Original CoreâShell Structure of Cubic CsPbBr3@Amorphous

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An Original Core-Shell Structure of Cubic CsPbBr3@amorphous CsPbBrx Perovskite Quantum Dots with High Blue Photoluminescence Quantum Yield over 80% Shixun Wang, Chenghao Bi, Jifeng Yuan, Linxing Zhang, and Jianjun Tian ACS Energy Lett., Just Accepted Manuscript • DOI: 10.1021/acsenergylett.7b01243 • Publication Date (Web): 28 Dec 2017 Downloaded from http://pubs.acs.org on December 28, 2017

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ACS Energy Letters

An Original Core-Shell Structure of Cubic CsPbBr3@amorphous CsPbBrx Perovskite Quantum Dots with High Blue Photoluminescence Quantum Yield over 80% Shixun Wang+, Chenghao Bi+, Jifeng Yuan, Linxing Zhang, Jianjun Tian* + These authors contributed equally to this work and should be considered co-first authors Institute for Advanced Materials and Technology, University of Science and Technology Beijing, 100083, China. *Corresponding authors: [email protected]

Abstract All-inorganic perovskite cesium lead halide quantum dots (QDs) have been widely investigated as promising materials for optoelectronic application, because of its outstanding photoluminescence (PL) properties and benefits from quantum effects. Although QDs with fullspectra visible emission have been synthesized for years, the PL quantum yield (PLQY) of pure blue emitting QDs still stays at low level in contrast to their green or red emitting counterparts. Herein, we obtained core-shell structured cubic CsPbBr3@amorphous CsPbBrx (A-CsPbBrx) perovskite QDs via a facile hot injection method and centrifugation process. The core/shell structure QDs showed a record blue emission PLQY of 84%, which is much higher than that of blue emitting cubic CsPbBr3 QDs and CsPbBrxCl3-x QDs. Furthermore, a blue emitting QDsassisted-LED with bright pure blue emission was prepared and illustrated the core-shell QDs’ promising prospect in optoelectrical application. TOC Graphic

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In the last few years, the newly emerging metal halide perovskite materials, especially organometallic halide perovskite (CH3NH3PbX3, X = I, Br and Cl), have drawn a great scientific attention in the field of photovoltaics1, light emitting diodes (LED)2,3, photodetectors4 and lasers5. Hitherto, dozens of studies on the new type of quantum dots (QDs) have been concentrated on the optimization of chemical composition and the method to improve crystal formation for different applications due to the unique physicochemical properties of organometallic

halide

perovskite,

the

superior

photoelectric

performance

and

high

photoluminescence quantum yield (PLQY).6,7 However, organometallic halide perovskite QDs are extremely vulnerable when they are exposed to several factors such as moisture, oxygen, temperature and light,8 thereby triggering a recent research interest in the optical applications of all-inorganic perovskite cesium lead halide QDs which have shown excellent properties including strong photoluminescence (PL) properties, tunable emission over the entire visible spectrum, efficient narrow band emission, negligible influence of self-absorption and Fröster resonance energy transfer as well as the facile synthesis process.9-13 The research for inorganic cesium lead halide perovskite QDs was initiated with adjusting halogen in the perovskite structure to achieve QDs with full-spectral visible emission. QDs with blue emission have attracted more attention among them because of their considerable potential for light-emitting application such as display device.14-16 Through indefatigable effort, researchers have gained green and red emitting CsPbX3 (X represents a halogen) nanocrystals, exhibiting superior thermal stability and high PLQY reaching 90% in solution.10 The highest PLQY of CsPbI3 QDs with red emission even reached 100% reported by Shen’s group recently.17 However, blue emitting CsPbX3 nanocrystals haven’t obtained a comparable value (emission peak: 450 ~ 470 nm, PLQY < 51%)18, 19. Besides, the mixed-halide CsPbBrxCl3-x nanocrystals with blue emission just obtained a lower PLQY less than 32% reported by several groups.20,21 It mainly attributes to their greater optical band gap for blue emission and lots of surface defects thereby inherently inducing the nonradiative recombination of holes and electrons. Up to the present day, the PLQY of those blue emitting perovskite nanocrystals or QDs still pretty low. Here, we successfully incorporated perovskite CsPbBr3 QDs with amorphous CsPbBrx (ACsPbBrx) shell, as shown in Scheme 1, through a facile hot injection method and centrifugation process. It is note that the byproduct of the novel core-shell QDs, existing in the substrate 2


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solution after centrifugation, is CsPbBr3 perovskite nanoplate (cubic crystal structure) with bright green emission (Figure S1, Supporting information). The highest PLQY of blue emitting CsPbBr3@A-CsPbBrx QDs dissolved in n-hexane solution is up to 84%. However, the blue emitting perovskite CsPbBr3 QDs synthesized through the conventional process, which lack an amorphous shell and have a similar PL emission peak, just show a PLQY of 54%. The high PL performance of CsPbBr3@A-CsPbBrx QDs is ascribed to its unique core-shell structure. The amorphous shell runs a protective strategy for perovskite CsPbBr3 core through enclosing its surface and protecting it against radiative corrosion to enhance the ability of the core crystal, and contributes to the formation of excitons as well, which results in the improvement of PLQY. Besides, the novel QDs are relatively free from the influence of lattice miss match between core and shell than multilayer crystal QDs, thereby enhancing the PL efficiency of the original coreshell structure.22 The A-CsPbBrx shell can also provide additional protection to prevent oxygen and water from diffusing into perovskite CsPbBr3 core.

Scheme 1. (a) diagrammatic sketch of all-inorganic perovskite CsPbBr3@A-CsPbBrx composite; (b) structural characterization of perovskite CsPbBr3.

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Figure 1. TEM images and corresponding HRTEM figures of CsPbBr3@ A-CsPbBrx QDs (a, b) originated from the supernatant solution, CsPbBr3 nanosheets (c) obtained from the substrate solution, and blue emitting perovskite CsPbBr3 QDs (d); (e) XRD pattern, (f) optical absorption, and (g) PL emission spectra of the CsPbBr3 QDs and CsPbBr3@A-CsPbBrx QDs. The insets show the emission images of QDs under the excitation of a 365 nm PL emitting UV lamp in ambient environment. Figure 1(a) shows the typical transmission electronic microscopy (TEM) image of CsPbBr3@ACsPbBrx QDs with an average diameter of 6 nm, which presents as flower cluster due to their amorphous structural surface. The high resolution transmission electron microscopy (HRTEM), as shown in Figure 1(b), depicts the core-shell structure of CsPbBr3@A-CsPbBrx QDs. It is plainly visible that the core possesses a well-defined crystalline structure with a characteristic lattice plane distance of about 0.29 nm, corresponding to the d-spacing of (200) crystal planes of cubic CsPbBr3, while the shell of it lacks crystalline structural characteristics. Notably, most of the core-shell QDs in spotlight were covered by the amorphous shell and only a small proportion had exposed crystalline cores. In addition, there is a sort of byproduct (CsPbBr3 nanosheets, Figure 1(c)) in the substrate solution after precipitation while CsPbBr3@A-CsPbBrx QDs existed in the supernatant solution. Figure S1(a) is used to illustrate this point. As the contrastive sample of core-shell QDs, the blue emitting perovskite CsPbBr3 QDs in Figure 1 (d) reveal a characteristic lattice plane distance of 0.18 nm corresponding to the (310) crystal facets of cubic CsPbBr3. It is noted that QDs in Figure 1 (b) are from the supernatant solution, whose substrate 4


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solution contains CsPbBr3 nanoplate (Figure 1(c)) with green PL. QDs in Figure 1(d) synthesized through the conventional method are used as the contrastive samples of the coreshell structured QDs in Figure 1(b). They both are synthesized through a similar hot injection method except different concentrations of precursors and heating process. Characterization was carried out with the help of the X-ray diffraction (XRD) method (Figure 1 (e)). The characteristic and primary diffraction peaks at 12.16° and 30.63° can be attributed to the diffractions from (100) and (200) planes of crystalline cubic CsPbBr3, respectively, which is consistent with the diffraction patterns of its bulk material obtained from PDF#54-752 database. The peak broadened around 20° of CsPbBr3 QDs can be attributed to their small size with average diameter of 2 nm, leading to serious lattice diffusion. Remarkably, the emission peaks of CsPbBr3@A-CsPbBrx QDs have distinct deviations toward a higher degree in comparison with CsPbBr3 QDs, thereby indicating the core’s smaller lattice constant and larger crystal size as well as the strain in the lattice.7 However, after centrifugation, CsPbBr3@A-CsPbBrx QDs existed in the supernatant solution, while CsPbBr3 QDs are extracted from the substrate solution. The main reason is that the CsPbBr3 QDs have larger crystal size than the core of CsPbBr3@A- CsPbBrx QDs, in line with the HRTEM figure above. To some extent, the aberration proved the existence of the amorphous shell which contributed to the peak deviations in the XRD figure. The ultraviolet-visible absorption (UV) and steady-state PL spectra were implemented at room temperature (RT) to further study their properties. So, a 365 nm PL emitting UV lamp was used as the excitation due to their considerable PL emission intensity (Figure S2). As shown in Figure 1(f), the UV absorption and PL spectrum of perovskite CsPbBr3@A-CsPbBrx QDs reveals an obvious excitation peak at 446 nm (2.78 eV) and a highly symmetric emission peak at 463.4 nm (2.68 eV) with a narrow FWHM of 32.87 nm. According to Figure 1(g), the UV absorption and PL spectra of perovskite CsPbBr3 QDs are dominated by a clear excitation peak at 424 nm (2.92 eV) and a sharp PL emission peak at 449 nm (2.76 eV) with a narrow full width at half-maxima (FWHM) of 31.9 nm, respectively, similar to the previously reported data of CsPbBr3 QDs processed through hot injection.13 The CsPbBr3@A-CsPbBrx QDs have shown a wider absorption area than pure CsPbBr3 QDs due to the amorphous shell, which could be worked as a buffer layer to decrease the energy of the UV light. The enhanced absorption ability would contribute to its radiative recombination and high PLQY, because the consumption of short-wavelength light would definitely fill defects through stimulating electron-hole pairs and 5



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provide an outstanding electron transfer pathway to produce bright PL. The small stokes shift of CsPbBr3@A- CsPbBrx QDs (0.1 eV) demonstrates that its less heat dissipation than that of CsPbBr3 (0.16 eV), and its relatively narrow FWHM reveals the uniform size of CsPbBr3@ACsPbBrx QDs (around 6 nm). Significantly, the PL spectra of perovskite CsPbBr3 QDs is not perfectly symmetrical due to the appearance of crystal defects. However, no sub-bandgap emission commonly associated with the defects in CsPbBr3@A-CsPbBrx QDs is observed in their PL spectra, indicating a pure PL emission originated from CsPbBr3 core instead of amorphous shell. To penetrate into the fundamental causes of this novel core-shell QDs, the composite nhexane solution was dipped on the surface of a slide and went through a XRD characterization during different annealing temperature: RT, 50

, 80

and 110

(Figure 2(a)). As mentioned

above, the original solution of the novel QDs consisted of the cubic-phase perovskite CsPbBr3 and amorphous CsPbBrx. Therefore, we can only find the crystal structure of cubic CsPbBr3 through XRD measurement as the blue line depicted in Figure 1(e). After annealing at 50

, the

sample still lacked an obvious diffraction peak of CsPbBrx. However, it exhibited a green PL instead of a blue PL, which means the grain growth of the cubic core CsPbBr3. When annealed at 80

, the CsPbBr3@A-CsPbBrx QDs exhibited a pure green PL. the diffraction peaks at 11.66°

and 46.90° are clearly observed, indicating the diffraction of (002) and (206) planes of crystalline CsPb2Br5, respectively. It means the crystallization of amorphous CsPbBrx and the formation of a new structure in form of mixed CsPbBr3/CsPb2Br5 nanocrystals. It is similar to the former reported green emitting dual-phase CsPbBr3-CsPb2Br5 QDs, which have the CsPb2Br5 phase in the crystal lattice of cubic CsPbBr3 because of metastable state in the cubic phase, nonstoichiometric material transfer or structural rearrangement.23 When the annealing temperature was risen to 110

, the formed mixed CsPbBr3/CsPb2Br5 nanocrystals became

blindingly obvious with the enhancement at 11.66°, 46.90° and 33.37° attributed to the diffraction of (210) plane of crystalline CsPb2Br5, as well as the uprush at 21.57° owning to the diffraction of (110) plane of cubic CsPbBr3. The result further implies the potential existence of the original core-shell structure of CsPbBr3@amorphous CsPbBrx QDs. In order to further study the annealing process, the TEM measurement was conducted to observe samples undertaking different annealing temperature, as shown in Figure 2(b-e). Figure 2(b) depicts the core-shell structure of CsPbBr3@A-CsPbBrx QDs and a well-defined crystalline structure with a 6


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characteristic lattice plane distance of about 0.18 nm, corresponding to the d-spacing of (310) crystal planes of cubic CsPbBr3 (white color). As shown in Figure 2(c), the core-shell QDs have grown into nanoplates after annealing at 50

. It is clearly visible that the center of the nanoplate

possesses a characteristic lattice plane distance of about 0.47 nm corresponding to the d-spacing of (102) crystal planes of CsPb2Br5 (yellow color). However, the peripheral vision shows an obvious characteristic lattice plane distance of about 0.58 nm corresponding to the d-spacing of (100) crystal planes of CsPbBr3 (white color). Figure 2 (d) and (e) represent the TEM figures of CsPbBr3@A-CsPbBrx QDs annealed at 80

and 110

, respectively. Their electron diffraction

diagrams show a distinctive characterization of cubic CsPbBr3. To further prove the growth procedure of the core-shell QDs during annealing process, the TEM measurements of samples annealed at 30 to 40

were carried out (as shown in Figure S3). It is obvious that the CsPbBr3

core would swallow the amorphous shell and form the nanoplates as shown in Figure 2(c) at the same time. Herein, it can be deduced that the cubic CsPbBr3 core would quickly grow into CsPbBr3 nanoplate while the amorphous CsPbBrx shell need more energy to form crystalline CsPb2Br5. The growth of the shell is more slowly compared to that of the CsPbBr3 core. The above results satisfied the Raman measurement (Figure 3(a)) which was conducted in nhexane solution. Here a longer laser excitation (633 nm) and a shorter laser excitation (532 nm) were applied to study the internal and surface of CsPbBr3@A-CsPbBrx QDs and pure CsPbBr3 QDs, respectively. When using the 633 nm light as excitation, both QDs revealed a strong vibrational model ν1 at 72 cm-1 indicating the existence of [PbBr6]4- octahedron at the center of QD samples based on the relevant literatures.24, 25 However, CsPbBr3@A-CsPbBrx QDs have no feature of ν1 compared to CsPbBr3 QDs when using a 532 nm laser excitation, which means the lack of crystalline CsPbBr3 phase in the surface of CsPbBr3@A-CsPbBrx QDs due to the existence of amorphous CsPbBrx. Therefore, the formation of CsPbBr3@A-CsPbBrx QDs can be reasonable deduced: a large number of cubic CsPbBr3 QDs nucleated at the first stage of hot injection process, then the QDs near the heating center (substrate of the three-necked bottle) formed CsPbBr3 nanoplate with green PL due to abundant precursors and enough energy for growth. On the contrary, the QDs in the supernatant were lack of sufficient energy to crystallization. So, the atomic layer of small cubic CsPbBr3 QDs forms amorphous CsPbBrx layer. To study the effects of the amorphous CsPbBrx layer, PL spectra of CsPbBr3@CsPbBrx QDs and CsPbBr3 QDs recorded under different light excitation wavelengths ranging from 365 7



to 400 nm were conducted as shown in Figure 3(b). Though the light excitation has little influence on their PL emission peak position, it still reveals the two samples’ distinct varying pattern of PL intensity due to their structural difference. The PL intensity of CsPbBr3 QDs gradually decrease with the reduction of light excitation, while the value of CsPbBr3@CsPbBrx QDs remains at a high emission level until the light excitation is lower than 3.1 eV (400 nm), indicating that the amorphous CsPbBrx shell can optimize the core’s absorption ability under certain excitation and make it easier to reach its absorption saturation through the process of changing light excitation from high energy to lower energy.

Figure 2. (a) XRD pattern of CsPbBr3@A-CsPbBrx QDs during different annealing process under ambient environment. The illustrated pictures in the right side represent the corresponding PL phenomenon under UV light; TEM images of CsPbBr3@A-CsPbBrx QDs processed under different annealing temperature: (b) RT, (c) 50 , (d) 80 and (e) 110 . The insets of (c) and (d, e) are the corresponding HRTEM figure and electron diffraction diagrams.

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Figure 3. (a) Raman spectra of CsPbBr3@A-CsPbBrx and CsPbBr3 QDs under the excitation of 633 nm and 532 nm PL emitting lamps respectively; (b) PL spectra of CsPbBr3@A-CsPbBrx QDs and CsPbBr3 QDs recorded under diverse light excitations ranging from 365 to 400 nm (the insets are their enlarged vision of emission peaks); PL emission spectra of (c) CsPbBr3 QDs and (d) CsPbBr3@A-CsPbBrx QDs under the excitation of a 365 nm PL emitting UV lamp with different excitation intensity. 9



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Table 1. PLQY of CsPbBr3@A-CsPbBrx QDs and pure CsPbBr3 QDs under the excitation of a 365 nm PL emitting UV lamp with different excitation intensities ( is the absorbance in UV spectra, is the integral absorption area in PL spectrum). Rhodamine B

Voltage Strength

CsPbBr3

300V 350V 400V 450V

13.0 45.6 127.9 164.2

0.01122

42.1 152.0 448.4 510.9

0.05293

PLQY (%) 48.57 49.94 52.52 46.62

CsPbBr3@A-CsPbBrx PLQY (%) 74 82.29 257.4 81.60 0.05485 740.8 83.73 952.3 83.84

With a referenced rhodamine B, the relative PLQY of QDs can be calculated as function (1) to evaluate their PL performance: =

where

⋅ ⋅

⋅

⋅

(1)

is refractivity, is the absorbance in UV spectra, is the integral absorption area in PL

spectrum. The absorption factor must be accurately determined for solutions with a low absorbance, typically A < 0.1, in order to avoid internal filter effects and errors arising from uneven distribution of the excited species in the detected volume.26 Thus, the two samples were all evenly distributed in n-hexane solution with absorption intensity far less than 0.1 at 365 nm to get a more precise relative PLQY. By calculation the result obtained: the PLQY of CsPbBr3@ACsPbBrx QDs (84%) is much higher than that of CsPbBr3 QDs solution (54%) under excitation of 365 nm PL emitting light. The influence of excitation intensity on PLQY has also been researched. As depicted in Figure 3(c, d), CsPbBr3@A-CsPbBrx QDs have a wider FWHM (around 31.88 nm) than that of CsPbBr3 QD (around 22.13 nm) due to their simple centrifugation process leading to larger size distribution and the favorable crystalline of CsPbBr3 QDs sample, though their FWHM is already low enough to produce pure blue emission. The red shift (< 3 nm) in Figure 3(d) may attributed to the growth of CsPbBr3 core during continuous radiation or reasonable measurement error which needs further systemic studies. However, the process didn’t weak it PL ability as can be seen in Table 1. The default excitation intensity of PL measurement is the value when a 400 V voltage power putted on a Xe lamp. The voltage strength has been adjusted from 300 to 450 V and the result shown a similar PLQY of CsPbBr3@A-CsPbBrx QDs 10


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(around 83%) which signifies the excitation intensity has little effect on the PLQY of CsPbBr3@A-CsPbBrx QDs. However, the PLQY of CsPbBr3 QDs decreased by around 5.9% when a 450 V voltage power was applied, indicating that the pure crystals’ PL abilities are not stable under a high radiative intensity, though their PL emission position remains stable. Meanwhile, according to the data in Table 1, the CsPbBr3 QDs have much less integral absorption area in PL spectrum than that of CsPbBr3@A-CsPbBrx QDs and obtain a lower PLQY (< 53%), though the two sample have similar absorbance at 365 nm (about 0.05 a.u.). In other words, the increase of the PLQY of CsPbBr3@A-CsPbBrx QDs is mainly because of the increase of , which means more photos would be excited. Thus the core-shell structural QDs have a higher ratio between the emitted photos and absorbed photos. Briefly, the CsPbBr3@A-CsPbBrx QDs have a better PL ability due to the core-shell system’s contribution to the formation of excitons.

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Figure 4. Time-resolved PL decay curves of CsPbBr3 QDs and CsPbBr3@A-CsPbBrx QDs tested (a) under various slit widths (2.5 nm, 5 nm and 10 nm), or (b) during different test times; (c) typical time-resolved PL decay curves of CsPbBr3 QDs and CsPbBr3@A-CsPbBrx QDs with the corresponding lifetimes shown in the inset The fitted curves are depicted in a black color for CsPbBr3 QDs and a red color for CsPbBr3@A-CsPbBrx QDs; (d) CIE chromaticity coordinates of the LED based on CsPbBr3@A-CsPbBrx QDs (circle) compared to the NTSC color standards (hollow circle). The inset in bottom-right indicates the LED fabricated with CsPbBr3@ACsPbBrx QDs working on the voltage of 2 V.

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Table 2. Time-resolved PL decay data of CsPbBr3 QDs and CsPbBr3@A-CsPbBrx QDs tested under different slit widths.

CsPbBr3 CsPbBr3 @ACsPbBrx

Slit Width/nm 2.5 5 10 2.5 5 10

!"

!#

!$

%&'()"

%&'()#

%&'()$

*#

!+.

1.43 1.65 1.53 1.74 1.71 1.69

5.42 4.83 5.29 5.38 5.52 5.56

24.09 21.31 20.55 23.35 22.13 20.52

9.97 12.36 8.27 9.19 10.52 8.93

85.68 83.00 86.18 86.07 85.11 86.04

4.36 4.64 5.56 4.74 4.37 5.03

1.12 1.14 1.07 1.02 1.07 1.04

7.23 6.68 7.02 7.15 7.25 7.22

To draw out the carrier recombination dynamics of the two different kinds of inorganic QDs, the time resolved PL measurement should be applied. The PL decay curves can be well-fitted with a biexponential function (2)9, 18: 6

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Original CoreâShell Structure of Cubic CsPbBr3@Amorphous

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