Flexible light-emitting diodes (LEDs) are thin, lightweight, and deformable, attracting considerable interest for their potential applications in wearable technology, flexible displays, and lighting [1–4]. Colloidal quantum dots (QDs) offer a distinctive class of emitters with high efficiency, a wide color gamut, stability, spectral tunability, and low fabrication costs [5–8]. QD-based LEDs (QLEDs) have emerged as a promising option for flexible and ultrathin displays due to their outstanding performance on different substrates [9–11]. Recently, various strategies have been employed to enhance the performance of QLEDs. These strategies include refining the nanostructure of QDs to boost hole injection [12,13], optimizing the design of ZnO-based electron transport layers for effective electron injection [14,15], and reducing device interfacial potential while passivating interfacial defects to maintain a balance in carrier injection and prevent interfacial exciton quenching [16,17]. These efforts have enabled red, green, and blue QLEDs with maximum external quantum efficiencies (EQEs) of 35.6%, 28.7%, and 23%, respectively [18–20], and long operating lifetimes of 3,300,000 h for red [21], 580,000 h for green [18], and 41,000 h for blue colors [20] have already satisfied the requirements for displays. Meanwhile, by optimizing the transparent electrodes and improving the light out-coupling efficiencies, the EQEs of flexible QLEDs for red [22], green [23], and blue devices [24] have reached 24.1%, 9.88%, and 4.5%, respectively, with a maximum lifetime of 50,000 h [25]. However, the electroluminescence (EL) efficiency and operating lifetime of state-of-the-art flexible QLEDs are still lower than those of rigid substrate devices. These facts necessitate the exploration of novel approaches to enhance the performance of flexible QLEDs.

Type-I nanostructure QDs (typical InP-based and CdSe-based QDs) have been successfully employed in flexible QLEDs [23,26]. However, the limitation in device performance is mainly because the repeated bending creates cracks in the active layer. This process hybridizes defects from the transport layer into the QD layer, resulting in increased operating current and accelerated device degradation. This is due to the weak confinement effect causing excitons to leak to the QD surface and then be trapped by defects, resulting in nonradiative recombination and decreasing the QD’s emission efficiency [27–30]. Moreover, the exciton wavefunction of type-I QDs tends to delocalize from the inner core to the outer shell. Hence, enhanced QD shell energy can confine the exciton wavefunction from escaping to the QD surface [29,31,32], effectively avoiding the decrease in photoluminescence quantum yields (PLQYs) due to surface contact defects [33,34]. Typically, in CdSe-based QDs, ZnS as the shell material (CdSe/ZnS QDs) is believed to offer enhanced exciton confinement characteristics, although this is limited by the low PLQY. Subsequently, the lattice fitness CdS and ZnSe as shell were designed for efficient CdSe/CdS [27], CdSe/ZnSe [12], and CdSe/gradient alloy shell/ZnS QDs [13,23,35], respectively [36–38]. Such QDs with type-I structures in flexible QLEDs are more favorable for the exciton delocalization, thus suffering from contact defect-induced efficiency degradation. These problems significantly reduce the operational stability of flexible QLEDs, impeding the commercialization of full-color QLEDs displays.

Here, we design, fabricate, and characterize high-performance red flexible QLEDs by manipulating the QD shell arrangement to confine the exciton wavefunction. The designed anti-type-I ZnCdSe/ZnS/ZnSe QDs (abbreviated as A-QDs) exhibit excellent performance in suppressing defect-induced QD degradation, enhancing exciton radiative recombination, and balancing charge injection in the device. The wide bandgap ZnS intermediate shell QDs show high PLQY and reliability comparable to that of ZnCdSe/ZnSe/ZnS QDs (C-QDs). When used as emitters, solution-processed flexible QLEDs have a maximum external quantum efficiency (EQE) of 23.0%, a high brightness of 108,100cdm−2, and long operational lifetimes of 63,050 h at T50@100cdm−2. The device with custom-shelled QDs reduces electron injection and enhances hole injection to balance charge injection, suppressing defect-induced nonradiative recombination during operation, resulting in highly efficient and stable flexible QLEDs. This study provides new insights into achieving superior performance of flexible QLEDs, a critical step toward realizing commercial flexible displays.

To confine the exciton wavefunction delocalization of core/shell QDs typically requires a strong confinement wide bandgap shell wrapped around the core emitter [Fig. 1(a)]. From the lattice constants, the ZnCdSe cores are promising as a candidate material to match the lattice constants of ZnS intermediate shell [19,38]. The ZnS shell wrapped around the ZnCdSe to form ZnCdSe/ZnS, and then the following ZnSe outer shell was grown to form the final anti-type-I ZnCdSe/ZnS/ZnSe QDs (A-QDs). Meanwhile, the same methods are applied to synthesize conventional type-I ZnCdSe/ZnSe/ZnS QDs (C-QDs) by continuously growing the ZnSe shell and ZnS outer shell on the initial ZnCdSe core. Both QDs have shown excellent PLQY of more than 90%, and corresponding emission peak and FWHM are 618 and 28 nm for A-QDs and 622 and 28 nm for C-QDs, respectively.

The band alignment diagrams of C-QDs and A-QDs heterojunctions are shown in Figs. 1(b) and 1(c). Based on the energy offsets at the C-QDs and A-QDs heterojunctions, the ZnS intermediate shell modulates the extent of electron and hole wavefunction delocalization [31]. The calculated spatial probability distributions of the hole and electron as a function of QDs radius are shown in Figs. 1(d) and 1(e). In contrast to C-QDs, the electron wavefunctions of A-QDs are confined within the intermediate shell. Due to the wide-bandgap ZnS intermediate shell, A-QDs exhibit strong electron-hole wavefunction confinement. The calculated electron and hole wavefunctions’ spatial areas in the core are 69% and 98% for A-QDs and 46% and 92% for C-QDs, respectively. These findings suggest that electron and hole wavefunctions can cross the potential barriers of type-I shells (C-QDs), while being significantly confined in anti-type-I shells (A-QDs).

In the fabrication of QLEDs, the typical structure of the electron transport layer/QD/hole transport layer (ETL/QD/HTL) makes the delocalized exciton formed by the C-QD susceptible to trapping by defect states at the heterointerfaces, which reduces the PLQY of the emitting layer [Fig. 1(b)]. In contrast, the strong exciton confinement effect of A-QDs effectively confines the exciton wavefunction within the inner core, thus efficiently suppressing contact defect induced nonradiative recombination [Fig. 1(c)]. Furthermore, the ZnSe outer shell of the A-QD is able to reduce the hole injection barrier (0.4 eV) and suppress electron injection in the device more effectively than the ZnS outer shell layer (C-QDs) [12,39], resulting in a more balanced charge injection. However, the type-I energy barrier of C-QDs reduces the electron-hole wavefunction overlap and promotes the electron delocalization to its surface, which increases the exciton sensitivity of the heterojunction films.

Aberration-corrected high-angle annular dark-field scanning transmission electron microscopy (AC-HAADF-STEM) images [Fig. 2(a)] show the QDs with uniform size (13.1 nm for A-QDs and 13.0 nm for C-QDs, respectively; see Fig. 6 in Appendix A). The simulated energy-dispersive X-ray spectroscopy (EDS) elemental mapping confirmed their composition and core/shell structures. Furthermore, the distribution of S and Cd elements in the A-QDs is definitively confirmed by AC-HAADF-STEM and EDS, aligning well with the nanostructure design [Fig. 2(b), and see Fig. 7 in Appendix A].

To examine the Auger recombination rate of these QDs, which is proportional to materials’ exciton binding energy (Eb), reducing the corresponding materials’ Eb can suppress Auger recombination [40]. We conducted temperature-dependent PL spectra and PL decay lifetime measurements of QD films at 300–430 K to quantitatively extract the Eb and emission states, respectively [Figs. 2(c) and 2(d)]. As shown in Fig. 2(c), the extracted Eb is estimated to be 66.4 and 49.2 meV for C-QDs and A-QDs, respectively. The results reconfirmed that the wide-bandgap ZnS intermediate shell decreases the Eb, due to the enhanced quantum confinement effect. Moreover, when the temperature is increased from 300 K to 430 K, the PL peak of A-QDs is redshifted by 6 nm less than that of C-QDs by 16 nm, suggesting that QDs with anti-type-I nanostructures have stronger quantum confinement, limiting the temperature-induced exciton delocalization. We extracted the QDs radiative decay channels and efficiencies at different temperatures by using the time-resolved PL (TRPL) spectra measurement [Fig. 2(d)]. When the temperature rises from 300 K to 400 K, the average lifetime decreases from 15.6 ns to 15.1 ns for A-QDs and 15.8 ns to 14.2 ns for C-QDs, while the corresponding radiative efficiencies decrease by 3.2% and 10%, respectively [33,41]. These differences indicate that the ZnS intermediate shell in the A-QDs effectively inhibits the thermally induced exciton quenching.

We further examine PLQY values and PL lifetimes of the A-QDs and C-QDs films to clarify their PL mechanism and emission states [Figs. 2(e) and 2(f)]. The A-QDs exhibit a PLQY and an average PL lifetime of 92% and 18.5 ns in solution, 85% and 15.8 ns on glass, 78% and 14.5 ns with Cl-treated ZnMgO ETL, and 76% and 13.3 ns on ETL, respectively. In contrast, the C-QDs show a PLQY and an average PL lifetime of 91% and 19.1 ns in solution, 78% and 11.2 ns on glass, 71% and 10.9 ns with Cl-treated ETL, and 66% and 9.5 ns on ETL, respectively. Notably, the PLQYs of the A-QDs films are higher than those of the C-QDs films under the same conditions [17,42]. The Cl-treated ZnMgO can effectively passivate the interface defects between the ETL and the QD layer, which is conducive to further improving the emission efficiency of the QDs. Combining the above findings, we can conclude that the excitons in C-QDs can be quickly separated and migrate to their surface due to the weak exciton confinement and could be captured by surface trap states, resulting in an increased nonradiative recombination rate. This result indicates that energy transfer (ET) influences the PLQY in a closely-packed QD film, accelerating the nonradiative recombination of excitons. The ET efficiency decreased for the A-QDs compared to the C-QDs when using the ZnS intermediate shell (ET efficiencies of 27% for C-QDs and 15% for A-QDs, respectively) [33], indicating reduced PL quenching for the A-QDs in film. Moreover, the ET efficiencies of QDs obtained from TRPL measurements are consistent with theoretical calculations; specifically, exciton delocalization to the surface of the QDs leads to improved ET. These results indicate A-QDs with better preservation of optical properties with a ZnS intermediate shell.

Modifying the nanostructure of QDs can influence their chemical state and energy level arrangement [43,44]. The chemical states of the QDs were investigated using X-ray photoelectron spectroscopy (XPS). As shown in Fig. 3(a), we extracted the binding energy from the high-resolution Zn 2p and Cd 3d spectra of the A-QDs, showing the peaks at 1045.2 eV (1022.1 eV) and 411.8 eV (405.1 eV) for Zn 2p1/2 (Zn 2p3/2) and Cd 3d3/2 (Cd 3d5/2), respectively. A comparison with the C-QDs reveals a slight increase in binding energy of approximately 0.3 eV for Zn 2p and 0.2 eV for Cd 3d in the A-QDs. A similar positive shift is also observed in the Se 3d and S 2p regions [Fig. 3(b), 0.1 eV for Se and 0.7 eV for S, respectively]. These findings suggest that the valence electrons of the A-QDs are more tightly bound than those of the C-QDs, indicating stronger electronic interaction. Further, we characterize the QD energy levels performed by ultraviolet photoelectron spectroscopy (UPS) measurements. As shown in Fig. 3(c), the valence band of the A-QDs was calculated at 5.8 eV, which is higher than that of C-QDs (6.2 eV). The UPS data demonstrate that the modulation nanostructure of QDs indeed causes a 0.4 eV upshift of the energy levels of A-QDs compared to C-QDs. This upshift boosts the hole current from the HTL to QDs and simultaneously decreases the electron overflow from the ETL into QDs, resulting in an improved balance between electron and hole injection.

The anti-type-I nanostructure design of A-QDs potentially facilitates hole injection while reducing electron injection. Therefore, we extracted the carrier mobility and trap state density of these QD films by using the space-charge-limited current (SCLC) technique [17]. The SCLC data show that the hole mobility and defect state density of C-QDs and A-QDs films are 9.37×10−5cm2V−1s−1 and 1.92×1016cm−3, and 1.46×10−4cm2V−1s−1 and 5.49×1015cm−3, respectively [Fig. 3(d)], suggesting that the A-QDs have higher carrier mobility and lower trap state density [45]. We compare the electron and hole transport into A-QDs and C-QDs films from neighboring charge transport layers in hole-only devices (HODs) and electron-only devices (EODs), respectively [Figs. 3(e) and 3(f)]. Notably, the current density–voltage characteristics of these devices show significant differences. The A-QDs-based single-carrier devices exhibit higher hole transport current and lower electron transport current. These results indicate that the A-QDs with modulating shell landscape confine the QDs exciton delocalization, which not only upshifts the electronic energy level of the QD layer but also promotes the charge injection balance of the QDs in the device (see Fig. 8 in Appendix A).

The flexible QLEDs with the configuration of indium tin oxide (ITO)/ZnMgO-Cl/QD/TFB/MoO3/Al were fabricated on a thin flexible PET substrate [Figs. 4(a) and 4(b)]. We utilized the inverted device structure of flexible QLEDs to integrate A-QDs and C-QDs between the HTL and ETL, respectively, achieving a significant improvement in the device performance. The current density and luminance versus voltage characteristics of the devices are presented in Fig. 4(c). Benefiting from the reduced hole injection barrier, the A-QDs device has a low turn-on (VT) voltage of 1.6 V, smaller than the C-QDs of 1.8 V. The flexible QLEDs with A-QDs and C-QDs showed a brightness of 108,100cdm−2 and 56,100cdm−2 and EQEs of 23.0% and 14.7%, respectively [Fig. 4(d)]. The injected current in A-QDs devices promotes exciton radiative recombination more effectively than in C-QDs devices, enhancing the electroluminescence (EL) efficiency of these devices. Accordingly, for the EL spectrum with an emission peak at 626 nm and a full width at half-maximum (FWHM) of 33 nm, the EL photographs show uniform and bright red emissions [2mm×3mm; Fig. 4(e)]. Furthermore, as the voltage increases, the brightness of the red flexible QLED photographs continues to increase (see Fig. 9 in Appendix A). Notably, a peak EQE of 23.0% is achieved for the A-QDs device, which represents one of the most efficient red flexible QLEDs reported to date.

The lifetime (T50) under different initial luminance after 100 bending cycles fits the acceleration factor (n) according to the empirical equation L0n·T50=constant [Fig. 4(f)] [40]. The flexible QLED with A-QDs showed a superior operational lifetime, and its T50 at 100cdm−2 is predicted to be 63,050 h, using the acceleration factor of 1.8 (given that we measured the lifetime at ∼100,000cdm−2).

The A-QDs devices exhibited a T50 lifetime of more than 60,000 h, which is a superior lifetime to that of C-QDs devices of 4650 h. This is the most stable flexible QLED ever reported (Table 1). Moreover, after nearly 3000 bending cycles of the flexible QLEDs, the brightness of the devices degraded by only 15% (see Fig. 10 in Appendix A). Thus, we confirm that manipulating the exciton wavefunction confinement of QDs, which can efficiently suppress defect-induced nonradiative recombination in the device, plays a crucial role in improving operational stability.Table 1.

Efficiency and Stability of Flexible QLEDs^a

To elucidate the effect of manipulating the exciton confinement of QDs on the charge transfer kinetics and stability of the flexible QLEDs, we first probed the mechanism by using capacitance-voltage (C–V) measurements [Figs. 5(a) and 5(b)] [48]. The C–V test of the devices was performed under a DC bias (VDC, from −1 to 3 V) and a small constant AC modulation voltage (VAC, 20 mV). The geometric capacitance of all devices was maintained at 3.0 nF, suggesting that all devices maintained structural consistency and similar dielectric constants [17]. The capacitance of the C-QDs device increased to a maximum of 3.8 nF when the VDC tuned from 1.4 up to 2.0 V. In terms of a direct comparison, the capacitance increased from 3.0 nF to a maximum of 3.5 nF (VDC increases from 1.4 to 1.6 V) for A-QDs device. A reduction in A-QDs device capacitance indicates that the device effectively avoids carrier accumulation at the heterointerfaces, resulting in a more balanced charge injection. However, after 100 bending operations, the peak capacitance of the C-QDs device increases by approximately 0.6 nF, reaching a maximum of 4.4 nF, whereas no significant difference is observed in the A-QDs device. This suggests that the C-QDs device experiences severe charge accumulation during operation due to the relatively high barrier of the ZnS outer shell.

A bending aging test could cause irreversible structural damage to device nanofilms, leading to mutual intrusion at the contact interface and rapid degradation of device performance. We have investigated the changes in carrier concentration (NCV) and built-in potential (Vbi) in flexible QLEDs by performing Mott–Schottky measurements on both pristine and aged devices. For both the pristine and aged A-QDs devices [Fig. 5(c)], the NCV is from 7.1×1016cm−3 to 8.0×1016cm−3 and the Vbi remains almost constant (from 1.17 V to 1.16 V). However, for C-QDs devices, the NCV increases dramatically from 5.7×1016cm−3 to 1.8×1017cm−3 while the Vbi increases from 1.12 V to 1.28 V. In comparison to A-QDs devices, the notable rise in carrier concentration in C-QDs devices suggests that a large amount of injected charge contributes to a current through nonradiative recombination. Furthermore, the rise in carrier concentration elevates the built-in electric field strength, promoting the separation of excitons. The high current density and low radiative recombination efficiency of the device result in rapid performance degradation.

We further applied thermal admittance spectroscopy (TAS) to probe the trap density of states (tDOS) and energy depth of defect states in the flexible QLEDs [49]. The TAS technique is recognized as an effective method for characterizing both shallow and deep energy depths of defect states. The energetic profile of tDOS can be deduced from the capacitance-frequency plots, where the energetic demarcation E(ω) depends on the trap state depth [40]. The peak tDOS of the A-QDs device is 1.32×1023cm−3eV−1, which is comparable to that of the C-QDs device [Figs. 5(e) and 5(f)]. After bending cycles, the tDOS and efficiency of the A-QDs device remain almost unchanged, while the tDOS of the C-QDs device more than triples over the entire trap depth, and the performance degrades rapidly compared to the pristine state. The changes in tDOS before and after the bending cycles were extracted in TAS analyses, consistent with the conclusions of the aforementioned C-V data, device efficiency, and stability. These results suggest that the QDs with strong exciton confinement of the ZnS intermediate shell can effectively eliminate the device efficiency degradation caused by bending-induced interfacial defects intrusion.

In summary, we demonstrated highly efficient and stable flexible QLEDs by manipulating exciton confinement. A novel design for the customization of red ZnCdSe/ZnS/ZnSe QDs with controlled shell energy alignment confines the exciton wavefunction into the core, improving PL stability and balancing carrier injection in the device. Encouraged by these findings, we finally achieved red flexible QLEDs with a maximum EQE up to 23.0%, brightness of 108,100cdm−2, and outstanding T50@100cdm−2 lifetime of 63,050 h, respectively. The operando characterization confirmed that the flexible A-QDs device can balance charge injection and resist defect-induced performance degradation, thus improving the efficiency and stability of QLEDs. We believe that the present innovative design and advances in manipulating QD exciton confinement represent a crucial step toward the realization of flexible QLEDs displays.

Zinc acetate [Zn(Ac)2, 99.99%], cadmium oxide (CdO, 99.99%), magnesium acetate tetrahydrate [Mg(Ac)2·4H2O, 99%], dimethyl sulfoxide (DMSO, 99.9%), and ethanol (99.9%) were purchased from Aladdin. Selenium (Se, 99.99%), sulfur (S, 99.99%), and tetramethylammonium hydroxide pentahydrate (TMAH, 98%) were purchased from Alfa Aesar. Oleic acid (OA, 90%), 1-octadecene (ODE, 90%), and tri-n-octylphosphine (TOP, 97%) were purchased from Acros. Poly-((9,9-dioctylfluorenyl-2,7-diyl)-co-(4,4′-(N-(4-secbutylphenyl) diphenylamine))) (TFB) was purchased from American Dye Source.

These QDs were prepared according to the previously reported literature with appropriate modifications [19,50,51]. For a typical synthesis of ZnCdSe/ZnS/ZnSe QDs (A-QDs), 0.7 mmol of CdO, 0.9 mmol of Zn(Ac)2, 10 mL of ODE, and 5 mL of OA were mixed in a 100 mL three-necked flask and heated to 150°C under nitrogen flow for 30 min. Then, the temperature was elevated to 310°C. A stock solution containing 1.5 mmol of Se dissolved in 3 mL of TOP was quickly injected into the flask, and the reaction was kept for 60 min to grow ZnCdSe core. After that, 2 mL of zinc precursor (1 mol/L) was added dropwise for 10 min [10 mmol Zn(Ac)2, 5 mL OA, and 5 mL ODE was mixed in a 100 mL flask, heated to 120°C, and exhausted for 10 min, followed by heating up to 310°C and boiling for 20 min to obtain a colorless solution], followed by 2 mmol of S dissolved in 1 mL of TOP being added dropwise into the flask, and the reaction was continued for 20 min to grow a ZnS intermediate shell. After that, 2 mL of zinc precursor (1 mol/L) was added dropwise for 10 min, followed by 2 mmol of Se dissolved in 4 mL of TOP added dropwise into the flask to react for 20 min. After completing the reaction, the temperature was rapidly reduced to room temperature. The resulting QDs were precipitated with acetone several times and finally dispersed in octane. For the synthesis of ZnCdSe/ZnSe/ZnS QDs (C-QDs), a similar procedure was performed except for the different order of adding the S and Se precursors of the growing shell layer.

ZnMgO NPs were synthesized using our previously reported method [17,43,52]. The synthesis process was as follows: 3 mmol of Zn(Ac)2 and 0.36 mmol of Mg(Ac)2·4H2O dissolved in 30 mL of DMSO and 5.5 mmol of TMAH dissolved in 10 mL of ethanol were mixed and stirred for 1 h in air at 22°C. Then, ZnMgO NPs were washed with ethanol/ethyl acetate and centrifuged (8000 r/min, 3 min) and dispersed in ethanol at a concentration of 30 mg/mL and stored at −4°C.

The flexible PET substrates with patterned ITO anodes (∼10Ω/sq) were cleaned sequentially in ultrasonic baths of ethanol, acetone, and isopropanol for 10 min each, and then the cleaned ITO was treated by plasma for 10 min. Then, ZnMgO NPs were spin-coated on the ITO substrates at 3000 r/min for 45 s, followed by baking at 80°C for 30 min in air. After that, a chloroform solution was spin-coated over the ZnMgO layer at 3500 r/min for 45 s and baked at 80°C for 5 min in an Ar-filled glove box. The QDs layers were spin-coated over ZnMgO films at 3000 r/min for 45 s using an 18 mg/mL solution and baking at 80°C for 15 min. Then, TFB layers were spin-coated at 3000 r/min for 45 s using an 8 mg/mL chlorobenzene solution, followed by baking at 80°C for 30 min. Finally, the top electrode layers, consisting of 10 nm of MoO3 and 100 nm of Al, were successively deposited onto the ITO/ZnMgO/QD/TFB structure by a thermal evaporator at a rate of 0.1 and 1 Å s−1, respectively. Then, the device was encapsulated with UV-cured epoxy.

The PL spectra, efficiency, and lifetime data of QDs were measured with an Edinburgh FLS1000 fluorescence spectrophotometer. AC-HAADF-STEM and EDS maps were recorded using an FEI Themis Z (300 kV). XPS data were collected by an ESCALAB 250 XI system with a monochromatic X-ray source (Al Kα source), and UPS data were obtained on an electron spectrometer using monochromatic HeI radiation at 21.22 eV. High-resolution XPS data were collected from the regions around Zn 2p, Cd 3d, Se 3d, and S 2p. The current density–voltage–luminance characteristics were measured and recorded using a Keithley 2400 source unit coupled with a fiber integration sphere and QE65 Pro spectrometer. The electrochemical measurements of the devices were carried out using a VMP-300 impedance analyzer (Biologic).

The exciton spatial distribution was determined by employing a simple model for QDs with a spherical potential well. The single-particle Hamiltonian in the effective mass approximation is given by H^=−12m∇2+V(r), where m is the effective mass and V(r) is the position-dependent potential energy. The time-independent Schrödinger equation is Hψ=Eψ, where E is the corresponding eigenenergy. In the calculations, the conduction band (CB) and valence band (VB) energy levels of bulk materials are extracted from a previous report [53]. The electronic and hole ground state energies are calculated for both ZnCdSe/ZnS/ZnSe and ZnCdSe/ZnSe/ZnS QDs with different shells. The effective masses of electrons (me) and holes (mh) are set according to literature reports as 0.12 and 0.57 for CdSe, 0.14 and 0.82 for ZnSe, and 0.34 and 1.78 for ZnS, respectively [18,54].

The trap state density (Ntraps) values are extracted from the current density–voltage profiles by the SCLC model, determined by the equation [40] Ntraps=2εε0VTFLeL2, where VTFL is the trap-filled limit voltage, e is the elementary charge, L is the film thickness, ε is the relative dielectric constant, and ε0 is the vacuum permittivity. The carrier mobility (μ, electron mobility) is extracted from the child region and determined by the equation [40] μ=8JL39εε0V2, where J is the device’s current density.

Thermal admittance spectra were measured by an impedance analyzer (VMP-300 impedance analyzer, 100 Hz to 7 MHz) at room temperature. We first tested the electrochemical impedance spectra of the devices by using a 20 mV AC signal to measure the capacitance of the devices as a function of AC signal frequency (C-f, 100 Hz to 7 MHz). The device is then tested in Mott–Schottky mode with an applied DC voltage of −1 to 3 V and an applied AC voltage of 20 mV (1000 Hz). Extraction of capacitance-voltage (C-V) curves comes from the Mott–Schottky test mode. The correlation between the DOS of trap charges indicates the frequency-dependent capacitance that can be expressed as Nt(Ew)=−Vbi2w[qVbi−(EF−Ew)]∂C∂wwkT, where Nt is the DOS of trap charges, Vbi is the built-in potential, L is the depletion width of the devices, EF is the Femi level, Ew is the energy level of the trap state relative to the valence band, k is the Boltzmann constant, and ∂C∂w is the partial derivative to the electrochemical impedance spectrum C-f [48,49].