set of more than 250 measured NPs, each analyzed for the first 10 s of its emission. As shown in Fig. 6, the top four NPs produced only band-edge type spectra, while the bottom five NPs produced only defect-type spectra. Statistically, 50.4% of the analyzed NPs showed band-edge emission, 36.8% showed defect emission and the remaining 12.8% NPs showed both types of emission simultaneously. The same sets of spectra as a function of photon energy are shown in Figure S5 in the supplementary information together with the spectrum of an NP ensemble used for the single-particle measurements. As shown in Fig. 6, the PL FWHMs from the individual NPs were smaller than the FWHM of the ensemble spectra. This observation indicates that the ensemble PL spectra include a distribution of peak wavelengths of individual particles and are thus inhomogeneously broadened, likely due to the size distribution of the core NPs. The PL peak energy and FWHM obtained from the observation of 250 individual NPs were analyzed and plotted as histograms, as shown in Figure S6 of the supplementary information. For the band-edge emission, the mean FWHM obtained from the histogram was as small as 80.0 meV (24.0 nm). Importantly, this value is close to the PL FWHM for single CdSe/CdS/ZnS core/shell/shell QDs embedded in polymer matrices (55–65 meV) recorded at room temperature38. Thus, the AgInS2/Ga2S3 core/shell NPs have considerable potential to serve as Cd-free alternatives in terms of emission properties.
Improvement of quantum yield
The low QY (12.4% for band-edge) and the remaining defect emission are major issues that must be improved for potential practical applications. Both are considered to originate from the insufficient removal of defect levels in both the core and the shell. The aforementioned ‘mix everything and heating up method’ used to synthesize AgInS2 core NPs generates particles with large size polydispersity, as shown in the TEM image and corresponding histogram presented in Fig. 4a, b. The polydispersity originates from several types of reactions occurring simultaneously during heating to 200 °C. When the as-synthesized core NPs were size separated before the formation of gallium sulfide shells, the smaller core fractions tended to yield more intense band-edge emission, suggesting that smaller NPs have fewer defect levels and/or are more favorably shelled. Therefore, we made modifications to the core synthesis conditions to selectively obtain smaller core NPs.
According to previous reports, AgInS2 NPs are most commonly generated through Ag2S as an intermediate product because Ag(I) species react more favorably with sulfur sources than with In(III) complexes. The two main roles of the intermediate Ag2S are to act as seed particles on which cation exchange of Ag+ and In3+ occurs39 and to serve as resource hosts that operate as catalysts for the reaction of indium and sulfur sources40. Theoretically, the former reaction leads to an approximately 15% decrease in particle volume, which has actually been observed39. The latter reaction would result in a three-fold increase in particle volume, which is undesirable in the present case. In our reaction system using thiourea as a sulfur source, the color of the solution began to change at 120 °C from yellow to orange and to dark brown, indicating the formation of Ag2S NPs. The target AgInS2 NPs were then formed at higher temperatures (>130 °C) either through cation exchange with In(III) species (at 130 °C–150 °C) or via reaction with indium sulfide source materials generally recognized as clusters before nucleation (above 150 °C). Therefore, in this study, the reaction temperature was kept low at 135 °C for 10 min to allow for the composition of the entire NP ensemble to change to AgInS2. This modified procedure successfully generated AgInS2 core NPs with a smaller particle diameter and a narrower size distribution than those of particles generated previously (Fig. 7a, b). The PL QY of the core NPs was 64% (Fig. 8a), which is a record high value for AgInS2 core-only NPs, even though the spectral shape is that of broad defect emission.
For the shell formation reaction, the sulfur source was changed from elemental sulfur to 1,3-dimethylthiourea, which reacts with Ga(acac)3 at higher temperatures (>230 °C) to generate gallium sulfide. Although elemental sulfur is a common sulfur source for many chalcogenide nanoparticles, we observed that it has a strong quenching effect on the PL of AgInS2 core NPs. Undesired contact with elemental sulfur in the early stage of the shell formation reaction could thus be avoided by the use of the sulfur compound. In addition, the change in atomic ratio between gallium and sulfur (from Ga:S = 1:1.5 to 1:1) led to gallium-rich conditions. The temperature gradient was also modified in accordance with the decomposition of 1,3-dimethylthiourea to avoid the nucleation of gallium sulfide as much as possible. These modifications to the synthesis protocol improved the PL QY of the AgInS2/GaSx core/shell NPs to 15–20% for the band-edge emission and to 25−30% for the entire spectral region (Fig. 8b). For the band-edge emission, this increase represents an improvement by a factor of 1.5 over the PL QY obtained using the original synthesis conditions. The thickness of the shell, calculated from the average diameter change before and after shell formation, was increased from 0.2 nm for the previous method to 0.6 nm (Fig. 7b, d), although the shell was still thinner than the expected value (1 nm) calculated based on the amount of materials. An HRTEM image clearly shows the existence of 0.5–1.0 nm amorphous shells around crystalline cores measuring 3–5 nm in diameter, depending on location (Fig. 7e). The difference between the expected bandgap and PL peak energy became even smaller (0.02 eV), indicating better passivation (Figure S3e, supplementary information). Further improvement of the PL QY was achieved by post-treatment of the AgInS2/GaSx core/shell NPs with TOP. The post-treatment was carried out after removing the NPs from the solution used for shell formation. The intensity of the band-edge emission nearly doubled within a few hours in the presence of TOP, with a slight decrease in the remaining defect emission, keeping the UV–vis absorption profile nearly unchanged (Fig. 8c). As a result, nearly single-band emission with a QY as high as 56% was obtained with high reproducibility, yielding a brilliant yellow emission bright enough to observe under ambient light (Fig. 8d and Figure S7, supplementary information). The crystal structures of the AgInS2 core and AgInS2/GaSx core/shell NPs were similar to those prepared by the original method, but a purer tetragonal phase was observed (Figure S8, supplementary information). These findings are well consistent with a previous report by Doh et al.39 that showed that 4 nm Ag2S NPs were converted to tetragonal AgInS2 NPs, while larger (8 nm) NPs were transformed into orthorhombic AgInS2 NPs, upon cation exchange with In3+ species. According to that report, Ag2S seed particles are prone to conversion to orthorhombic AgInS2, which is structurally more favorable, whereas the thermodynamically more stable tetragonal phase can be generated when the particle size is small enough (<4.5 nm for seed Ag2S) to invoke a rearrangement.
The remarkable improvement in PL QY by the TOP treatment was surprising, and we offer the following hypothesis. The gallium sulfide shell passivates the defect levels on the surface of the AgInS2 core NPs, which leads to the generation of the band-edge emission. However, excitons could still be leaking outside of the core due to an insufficient potential barrier between the core and the thin amorphous gallium sulfide shells. The major traps for electrons were, potentially, zero-valent sulfur sites existing on the surface of the gallium sulfide shell. These traps operate as acceptors for photogenerated electrons. TOP is well known to coordinate with chalcogens; therefore, the TOP treatment has the potential to passivate the electron traps on the surface of the gallium sulfide shells, either by filling electrons in the LUMO level of the surface sulfur sites or by removing the sulfur sites via conversion to a more stable tri-n-octylphosphine sulfide. Elemental analyses of the core/shell NPs synthesized with the modified procedure showed a decrease in the ratio of sulfur in the shell (to Ga:S = 1:0.84) from that observed for the shells synthesized using elemental sulfur with a 1.5 equivalent of Ga(acac)3 (Table S2 and Figure S9, supplementary information). The abovementioned observations suggest that the amount and valency of sulfur may be key factors controlling the PL QY of core/shell NPs.
Unexpectedly, however, the PL decay curves obtained before and after the TOP treatment did not show such drastic changes (Fig. 9). The PL lifetimes, which nearly doubled those of the core/shell NPs synthesized using >6 nm cores and elemental sulfur for the shell source material (Table 1), remained unchanged after the TOP treatment (Table 2). These results suggest suppression of blinking, which itself is generally caused by NP charging. We indeed observed strong blinking with long off-times for the AgInS2/GaSx core/shell NPs synthesized by the first method (Figure S1, supplementary information), similar to the blinking behavior of uncoated core NPs37. However, Figure S10 shows the time traces of PL intensity of individual core/shell NPs synthesized with the modified method before and after TOP treatment. The core/shell NPs without TOP treatment (Figure S10, bottom, red traces) showed short ‘on’ durations with an on/off frequency often higher than the recording frame rate (20 Hz). This behavior was substantially altered by the addition of TOP, which resulted in longer and more stable ‘on’ states (Figure S10, top, blue traces). Therefore, the doubling of the PL QY by TOP treatment is mostly attributed to the increase in ‘on’ duration and was likely caused by the effective blocking of charges from escaping from the core. In addition, it is worth mentioning that the TOP-treated core/shell NPs exhibited a PL QY of 56.0%, which is comparable to that of the core NPs (64%), suggesting that we were successful in altering the pathway of photogenerated carriers without causing a significant efficiency loss. Studies aimed at further increasing the PL QY and identifying the quantitative aspects of the corresponding excitonic characteristics are underway.
In summary, we have demonstrated band-edge emission from AgInS2 NPs by coating them with III–VI semiconductors, GaSx shells, with a composition ranging from x = 0.8 to x = 1.5. The new band-edge PL band has a FWHM as small as 30 nm in an ensemble and 20 nm at the single-particle level. The core/shell structure was successfully observed by HRTEM and STEM, which showed that, unlike conventional crystalline II–VI semiconductors shells, the GaSx shell was amorphous. The successful demonstration of surface trap removal by the amorphous layer may indicate a new direction in the design of core/shell nanostructures. Similarly to the role of organic ligands that bind to specific elements to remove intra-bandgap defect levels, ‘soft’ non-crystalline inorganic shells might prove advantageous in stabilizing ternary, quaternary semiconductor NPs that have complex surface structures. In terms of PL QY, the AgInS2 NP cores obtained using the ‘mix everything and heating up method’ demonstrated a value representing 12.4% of the band-edge emission after coating with gallium sulfide. This value was further improved dramatically by several modifications to the synthetic conditions, for both core and shell syntheses, and by post-treatment of the core/shell NPs with TOP. Nearly single-band-edge emission with a PL QY as high as 56% was obtained by the modified synthetic procedures. These findings indicate that the core/shell AgInS2 NPs are nearly comparable to cadmium chalcogenide QDs. The elucidation of the band-edge emission mechanism is underway, with the goal of further improving the PL QY and monochromaticity.
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This work was supported by the MEXT Photonics Advance Research Center Program (Osaka University); by Grant-in-Aid for Scientific Research (A) Grant Number JP18H03863; and by JSPS KAKENHI Grant Numbers JP26107014, JP17H05254 in Scientific Research on Innovative Areas “Photosynergetics”, and JP16H06507 in Scientific Research on Innovative Areas “Nano-Material Optical-Manipulation”, JSPS Advanced Nano Photonics Research and Education Center in Asia. This work was also supported by the Nichia Corporation. Part of this work (HRTEM and STEM observations) was supported by MEXT “Nanotechnology Platform” (project No. 12024046). The authors thank Professor Satoshi Minakata and Professor Youhei Takeda (Osaka University) for the PL decay measurements.
Conflict of interest
Two of the research groups represented by S.K. (Osaka University) and T.T. (Nagoya University) have been funded by the Nichia Corporation. The other authors declare that they have no conflict of interest.
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Uematsu, T., Wajima, K., Sharma, D. et al. Narrow band-edge photoluminescence from AgInS2 semiconductor nanoparticles by the formation of amorphous III–VI semiconductor shells. NPG Asia Mater 10, 713–726 (2018). https://doi.org/10.1038/s41427-018-0067-9
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