*Article* **Numerical Study to Enhance the Sensitivity of a Surface Plasmon Resonance Sensor with BlueP/WS2-Covered Al2O3-Nickel Nanofilms**

**Shivangani 1, Maged F. Alotaibi 2, Yas Al-Hadeethi 2, Pooja Lohia 1,\*, Sachin Singh 3, D. K. Dwivedi 3,\*, Ahmad Umar 4,5,\*, Hamdah M. Alzayed 2, Hassan Algadi 5,6 and Sotirios Baskoutas <sup>7</sup>**


**Abstract:** In the traditional surface plasmon resonance sensor, the sensitivity is calculated by the usage of angular interrogation. The proposed surface plasmon resonance (SPR) sensor uses a diamagnetic material (Al2O3), nickel (Ni), and two-dimensional (2D) BlueP/WS2 (blue phosphorous-tungsten disulfide). The Al2O3 sheet is sandwiched between silver (Ag) and nickel (Ni) films in the Kretschmann configuration. A mathematical simulation is performed to improve the sensitivity of an SPR sensor in the visible region at a frequency of 633 nm. The simulation results show that an upgraded sensitivity of 332◦/RIU is achieved for the metallic arrangement consisting of 17 nm of Al2O3 and 4 nm of Ni in thickness for analyte refractive indices ranging from 1.330 to 1.335. The thickness variation of the layers plays a curial role in enhancing the performance of the SPR sensor. The thickness variation of the proposed configuration containing 20 nm of Al2O3 and 1 nm of Ni with a monolayer of 2D material BlueP/WS2 enhances the sensitivity to as high as 374◦/RIU. Furthermore, it is found that the sensitivity can be altered and managed by means of altering the film portions of Ni and Al2O3

**Keywords:** surface plasmon resonance sensor; blue-phosphorus tungsten di-sulfide; Al2O; nickel; sensitivity

#### **1. Introduction**

A method named surface plasmon resonance has arisen as an incredibly sensitive procedure for recognizing a very significant alteration in the refractive indexes of a detecting medium while communicating with the metal layer [1–3]. Enzyme detection, drug detection, medical diagnostics, and food safety are some of the biosensing applications of SPR-based biosensors [4–8]. Without any need for biomolecule labeling, a minute change in the refractive index (RI) can be detected in the detecting medium [9]. SPR is a highly sensitive technology that can detect very small fluctuations in the refractive index (RI) for biomolecule absorption of the order of 10−<sup>7</sup> on the sensing interface. At the metals' dielectric contact, the collective oscillation of free electrons generates a transverse magnetically polarised electromagnetic wave known as a surface plasma wave (SPW). SPWs are made from metals with negative permittivity, such as gold, silver, copper, and aluminium, as

**Citation:** Shivangani; Alotaibi, M.F.; Al-Hadeethi, Y.; Lohia, P.; Singh, S.; Dwivedi, D.K.; Umar, A.; Alzayed, H.M.; Algadi, H.; Baskoutas, S. Numerical Study to Enhance the Sensitivity of a Surface Plasmon Resonance Sensor with BlueP/WS2-Covered Al2O3-Nickel Nanofilms. *Nanomaterials* **2022**, *12*, 2205. https://doi.org/10.3390/ nano12132205

Academic Editors: Protima Rauwel and Erwan Rauwel

Received: 19 May 2022 Accepted: 8 June 2022 Published: 27 June 2022

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2022 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https:// creativecommons.org/licenses/by/ 4.0/).

well as dielectric materials, which can be liquid, gas, or solid. Researchers suggest that the Kretschmann arrangement brings about the productive coupling of light first from the crystal to the metal region, which is rooted in attenuated total reflection (ATR) [10]. The metallic regions of surface plasmons must be energized by the p-polarized light, while the s-polarised part is utilized as a source for the reference signal [11].

The SPR is acquired by the horizontal part of the evanescent wave (kev) being stagematched to the surface plasmon wave vector (ksp). For example,:

$$\mathbf{k\_{ev}} = \sqrt[\mathbf{k}]{\in p} \mathbf{k\_{p}} \sin \theta\_{\text{res}} = \mathbf{k\_{sp}}$$

where the incident wave vector is addressed by k0 = ω/c, crystal permittivity is addressed by ∈p, and the resonance point–angle is addressed by θres. Because of the total change of the p-polarized wave to the SP waves, the SPR condition causes a reduction in the depth of mirrored light (R).

In SPR sensors, normal metal layers of gold, silver, copper, and platinum are utilized. These are also plasmon-active metals that are used to generate surface plasmon waves in the sensing medium. The most promising metal is gold (Au), which has excellent optical properties, good chemical stability, and high oxidation and corrosion-resistant properties, although gold is the most expensive metal and lowers the biological molecules' absorption rate as compared to Ag metal [12–14]. The precision of an SPR sensor built from Ag film is better than that of an Au-film-based sensor. Because silver film is less costly than gold film, and as it was observed that silver has better sensitivity than Au metal, silver's SPR curve dip is narrower than gold's, meaning the sensitivity is improved. However, the chemical solidity problem of Ag should be alleviated, and the protecting layers must be explored for favorable optical properties [15–17]. As a result, this work introduces a novel aluminium oxide dielectric material to improve the SPR performance (Al2O3). It must be used to improve the SPR biosensor's performance parameters, such as the figure of merit (FOM), sensitivity, and so on. In fact, Al2O3 is widely used in mechanical areas due to its well-known properties, such as its superior corrosion resistance, high ductility, and high hardness, as well as in optical devices due to its high transparency and low refractive index. Certain 2D materials, such as graphene, WS2, and the 2D heterostructure BlueP-WS2, are the implicit aspirants used as a self-productive layer to give stability and boost the SPR sensor as it is slowly oxidized by air at room temperature and are considered corrosionresistant. Transition metal dichalcogenide (TMDC) substances have a curial position in SPR sensors. PtSe2, Ti3C2Tx-Mxene, blue phosphorus, black phosphorus, and transition metal dichalcogenides (TMDC) are examples of substances that have stood out in the past twenty years due to their notable optical and electrical nature and have been used to make optoelectronic devices. Forty unique synthetic substances are presently incorporated in the TMDC family. MX2 is another identifier for them, where M denotes metals such as tungsten, molybdenum, and niobium; and X denotes the chalcogen substances such as sulfur, selenium, and tellurium. MX2's monolayer has three nuclear layers, with an alternate metal layer embedded between two chalcogen substance layers [18]. These nanomaterials collaborate with the metal layer and enhance the co-operation of the particles. If the oxidation of Al2O3 (2D material) can be minimized by means of coating it with another layer, we can use it as a high-sensitivity sensor. Because of its refractive index, Al2O3 ensures that the biosensor's performance is no longer affected. The homogeneity of the Al2O3 layer is a characteristic that leads to sensitivity enhancements [19]. Nickel (Ni), a ferromagnetic metal, is also gaining interest due to its notable magneto-optical and magnetic properties and being a good light absorber. Using inert magnetic metals minimizes the cost of the SPR sensor while simultaneously considerably improving its performance [20–22].

Moreover, because the TMDCs and blue phosphorene have a similar hexagonal crystallike structure, BlueP/TMDCs can be formed [23]. Therefore BlueP/WS2's heterostructure shows greater sensitivity. The sensing medium, also known as a sensing analyte, is the outermost layer in the proposed SPR biosensor design. There are 2 main SPR geometry configurations: the Kretschmann configuration and Otto's design. The Kretschmann

configuration's benefits over Otto's arrangement have made it more widely applicable [24]. The SPR-based sensor's biosensing application covers the foundations of detecting the concentrations of biological things on a very small scale, such as bacteria, viruses, DNA, and proteins. Apart from biosensing and biomolecular analysis, the SPR sensor may also be used to detect nanostructured film depositions, as well as to quantify displacement and angular position [25–27].

In the present paper, calcium fluoride glass crystal (CaF2) is taken for the proposed SPR biosensor since it gives maximal sensitivity and an enormous change in reflectivity when contrasted with different crystals (BK7, SF10, SF11, and so forth). A framework for an SPR sensor based on heterostructure containing Al2O3, BlueP/WS2 with metal (Ag), and Ni (Nickle) is proposed in the present work to achieve improved sensitivity by altering the thickness of the Ni layers. The results demonstrate that adding the Al2O3 layer and BlueP/WS2 to this structure boosts the sensitivity substantially. In this work, COMSOL Multiphysics 5.3a and MATLAB 2016a softwares are used to draw the plot of the reflectance curve, sensitivity curve, and electric field curve, which are calculated using these software programs.

**Figure 1.** Sketch diagram of the proposed SPR biosensors.

#### **2. Mathematical Modeling for the Proposed SPR Biosensor**

*2.1. Device Structure*

In the present paper, the SPR sensor is composed of multiple layers primarily based on the Kretschmann configuration, in which the prism of CaF2, Ag as a metal, Al2O3 (diamagnetic material), BlueP/WS2, and SM on the top layer is used as shown in Figure 1. Table 1 refers to the width and different refractive indexes of the layers.

**Table 1.** Details of each layer of the proposed biosensor at 633 nm wavelength.


The structural diagram using Ag-Al2O3-Ni-BlueP/WS2 is depicted in Figure 1 for this model. The 633 nm wavelength is utilized in SPR sensors for optimal results. The CaF2 prism is employed in metal Ag with a refractive index of 1.4329 and a thickness of 50 nm, in Al2O3 with a thickness of 20 nm, in nickel with a thickness of 1 nm, and in BlueP/WS2 with a thickness of 0.75 nm. Finally, the sensing medium has a refractive index of 1.330–1.335. The variation in the detecting medium is caused by adsorption, which occurs when biomolecules in the sensing medium contact with the BlueP/WS2 layer, causing the sensing medium's RI to alter. The first film is the CaF2 prism, whose refractive index value may be computed using the Sellmeier relation.

$$\ln^2 = 1.33973 + \frac{0.69913 \times \lambda^2}{\lambda^2 - (0.09374)^2} + \frac{0.11994 \times \lambda^2}{\lambda^2 - (21.18)^2} + \frac{4.35181 \times \lambda^2}{\lambda^2 - (38.46)^2} \tag{1}$$

Here, 'λ' is the wavelength in nanometers. The RI of Ag is calculated using the Drude–Lorentz model:

$$\mathbf{n}\_{\text{metal}}(\lambda) = \left(1 - \frac{\lambda^2 \times \lambda\_{\text{C}}}{\lambda\_{\text{P}}^2 (\lambda\_{\text{C}} - \lambda \times \text{i})}\right)^{\frac{1}{2}} \tag{2}$$

where λ<sup>c</sup> and λ<sup>P</sup> are the collision and plasma wavelengths, respectively. These are dispersion coefficients, while the used wavelength is 633 nm.

For Ag, λ<sup>P</sup> = 145.41 nm and λ<sup>C</sup> = 176.14 nm [33]. The transfer matrix method for the n-layer modeling and Fresnel equations is used throughout the numerical analysis. Using the reflectance curve, all the SPR sensor's performance characteristics in MATLAB software are evaluated. The graph is drawn using Origin software for all parameters and the FWHM values. Here, additionally compared distinct SPR sensors with the proposed model are mentioned in Table 2.


**Table 2.** Comparative study of the proposed SPR model.

#### *2.2. Mathematical Expression for Reflectivity*

The transfer matrix approach is utilized in the present work to obtain the reflection coefficient of the projected multilayer design (crystal, metal, Al2O3, Ni, and BlueP/WS2 film). For the calculation of the reflectivity of the reflected light, the matrix approach for the N-film design is used. This method is quick and easy to use, and it does not involve any approximation. Along the z-axis, the layer thicknesses, dk, is considered. The kth layer's dielectric constant and RI are denoted by k and n, respectively. The tangential fields at Z=Z1 = 0 are expressed in terms of the tangential field at Z = ZN−<sup>1</sup> using the boundary condition [34]:

$$
\begin{bmatrix} \mathbf{U}\_1 \\ \mathbf{V}\_1 \end{bmatrix} = \mathbf{M} \begin{bmatrix} \mathbf{U}\_{\mathbf{N}-1} \\ \mathbf{V}\_{\mathbf{N}-1} \end{bmatrix} \tag{3}
$$

The places U1 and V1 address the tangential element of electric and magnetic fields, respectively, and the first and last layer of boundary are denoted by UN−1,VN−1, respectively. Mij is the element for which the characteristics matrix is as follows [35]:

$$\mathbf{M}\_{\ddot{\mathbf{j}}} = \left(\prod\_{\mathbf{k}=2}^{\mathbf{N}-1} \mathbf{M}\_{\mathbf{k}}\right)\_{\dddot{\mathbf{i}}\mathbf{j}} = \begin{bmatrix} \mathbf{M}\_{11} & \mathbf{M}\_{12} \\ \mathbf{M}\_{21} & \mathbf{M}\_{22} \end{bmatrix} \tag{4}$$

$$\mathbf{M}\_{\mathbf{k}} = \begin{bmatrix} \cos \beta \,\kappa & (-\dot{\mathbf{s}} \sin \beta \,\kappa) / \mathbf{q}\_{\mathbf{k}} \\ -\dot{\mathbf{i}} q\_{\mathbf{k}} \sin \beta \,\kappa & \cos \beta \,\kappa \end{bmatrix} \tag{5}$$

$$\mathbf{q}\_{\mathbf{k}} = \left(\frac{\mathbf{u}\_{\mathbf{k}}}{\varepsilon\_{\mathbf{k}}}\right)^{1/2} \cos \theta \mathbf{x} = \frac{\left(\varepsilon\_{\mathbf{k}} - \sin \theta\_1 \mathbf{n}\_1^2\right)^{1/2}}{\varepsilon \mathbf{k}} \tag{6}$$

and

$$\mathbf{q}\_{\mathbf{k}} = \left(\frac{\mathbf{u}\_{\mathbf{k}}}{\varepsilon\_{\mathbf{k}}}\right)^{1/2} \cos \theta \mathbf{x} = \frac{\left(\varepsilon\_{\mathbf{k}} - \sin \theta\_1 \mathbf{n}\_1^2\right)^{1/2}}{\varepsilon \mathbf{k}} \tag{7}$$

Following the mathematical steps, one can attain the reflection coefficient for ppolarized light, which is given below:

$$\mathbf{r}\_{\mathbf{P}} = \frac{(\mathbf{M}\_{11} + \mathbf{M}\_{12}\mathbf{q}\_{\mathbf{n}})\mathbf{q}\_{1} - (\mathbf{M}\_{21} + \mathbf{M}\_{22}\mathbf{q}\_{\mathbf{n}})}{(\mathbf{M}\_{11} + \mathbf{M}\_{12}\mathbf{q}\_{\mathbf{n}})\mathbf{q}\_{1} + (\mathbf{M}\_{21} + \mathbf{M}\_{22}\mathbf{q}\_{\mathbf{n}})} \tag{8}$$

The multilayer configuration of the reflectivity Rp is given as:

$$\mathcal{R}\_{\mathbb{P}} = |r\_P|^2 \tag{9}$$

The conventional and three revised characteristics plots of the specific SPR sensor are shown in Figure 2. For the conventional SPR sensor shown in Figure 2a, the variation of the refractive index is Δn = 0.005, while the Δθ and sensitivity and full width at half maximum (FWHM) are 0.72(deg), 144◦/RIU, and 1.04677, respectively. Further, Figure 2b–d presents a comparison with the SPR reflectance curve from Figure 2a. The reflectance curves of SPR designs 2 (Ag/Al2O3) and 3 (Ag/Al2O3/Ni) are displayed in Figure 2b,c, respectively. The angular shifts (Δθ) for Figure 2b,c are 1.48 and 1.6, respectively, leading to sensitivities and FWHMs of 296◦ RIU−<sup>1</sup> and 320◦ RIU−1, and 2.05852 deg and 2.25936 deg, respectively. Figure 2d suggests a superior resonance angle, sensitivity, and FWHM in contrast to the other 3 designs. Figure 2d has the most extreme sensitivity of 397◦/RIU among all other structures. Table 3 shows not only the maximum sensitivity but also the maximum resonance angle and a larger figure of merit compared to the other 3 designs.

**Figure 2.** *Cont*.

**Figure 2.** *Cont*.

**Figure 2.** SPR reflectance curves: (**a**) design 1; (**b**) design 2; (**c**) design 3; (**d**) design 4.

**Table 3.** Comparative study of different proposed SPR sensors at a wavelength of 633 nm and Δn = 0.005.


Consequently, design 4 is considered the most reasonable decision among all designs.

#### **3. Results and Discussion**

The sensitivity properties of the biosensor with the changed Kretschmann design, which incorporates Al2O3 and Ni, are discussed here. To show how the sensitivity has reached the next level, from the reflectance curve the sensitivity is assessed according to the change in the resonance point. The excitement of the SPR causes a sharp drop in reflectance at a given point, which is clearly apparent. This event shows that the light is consumed by initiating the SPR in the biosensor arrangement, while the atom connection causes minor shifting in the refractive index of the sensor, which has a slight resonance dip around 1.87◦. Thus, the design's sensitivity is achieved (Sn = 374◦/RIU) by utilizing the connected computation articulation Sn =S= <sup>Δ</sup>θres <sup>Δ</sup><sup>n</sup> as displayed in Figure 2.

The effects of the thickness variation of Al2O3 and Ni on the performance of the projected SPR sensor are shown in Table 4. The execution constraints of the proposed SPR sensor for the variety of thicknesses of Al2O3 (14–20 nm) and Ni (1, 3, 5 nm) are displayed in Table 4. The greatest sensitivity of 374◦/RIU is obtained with a thickness for Al2O3 of 20 nm and for Ni of 1 nm at a working wavelength of 633 nm.


**Table 4.** The optimized thickness values of Al2O3 and Ni with respect to the other parameters such as ΔθSPR, S, DA, FOM, and FWHM.

#### *3.1. Use of CaF2 Crystal*

The refractive index directly affects the performance of the SPR sensor. Since the shift with the incident point in the reflectance bend is high and a sharp plunge is acquired, a CaF2 crystal is utilized in the present situation. The sensitivity of this crystal material is incredible, with a lower RI than the CaF2 crystal. Therefore, the CaF2 glass crystal is at last utilized in the SPR sensor in the present theoretical investigation [36].

#### *3.2. Performance Constraints of the SPR Sensor*

The sensitivity, full width at half maximum, quality factor, detection accuracy, and limit of detection of the SPR sensor rely on certain variables. All of these constraints are dependent upon one another, and the reflectivity curve versus the incident angle determine the mathematical study of the SPR design.

#### *3.3. Sensitivity (S)*

The variation in resonance angle (Δθres) with respect to the variation in the refractive index (Δn) decides the sensitivity and is defined as [37]:

$$\mathbf{S} = \frac{\Delta\theta\_{\rm res}}{\Delta\mathbf{n}} \left( \mathbf{U} \text{mit} : \text{\textdegree RIU}^{-1} \right) \tag{10}$$

The variation of the reflectance with the incident angle for the proposed SPR sensor is shown in Figure 3.

Figure 3 shows the shifts in resonance angle with the incident angle at different RI values (nSM = 1.330–1.335). From Figure 3, the greatest alteration in resonance angle (1.87) is acquired for the present SPR design. The most extreme change in the resonance point shows the adjustment of the coupling state of the surface plasmon wave (SPW). The device's sensitivity should always be high. This means that the higher sensitivity sensor detects the minute variations in analyte (biomolecules) concentration, which shows that the sensor has superior sensing capabilities because it can easily detect minute RI variations in the structure.

**Figure 3.** Variation of resonance angle vs. incident angle.

#### *3.4. Quality Factor (QF)*

The division of sensitivity with the full width at half maximum is defined as the quality factor. It is also called the figure of merit (FOM). On the other hand, the quality factor is the multiplication of the sensitivity by the detection accuracy. The figure of merit (FOM) is a quantity used to characterize the performance as the device increases [38].

$$\text{QF} = \frac{\Delta\theta\_{\text{res}}}{\Delta\mathbf{n} \times \text{FWHM}} \text{ (unit}: \text{ RIU}^{-1} \text{)}\tag{11}$$

The variation of the FOM with the different layer thicknesses of Al2O3 is shown in Figure 4. It can be observed that the FOM is at its maximum for the Ni thickness of 1 nm.

**Figure 4.** Variation of FOM vs. different layer thicknesses of Al2O3.

#### *3.5. Full Width at Half Maximum (FWHM)*

The incidence angle changes at the halfway point of the reflectance curve can be used to calculate the full width at half maximum (FWHM), and its values should be low to boost the FOM. The FWHM assumes a significant role in the sensor's execution because the majority of the parameters rely upon it. The width of the reflectance curve as the resonance angle shifts is measured by the FWHM. A low FWHM reduces the uncertainty in determining the resonance dip, and as a result improves the sensor's resolution. It is defined by [39]:

$$\text{FWHM} = \frac{1}{2} \left( \theta\_{\text{max}} + \theta\_{\text{min}} \right) , \text{ (unit} : \text{ degree}) \tag{12}$$

The FWHM with respect to the Ni thickness is shown in Figure 5.

**Figure 5.** Variation of thickness of Ni vs. FWHM.

Figure 5 shows that the variations of Ni vs. FWHM give the best result for the FWHM at the 20 nm thickness of Al2O3 when the thickness of Ni varies from 1 nm to 5 nm. The FWHM is also used to calculate the quality factor and detection accuracy.

#### *3.6. Detection Accuracy (DA)*

This is conversely connected with the full width at half maximum. It is also termed the signal-to-noise ratio, and the ratio of the signal to the noise should be high as possible to better the device quality. Basically, it shows how the noise level is impacting the device structure. It is defined as [30]:

$$\text{DA} = \left[ \frac{1}{\text{FWHM}} \right] \Big/ \left( \text{degree}^{-1} \right) \tag{13}$$

The detection accuracy vs. layer thickness plot for Al2O3 is shown in Figure 6. It can be observed that the detection accuracy decreases with the layer thickness.

**Figure 6.** Plot of DA vs. different layer thickness for Al2O3.

#### *3.7. Limit of Detection (LOD)*

This is the difference in biomolecule fixation or analyte concentration in the detecting area, and it is also defined as the proportion of progress in the RI (Δn) with the change in resonance angle (Δθres) [30]:

$$\text{LOD} = \frac{\Delta \text{n}}{\Delta \theta\_{\text{res}}} \times 0.001^0 \tag{14}$$

where 0.001<sup>0</sup> is a very small shift in the detecting medium.

#### *3.8. The Impact of the Refractive Index on the Reflectance Curves of Different SPR Sensor Structures*

The curves of SPR reflectivity for the sensor in the detecting medium with the RI range from 1.330 to 1.335. The proposed SPR curve for the changing RI of an analyte (ns = 0.005) with adjustment of the resonance angle (1.87◦), sensitivity (374◦/RIU), quality factor (56.211 RIU<sup>−</sup>1), and DA (0.281 deg−1) for the recommended design might be found in the SPR reflectance bends. The sensitivity and resonance angle shift are significantly bigger than with conventional sensors, as displayed in Table 3. The variation of the resonance angle with the RI of the sensing medium for the conventional and proposed sensor designs are plotted in Figure 7.

**Figure 7.** Variation of the resonance angle with Ri for different sensor designs.

#### *3.9. Optimization of the Thicknesses of Al2O3 and Metal (Ni)*

The refractive index of Al2O3 is small, so it should be selected because it minimizes the losses in performance of the biosensor, which is another characteristic contributing to the sensor's improvement [32]. Since the Al2O3 layer has small damping properties with a higher penetration rate of the surface plasmon in the sensing medium, this property helps to stop the corrosion phenomena and enhances the performance of the SPR sensor [40]. On the other hand, the cost of the SPR sensor is reduced by using the ferromagnetic material nickel (Ni) because it has incredible magneto-optical characteristics, magnetic qualities, and small optical losses. At the optimum thickness of Ni (1 nm), the molecular absorption of light increases with minimum reflectance. The Ni layer is also used as a protective layer, which helps to raise the sensitivity of the SPR sensor [41]. Figure 8 shows a graph demonstrating the fluctuations in Al2O3 and Ni thickness and the varying sensitivities at different thicknesses. Figure 8a shows that the maximum sensitivity is achieved with a layer of Al2O3 of 20 nm. For metal, the best sensitivity is 374◦/RIU (Ag). To determine the best sensitivity for the suggested SPR sensor, the thickness of Ni is optimized, as Figure 8b illustrates, whereby the sensitivity changes with the changes in Ni thickness. The optimal thickness ranges from 1 to 5 nm, with a maximum sensitivity of 374◦/RIU and minimal reflection (Rmin). Consequently, the proposed SPR sensor shows the best results when the thickness of the Ni layer is 1 nm.

**Figure 8.** (**a**) Variation of sensitivity vs. Al2O3 (nm). (**b**) Variation of sensitivity vs. Ni (nm).

#### *3.10. Parameter Analysis of the Projected SPR Biosensor*

The design (CaF2/Ag/Al2O3/Ni/BlueP/WS2/sensing medium) arrangement has the most elevated sensitivity (374◦RIU−1) for sensor applications. It is additionally advantageous for energizing surface plasmons by changing over from the crystal-directed mode to the surface plasmon polariton (SPP) mode proficiently. The performance parameters, sensitivity, and QF show increments with high RI values. However, the DA (detection accuracy) diminishes, which causes the increase in FWHM increments with a huge change in the reflectivity bend. The FWHM and DA have an opposite relationship.

#### *3.11. Clarification of Transvers Magnetic (TM) Field and Penetration Depth*

The transverse magnetic (TM) field plot is discussed in this part. Utilizing the COM-SOL Multiphysics programming software, the recommended SPR biosensor's transverse magnetic (TM) field variation is shown in Figure 9. The TM field likewise helps with the estimation of a vital measurement, the sensor's penetration depth. The distance of the electric field force diminishes to 1/e, which is called the penetration depth, and its value is 108.25 nm. Subsequently, when compared with past SPR sensor calculations, the proposed sensor has the best penetration depth, and along these lines is more delicate. To show the electromagnetic field, the effective mode index (EMI = 1.0926 − 0.31363 × i) is utilized [42].

**Figure 9.** Penetration depth variation and transverse magnetic field.

#### **4. Conclusions**

In the work, the sensitivity was improved by utilizing a layer of Al2O3 over the metals silver (Ag) and nickel (Ni) on the top. A layer of 2D material, BlueP/WS2, was utilized to upgrade the sensitivity and safeguard the device from corrosion. The greatest sensitivity was found for Ag metal (374◦/RIU) and the primary arrangement for the most extreme sensitivity was CaF2/Ag/Al2O3/Ni/BlueP/WS2/SM. The performance qualities of the heterostructure-based SPR sensors, such as FWHM, detection accuracy (DA), and LOD, showed great correlations with traditional sensors for the appropriate scope of the RI from 1.330 to 1.335. From the above study, the projected SPR sensor configuration has incredible sensitivity and could be utilized in the biosensing field.

**Author Contributions:** Data curation: S., Y.A.-H., P.L., S.S., D.K.D. and A.U.; Formal analysis: M.F.A., P.L., H.M.A., H.A. and S.B.; Methodology: S., M.F.A., S.S. and H.M.A.; Project administration: Y.A.-H., D.K.D. and A.U.; Supervision: S.B.; Writing—original draft, P.L. and H.A. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research work was funded by Institutional Fund Projects under grant no. (IFPIP: 1663-130-1442).

**Data Availability Statement:** Not applicable.

**Acknowledgments:** Authors gratefully acknowledge technical and financial support from the Ministry of Education and King Abdulaziz University, DSR, Jeddah, Kingdom of Saudi Arabia.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**


### *Article* **Significance of Hydroxyl Groups on the Optical Properties of ZnO Nanoparticles Combined with CNT and PEDOT:PSS**

**Keshav Nagpal 1, Erwan Rauwel 1, Elias Estephan 2, Maria Rosario Soares <sup>3</sup> and Protima Rauwel 1,\***

<sup>1</sup> Institute of Forestry and Engineering, Estonian University of Life Sciences, 51014 Tartu, Estonia


**Abstract:** We report on the synthesis of ZnO nanoparticles and their hybrids consisting of carbon nanotubes (CNT) and polystyrene sulfonate (PEDOT:PSS). A non-aqueous sol–gel route along with hydrated and anhydrous acetate precursors were selected for their syntheses. Transmission electron microscopy (TEM) studies revealed their spherical shape with an average size of 5 nm. TEM also confirmed the successful synthesis of ZnO-CNT and ZnO-PEDOT:PSS hybrid nanocomposites. In fact, the choice of precursors has a direct influence on the chemical and optical properties of the ZnO-based nanomaterials. The ZnO nanoparticles prepared with anhydrous acetate precursor contained a high amount of oxygen vacancies, which tend to degrade the polymer macromolecule, as confirmed from X-ray photoelectron spectroscopy and Raman spectroscopy. Furthermore, a relative increase in hydroxyl functional groups in the ZnO-CNT samples was observed. These functional groups were instrumental in the successful decoration of CNT and in producing the defect-related photoluminescence emission in ZnO-CNT.

**Keywords:** ZnO; ZnO-CNT; ZnO-PEDOT:PSS; nanoparticles; hybrids; hydroxyl groups; non-aqueous sol–gel; surface defects; photoluminescence

#### **1. Introduction**

Hybrid nanocomposites combining organic and inorganic counterparts have a multitude of applications, e.g., light-emitting diodes (LED), solar cells, and photodetectors [1–3]. Organic materials consist of polymers possessing remarkable properties, such as easy processing, flexibility, and good conductivity [4]. However, their high cost and lack of stability are obstacles for practical devices. On the other hand, inorganic materials present higher structural, chemical, and functional stability, as well as a high charge mobility, making them suitable for optoelectronic applications [5]. Therefore, the combination of organic with inorganic materials provides robust multifunctional nanocomposites with applications in flexible electronic and photonic devices [6–8].

Conducting polymers such as poly (3,4-ethylenedioxythiophene) poly(styrenesulfonate) (PEDOT:PSS) are already being incorporated into organic thin film transistors, organic LED, organic solar cells, capacitors, batteries, and thermoelectric devices, as well as technologies such as touch screens and electronic papers [9,10]. In addition, PEDOT:PSS is mechanically stable and highly flexible. Various combinations of PEDOT:PSS with inorganic materials, such as SnO2, TiO2, CdS, CdSe, ZnO and metal nanostructures, have been investigated to that end [11,12]. Among these inorganic nanomaterials, ZnO is promising due its wide band gap of 3.37 eV, large exciton binding energy of 60 meV, high chemical stability, and remarkable electrical and optical properties [13]. Moreover, the high surface-to-volume ratio of ZnO nanoparticles implies a spontaneous presence of surface defects, including oxygen vacancies (VO), oxygen interstitials (Oi), and zinc interstitials (Zni). Therefore, in addition to the UV emission, known as the near-band emission (NBE), ZnO nanoparticles

**Citation:** Nagpal, K.; Rauwel, E.; Estephan, E.; Soares, M.R.; Rauwel, P. Significance of Hydroxyl Groups on the Optical Properties of ZnO Nanoparticles Combined with CNT and PEDOT:PSS. *Nanomaterials* **2022**, *12*, 3546. https://doi.org/10.3390/ nano12193546

Academic Editor: Michael Tiemann

Received: 20 September 2022 Accepted: 7 October 2022 Published: 10 October 2022

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emit within the entire visible spectrum, also known as defect-level emission (DLE) [14]. The latter depends on both the surface and the volume defects introduced during the synthesis of the nanoparticles [15,16]. For example, ZnO nanoparticles prepared by aqueous sol– gel routes tend to emit higher NBE and a negligible DLE [17]. On the other hand, for ZnO nanoparticles prepared by non-aqueous sol–gel routes, the emission depends on the presence of hydrates in the precursor [18,19]. In fact, hydrates in the precursor contribute to the enhancement of NBE due to improved oxidation of ZnO during synthesis [18]. On the other hand, adsorption of hydroxyl groups on the surface of ZnO nanoparticles has been shown to increase the visible PL emission [20]. The chemisorption of oxygen radicals from air on the surface of ZnO nanoparticles also augments green emission from them [21]. In general, small nanoparticles possess a high surface-to-volume ratio, and therefore harbor higher amounts of surface defects. For larger ZnO nanoparticles, defects can be both surface and volume related [22].

Recently, ZnO-CNT nanohybrids have attracted considerable interest due to their high stability and superior photonic, electrochemical and electromagnetic properties, which originate from interfacial effects. In a previous study, we successfully passivated ZnO surface states by combining them with CNT via sonication [17]. In this work, we carry out ZnO nanoparticle synthesis via a non-aqueous sol–gel route with hydrated (zinc acetate dihydrate (Zn(CH3CO2)2.2H2O) and anhydrous (zinc acetate anhydrous (Zn(CH3CO2)2) precursors. In a study by Šari´c et al., it was shown that with similar precursors, ZnO precipitation could be promoted through an esterification reaction that generates water upon the addition of acetic acid [23]. In our reaction, only absolute ethanol is used as a solvent. It plays a crucial role in controlling the size and shape of ZnO nanoparticles, and consequently, in the formation of various surface defects. Due to the addition of sodium hydroxide in this work, the basic character of the solution prevents any esterification reaction and in turn, no water molecules are formed. This route therefore enables the formation of very small ZnO nanoparticles functionalized with hydroxyl groups, promoting the decoration of CNT with ZnO. We then decorated CNT with ZnO nanoparticles in order to create a hybrid nanocomposite. Subsequently, we fabricated a second type of hybrid nanocomposite consisting of ZnO-PEDOT:PSS and compared the evolution of the surface defects to ZnO-CNT. These optical properties are discussed in terms of synthesis conditions, crystal structure, chemical properties, and the morphology of ZnO nanoparticles and their hybrids. The objective is to use these hybrid materials in LED. Therefore, finding a way to control these surface defects or trap states for LED applications is a priority, as they are detrimental to device properties.

#### **2. Materials and Methods**

#### *2.1. Synthesis*

#### 2.1.1. ZnO

Two different zinc precursors, Zn(CH3CO2)2.2H2O (99.5%, Fisher Scientific, Loughborough, UK) and Zn(CH3CO2)2 (99.9%, Alfa Aesar, Kandel, Germany), were used for the synthesis of ZnO nanoparticles via non-aqueous sol–gel routes. Sodium hydroxide (NaOH) (99.9%, Aldrich) was used as a reducing agent. All the chemicals used were of analytic reagent grade. To prepare 0.05 M solutions of zinc precursors, 219.5 mg of Zn(CH3CO2)2.2H2O or 183.48 mg of Zn(CH3CO2)2 were dissolved in 20 mL absolute ethanol in a beaker placed in a water bath. The solutions were maintained at 65 ◦C under continuous magnetic stirring until the precursors were completely dissolved in absolute ethanol. Furthermore, a solution of 0.10 M NaOH in 20 mL absolute ethanol was prepared. The NaOH solution was added dropwise to the zinc precursor solutions. Thereafter, the mixtures were maintained at 65 ◦C for 2 h after which they were cooled to ambient temperature. White ZnO precipitates settled at the bottom of the reaction vessel. The resulting solutions containing ZnO nanoparticles were then centrifuged at 4500 rpm for 6 min, followed by drying for 24 h in air at 60 ◦C. This resulted in an agglomeration of ZnO nanoparticles in the form of a pellet, which is typical after drying nanoparticles synthesized

via sol–gel routes. These pellets were thereafter crushed using a pestle and mortar to obtain a very fine powder of ZnO nanoparticles.

#### 2.1.2. ZnO-CNT Hybrids

For the preparation of ZnO-CNT hybrids, firstly, a solution of CNT was prepared by mixing 4 mg of CNT in 50 mL absolute ethanol and sonicating until a homogenous mixture was obtained. As before, 0.05 M zinc precursor solutions were prepared in 20 mL absolute ethanol. To prepare 0.10 M NaOH solutions, ~80 mg of NaOH was added to 19 mL absolute ethanol, in which 1 mL CNT mixture (~0.08 mg) was added. The final mixtures were sonicated and added dropwise to zinc precursor solutions. Thereafter, the reaction was completed as described earlier, and ZnO-CNT hybrid pellets were obtained. These pellets were gently crushed to obtain fine black powders of ZnO-CNT nanohybrids.

#### 2.1.3. ZnO-PEDOT:PSS Hybrids

For the preparation of ZnO-PEDOT:PSS hybrids, 20 mg of the as-synthesized ZnO nanoparticles were taken, to which 400 mg of PEDOT:PSS (as purchased) was added. The mixtures were sonicated for 1 h and dried at 70 ◦C for 24 h. The dried mixtures were further gently crushed to obtain blue powders with agglomerated particles of ZnO-PEDOT:PSS nanohybrids.

#### *2.2. Characterization*

X-ray diffraction patterns were collected in Bragg–Brentano geometry using a Bruker D8 Discover diffractometer (Bruker AXS, Germany) with CuKα1 radiation (λ = 0.15406 nm) selected by a Ge (111) monochromator and LynxEye detector. Transmission electron microscopy (TEM) was carried out on a Tecnai G2 F20 (Netherlands) is a 200 kV field emission gun (FEG) for high-resolution and analytical TEM/STEM. It provided a point-topoint resolution of 2.4 Å. XPS measurements were performed at room temperature with a SPECS PHOIBOS 150 hemispherical analyzer (SPECS GmbH, Berlin, Germany)) with a base pressure of 5 × <sup>10</sup>−<sup>10</sup> mbar using monochromatic Al K alpha radiation (1486.74 eV) as excitation source operated at 300 W. The energy resolution as measured by the FWHM of the Ag 3d5/2 peak for a sputtered silver foil was 0.62 eV. The spectra were calibrated with respect to the C1s at 284.8 eV. The optical absorbance of ZnO, ZnO-CNT and ZnO-PEDOT:PSS nanohybrids was determined using a NANOCOLOR UV-VIS II spectrometer (MACHEREY-NAGEL, Germany) in the 200–900 nm region. The band gap of ZnO, ZnO-CNT and ZnO-PEDOT:PSS nanohybrids was subsequently calculated using Tauc plots. PL spectroscopy was carried out at room temperature with an excitation wavelength of 365 nm of an LSM-365A LED (Ocean Insight, USA) with a specified output power of 10 mW. The emission was collected by FLAME UV-Vis spectrometer (Ocean optics, USA) with a spectral resolution 1.34 nm. Optical images of ZnO were taken under a UV lamp ZLUV220 (China) with an excitation source of 365 nm. Raman spectra were collected using a WITec Confocal Raman Microscope System alpha 300R (WITec Inc., Ulm, Germany). Excitation in confocal Raman microscopy is generated by a frequency-doubled Nd:YAG laser (New-port, Irvine, CA, USA) at a wavelength of 532 nm, with 50 mW maximum laser output power in a single longitudinal mode. The system was equipped with a Nikon (Otawara, Japan) objective with a X20 magnification and a numerical aperture NA = 0.46. The acquisition time of a single spectrum was set to 0.5 s.

#### **3. Results**

Table 1 provides a list of ZnO, ZnO-CNT and ZnO-PEDOT:PSS samples synthesized in this work. Samples ZnO-D and ZnO-A correspond to ZnO nanoparticles synthesized from Zn(CH3CO2)2.2H2O and Zn(CH3CO2)2 precursors, respectively. Samples ZnO-D-CNT, ZnO-A-CNT, ZnO-D-PEDOT:PSS and ZnO-A-PEDOT:PSS correspond to the hybrids of samples ZnO-D and ZnO-A with CNT and PEDOT:PSS, respectively. In this study, the terms ZnO samples refer to samples ZnO-D and ZnO-A; ZnO-CNT hybrids refer to

#### samples ZnO-D-CNT and ZnO-A-CNT and ZnO-PEDOT:PSS hybrids refer to samples ZnO-D-PEDOT:PSS and ZnO-A-PEDOT:PSS.


ZnO-A-PEDOT:PSS - ~95

**Table 1.** List of ZnO-CNT hybrids and ZnO-PEDOT:PSS hybrids synthesized in this work.

#### *3.1. Structure and Morphology*

The XRD patterns of ZnO samples ZnO-D and ZnO-A are shown in Figure 1. The peaks (100), (002), (101), (102), and (110) correspond to the hexagonal Wurtzite structure (a = 3.25 Å and c = 5.20 Å) of ZnO (JCPDS, Card Number 36-1451). No secondary phases are visible in the XRD patterns, indicating that single-phase ZnO nanoparticles were formed. In addition, XRD patterns illustrate that both samples ZnO-D and ZnO-A exhibit very small particle sizes due to broader XRD peaks. The size of nanoparticles was estimated using the Scherrer equation [24].

$$\mathbf{D} = \frac{0.9\lambda}{\beta \cos \theta} \tag{1}$$

where D is particle size, λ (=0.15406 nm) is the wavelength of incident X-ray beam, β is FWHM in radians, and θ is Bragg's diffraction angle. Size calculation was carried out by considering the highest-intensity (101) peak. The calculated ZnO nanoparticle sizes of samples ZnO-D and ZnO-A were ~9 nm and ~5 nm, respectively. However, the actual size and shape of ZnO nanoparticles were confirmed from TEM studies as discussed below.

**Figure 1.** Normalized XRD patterns of samples ZnO-D and ZnO-A.

The morphological features of the as-synthesized ZnO samples, ZnO-CNT hybrids and ZnO-PEDOT:PSS hybrids were studied by TEM, as shown in Figure 2. TEM images in Figure 2a,b consist of overviews of the as-grown samples ZnO-D and ZnO-A, respectively. The micrographs reveal spherical nanoparticles of uniform size that tend to agglomerate. With the help of size distribution histograms of the as-synthesized ZnO samples, we estimate an average nanoparticle size of ~5.2 nm and ~4.8 nm for ZnO-D and ZnO-A, respectively. Figure 2c is a high-resolution TEM (HRTEM) image of sample ZnO-A, where

two ZnO nanoparticles are oriented along the [0001] zone axis of the basal plane of the Wurtzite structure. Figure 2d,e are low-magnification TEM images of the samples ZnO-D-CNT and ZnO-A-CNT, respectively. ZnO nanoparticles dominate the TEM images due to the low wt% (~1 wt%) of CNT in the samples. HRTEM images of ZnO-CNT are presented in Figure 2g,h. The walls of the CNT are clearly visible along with nanoparticles decorating them. We observe that the presence of CNT does not alter the crystallinity or the size distribution of the nanoparticles, and average sizes of ~5.7 nm and ~4.7 nm were retained. The micrographs therefore clearly indicate successful decoration of the nanoparticles on the walls of the CNT. In our study, the nanotubes were functionalized by sonication in pure ethanol; hence, the most likely functional groups present are carboxyl (COOH) that can be broken down into carbonyl (C-O) and hydroxyl (OH) [25]. These functional groups promote a covalent bonding between the CNT and ZnO nanoparticles, necessary for the decoration of CNT. Figure 2f,i are the TEM and scanning transmission electron microscopy (STEM) images of ZnO-D-PEDOT:PSS and ZnO-A-PEDOT:PSS samples, respectively. PE-DOT:PSS appears as flakes without any noticeable agglomeration of ZnO nanoparticles in the polymer layer. However, some areas of PEDOT:PSS are more densely packed with ZnO nanoparticles. The insets of Figure 2f,i are high-magnification TEM images of the samples emphasizing on their homogeneous distribution in the PEDOT:PSS matrix.

**Figure 2.** Overview TEM images of samples (**a**) ZnO-D, (**b**) ZnO-A, (**c**) HRTEM image of ZnO-A nanoparticles. Overview TEM images of (**d**) ZnO-D-CNT, (**e**) ZnO-A-CNT, (**f**) ZnO-D-PEDOT:PSS. HRTEM images of (**g**) ZnO-D-CNT, (**h**) ZnO-A-CNT and (**i**) STEM image of ZnO-A-PEDOT:PSS.

The high-resolution XPS spectra of the C 1s and O 1s regions of the as-synthesized and hybrid ZnO nanoparticles are shown in Figures 3 and 4, respectively. For the C 1s spectra of the as-synthesized samples in Figure 3a,d, the photoelectron peak at 284.8 eV corresponds to adventitious carbon [26]. Several carbon bonds are present in the samples, such as C-OH, O=C-O originating from the NaOH and acetate precursors used in the syntheses [27]. Both ZnO-A and ZnO-D contain oxygen and hydroxyl groups that are chemisorbed. The C 1s region of ZnO-CNT in Figure 3b,e manifests an additional peak corresponding to sp2

hybridization of C atoms in the CNT at a binding energy of 283 eV. In addition, the C-OH peak is relatively more intense for the ZnO-A-CNT sample compared to the as-synthesized ZnO sample in Figure 3d, which could indicate an increase of hydroxyl groups or oxygen vacancies [18]. In fact, sonication of CNT in ethanol engenders a breakdown of the sidewalls, which then produces C-dangling bonds [28]. After sonication and during the initial stages of synthesis, CNT were mixed in ethanol and heated to a temperature of 65 ◦C for 2 h during which a solution of NaOH was added dropwise. Considering the hydroxyl rich conditions, the attachment of OH groups to C-dangling bonds is likely. For the ZnO-D-based samples, in Figure 3a–c, the relative intensities of the various peaks in the C1s region are similar, unlike the ZnO-A-based samples. In addition, in Figure 3f, a decrease in the O=C-O and C-OH peak intensities relative to the C-C peak for sample ZnO-A-PEDOT:PSS is observed, indicating an oxygen-deficient or -reduced PEDOT:PSS polymer.

**Figure 3.** High-resolution XPS spectra of the C1s region of (**a**) ZnO-D, (**b**) ZnO-D-CNT, (**c**) ZnO-D-PEDOT:PSS, (**d**) ZnO-A, (**e**) ZnO-A-CNT and (**f**) ZnO-A-PEDOT:PSS.

The high-resolution spectra of the O 1s region of the ZnO samples, ZnO-CNT hybrids and ZnO-PEDOT:PSS hybrids in Figure 4 consist of several peaks, including lattice oxygen peak of ZnO or the Zn-O bond. Additionally, for the samples ZnO-D (Figure 3a), ZnO-D-CNT (Figure 3b), ZnO-A (Figure 3d) and ZnO-A-CNT (Figure 3e), photoelectron peaks that correspond to hydroxyl groups are also visible. In particular, the photoelectron peak at around 531.5 eV is attributed to Zn-OH bonds as well as oxygen vacancies [22]. TEM analysis estimated an average ZnO nanoparticle size of 5 nm, implying a very high surfaceto-volume ratio. In such small nanoparticles, surface oxygen vacancies are prevalent. Since the C 1s region contains oxygen or hydroxyl components and the O 1s region contains carbon and hydroxyl components, it therefore suggests that hydroxyl groups are responsible for the decoration of CNT with ZnO. This directly implies that hydroxyl groups enable the anchoring of ZnO on CNT surface through covalent bonding with carbon, as there is no indication of Zn-C bonds. TEM images clearly indicate that ZnO nanoparticles grow directly on the CNT sidewalls through Zn-O/OH-C bonds. Furthermore, the O 1s region of the ZnO-PEDOT:PSS hybrids of Figure 4c,f display additional peaks, along with differences in relative intensities of peaks compared to ZnO and ZnO-CNT hybrids. In these samples, the characteristics of the PEDOT:PSS polymer is more dominant. In fact, two peaks—C-O-C of PEDOT at 532.7 eV and O=S of PSS at 531.7 eV—are visible as well as a third peak of Zn-O [29]. In general, the ZnO-D lattice, i.e., as-synthesized ZnO-D or ZnO-D in the nanohybrids, shows a more stable oxygen component, when considering the C 1s spectra of Figure 5a–c, where the relative intensities of O=C-O, C-OH and C-C

peaks are rather constant. However, for the ZnO-A-PEDOT:PSS, the PEDOT peak is less intense than the PSS peak. In fact, PEDOT:PSS macromolecule consists of PEDOT that is positively charged, highly conductive, and hydrophobic. On the other hand, PSS is negatively charged, insulating and hydrophilic. If we consider that the nanoparticles were dispersed in an aqueous solution of PEDOT:PSS, then the adsorption of hydroxyl groups on the surface of the ZnO nanoparticles is inevitable. From the relative intensities of various peaks of the O1s region in Figure 4a,d, ZnO-A tends to adsorb a higher quantity of hydroxyl groups than ZnO-D. Consequently, the surface of ZnO-A is more electronegative with a propensity to the positively charged PEDOT. The O 1s region of ZnO-A-PEDOT:PSS consists of a less intense C-O-C peak and a highly intense O=S peak compared to ZnO-D-PEDOT:PSS. The reduction in the relative intensity of the C-O-C peak suggests that either PEDOT was removed or degraded on adding ZnO-A. In addition, the shift in the Zn-O and C-O-C peaks to higher binding energies confirms the formation of a covalent bond between the C of PEDOT and OH groups present on the ZnO surface. The higher binding energy of the Zn-O peak along with an increase in its intensity indicates that the configuration for the lattice oxygen of ZnO-A becomes more stable.

**Figure 4.** High-resolution XPS spectra of the O 1s region, (**a**) ZnO-D, (**b**) ZnO-D-CNT, (**c**) ZnO-D-PEDOT:PSS, (**d**) ZnO-A, (**e**) ZnO-A-CNT and (**f**) ZnO-A-PEDOT:PSS.

The vibrational properties of the ZnO nanoparticles and their hybrids were investigated using Raman spectroscopy. The results obtained from Raman spectroscopy complement those obtained via XPS. In fact, chemical and structural changes can be evaluated simultaneously on ZnO and CNT or PEDOT:PSS using Raman spectroscopy. Figure 5a compares the different vibrational modes obtained from these samples in the range of 100–800 cm−1. The first-order phonon modes obtained at ~440 cm−1, ~585 cm−<sup>1</sup> and ~667 cm<sup>−</sup>1, correspond to E2H, E1 (LO) and E2 (TO) modes, respectively [21]. Other modes obtained at ~320 cm−<sup>1</sup> and 506 cm−<sup>1</sup> are multiphonon scattering modes that correspond to the E2H–E2L and E1(TO) + E2L modes, respectively [30]. The E2H, E2H–E2L, E1 (LO) modes involve the oxygen component of ZnO. More specifically, the E2H at 440 cm−<sup>1</sup> corresponds to lattice oxygen, whereas the E1 (LO) corresponds to oxygen-related defects [31]. For all the samples, the E2H mode intensities are high, implying that the ZnO lattice structure is unaffected on hybridizing with CNT or PEDOT:PSS. However, the relative intensity of the E1 (LO) band increases in the nanocomposites, indicating an increased number of surface defects [32]. The attachment of ZnO on the sidewalls of the CNT through hydroxyl functional groups indicates that the interfacial region and, therefore, the surface of ZnO are highly defective. Additionally, the E2H peak for ZnO-PEDOT:PSS samples has shifted to

a higher wavenumber of 445 cm−<sup>1</sup> owing to chemical interactions between PEDOT and ZnO. This peak is more intense for ZnO-A-PEDOT:PSS than ZnO-D-PEDOT:PSS, which once again supports that ZnO-A has a more stable lattice configuration in PEDOT:PSS.

Raman signatures lower than 300 cm−<sup>1</sup> are assigned to the vibrations of Zni, and those above 300 cm−<sup>1</sup> are assigned to the vibrations of oxygen atoms [33]. The peak at 275 cm−<sup>1</sup> has been attributed to Zni or Zni clustering [34,35]. The intensity of this mode increases relative to the other modes in the CNT-based nanocomposites and is the highest for PEDOT:PSS-based nanocomposites. This suggests that the amount of Zni is higher than Vo in the hybrid samples. Another mode at 526 cm−<sup>1</sup> is observed for the PEDOT:PSS nanocomposites, corresponding to the combination of Vo and Zni [34]. A lower-intensity peak at the same localization is also visible in the ZnO-CNT-based samples. In general, the relative intensity of this combined mode increases in the hybrid samples owing to an increase in Zni. In Figure 5b, Raman bands from 1200–1800 cm−<sup>1</sup> of pristine CNT are compared to those of ZnO-CNT. The D-band at 1341 cm−<sup>1</sup> for pristine CNT redshifts for ZnO-CNT to ~1351 cm<sup>−</sup>1. A similar redshift in the G-band from 1579 cm−<sup>1</sup> to ~1592 cm−<sup>1</sup> is also observed. These redshifts further confirm the presence of oxygen or hydroxyl groups on the CNT surface [36]. The (\*) marked peaks in ZnO-CNT samples are assigned to C-O bond vibrations from the acetate precursor used during synthesis [21]. Infrared spectroscopy studies of hydrogen adsorption on ZnO suggest that OH and H are adsorbed simultaneously [37]. In fact, the dissociation of hydrogen followed by its adsorption manifests as a change in the corresponding vibrational frequency, including stretching vibrations of Zn-H and O-H, which are very different from the free hydroxyl group vibrational frequency [38]. However, hydrogen adsorption is more likely on prismatic surfaces, implying that facetted ZnO nanoparticles would be more susceptible to hydrogen adsorption [39]. However, for successful hydrogen adsorption, firstly, a more acidic environment is required when working in aqueous media, or a high pressure when working in gaseous media. In addition, the nanoparticles presented in this study are spherical and not facetted. In our case, the NaOH-rich conditions provide a basic environment that is advantageous to the adsorption of hydroxyl groups, further promoted by the presence of Vo.

Figure 5c shows the Raman spectra of ZnO-PEDOT:PSS samples in the range (900–1700 cm−1), where the contributions from PSS and PEDOT vibrational modes are the most significant. Two typical PSS vibrational modes at 988 cm−<sup>1</sup> and 1097 cm−<sup>1</sup> are observed [40]. The vibrational modes of PEDOT observed at 1263 cm−1, 1369 cm−1, 1436 cm−<sup>1</sup> and 1517 cm−<sup>1</sup> correspond to Cα-Cα, Cβ-Cβ, symmetrical Cα=C<sup>β</sup> and asymmetrical Cα=C<sup>β</sup> stretching vibrational modes, respectively. In the ZnO-D-PEDOT:PSS samples, the symmetrical vibrational mode at 1436 cm−<sup>1</sup> is redshifted compared to the pristine PEDOT:PSS (~1440 cm−1) [41]. However, this mode is slightly more redshifted in the ZnO-A-PEDOT:PSS, suggesting a slightly higher benzoid (coil) to quinoid (linear) structural transition [41,42]. The PEDOT chains of linear conformation tend to increase the conductivity of the polymer due to a stronger covalent bonding with ZnO. Additionally, asymmetrical Cα=Cβ bonds of PEDOT have similar intensities for both samples, whereas, Cα-C<sup>α</sup> and Cβ-C<sup>β</sup> bonds for sample ZnO-D-PEDOT:PSS are more intense than sample ZnO-A-PEDOT:PSS. On the other hand, the asymmetrical Cα=C<sup>β</sup> bond of PEDOT at 1517 cm−<sup>1</sup> is more intense for ZnO-A-PEDOT:PSS samples. This implies that the bonds in the PEDOT chain have undergone structural modification provoking a breakdown in symmetry. This again suggests that PEDOT was degraded or removed from the macromolecule upon combining with ZnO-A.

**Figure 5.** Normalized Raman spectra of (**a**) ZnO-D, ZnO-A, ZnO-D-CNT, ZnO-A-CNT, ZnO-D-PEDOT:PSS and ZnO-A-PEDOT:PSS from 100-800 cm−1, (**b**) pristine CNT, ZnO-D-CNT and ZnO-A-CNT from 1200-1800 cm<sup>−</sup>1, and (**c**) ZnO-D-PEDOT:PSS and ZnO-A-PEDOT:PSS from 900–1600 cm<sup>−</sup>1.

#### *3.2. Optical Properties*

The band gaps of the ZnO samples, ZnO-CNT hybrids and ZnO-PEDOT:PSS hybrids were calculated via UV-Vis absorption spectroscopy followed by Tauc plots, presented in Figure 6. The band gaps of these samples range from 3.11 to 3.3 eV, which correspond to the theoretical band gap of ZnO, implying that the absorbance in the nanocomposites is dominated by ZnO. Depending on the synthesis routes, variations in the band gaps of ZnO have been observed [43]. The absorption spectra of ZnO samples revealed a sharp shoulder at ~3.3 eV, stretching down to 2.0 eV, whereas, a broader shoulder at ~3.3 eV stretching down to ~1.5 eV was observed for the CNT hybrids [44].

**Figure 6.** Tauc plot of (**a**) ZnO-D and ZnO-D-CNT, (**b**) ZnO-D-PEDOT:PSS, (**c**) ZnO-A and ZnO-A-CNT, and (**d**) ZnO-A-PEDOT:PSS.

The emission properties of the as-synthesized ZnO samples, ZnO-CNT hybrids, and ZnO-PEDOT:PSS hybrids were investigated at room temperature. A 365 nm (3.4 eV) excitation source was used to induce band-to-band transitions in these samples with band gaps between 3.11 eV and 3.3 eV. The PL spectra in Figure 7a–f present typical PL emission characteristics of ZnO nanoparticles, consisting of the NBE and DLE [45]. The similarities in emission peak localizations indicate that the emissions mainly originate from ZnO nanoparticles, for both the freestanding and hybrid nanocomposites. However, there are significant changes in the overall quantum efficiencies and intensities of certain emission peaks of the hybrid samples. This suggests that interfacial bonding between ZnO nanoparticles and CNT or PEDOT:PSS via OH groups plays an important role in excitonic separation and recombination. The probable origin of the DLE is the combination of several point defects, such as oxygen interstitials (Oi), oxygen vacancies (VO), zinc vacancies (VZn), zinc interstitials (Zni), and their complexes [22,46] that are related to the presence of hydrates in the ZnO precursor. In addition, the NBE to DLE ratio is useful in evaluating the crystalline quality of ZnO. Moreover, an average nanoparticle size of 5 nm indicates a high surface-to-volume ratio, which in turn denotes a high amount of surface defects. These surface traps also consist of chemisorbed species, allowing additional radiative or non-radiative recombination mechanisms, which would alter the quantum efficiency of the DLE.

The doubly ionized oxygen vacancy, i.e., VO++, or surface oxygen vacancy at 2.2 eV is dominant in all the samples due to the high surface-to-volume ratio. The 2.2 eV transition is associated with the capture of a hole by VO<sup>+</sup> from surface charges to form VO++ [47]. The single ionized oxygen vacancy VO<sup>+</sup> or volume oxygen vacancy emits at ~2.5 eV. Both types of vacancies produce green luminescence in ZnO. Additionally, the intensity of the green emission can be strongly influenced by free carriers on the surface, especially for nanoparticles with very small sizes [47,48]. Since PL measurements were performed in air, it is likely that hydroxyl groups or oxygen molecules are adsorbed on the surface of the nanoparticles. The chemisorbed oxygen species provoke an upward band bending in the as-synthesized ZnO nanoparticles (Figure 7g), which allows VO<sup>+</sup> to convert into VO++ through the tunneling of surface-trapped holes to deep levels. Therefore, the observed dominant green emission in ZnO-D and ZnO-A samples is mainly surface related. In a previous study, the chemisorption of hydroxyl groups/oxygen species was suppressed by covering the nanoparticles with CNT [17]. In that study, non-functionalized CNT were used, and a successful passivation of surface states was obtained. In the present case, the CNT were functionalized with OH functional groups, as discussed previously. Therefore, in the present case, the upward band bending is enhanced (Figure 7h). The increased upward band bending leads to further increase in the depletion region size, whereupon the probability of electron capture at the defect sites increases. This mechanism also reduces the probability of band-to-band transitions and the NBE is diminished.

For ZnO-PEDOT:PSS samples, a complete coverage of ZnO with PEDOT:PSS is visible in the TEM images. In general, there is a reduction in the overall emission compared to the as-synthesized and ZnO-CNT samples, due to the low amount of ZnO nanoparticles. However, the NBE-to-DLE ratio is higher in these samples. The increase in NBE can be attributed to the reduced surface hydroxyl groups, relative to the as-synthesized and ZnO-CNT samples, leading to lower upward band bending (Figure 7i). More particularly, the NBE-to-DLE ratio is higher for the ZnO-D-PEDOT:PSS than ZnO-A-PEDOT:PSS. An increase in DLE for the latter can be attributed to the higher amount of hydroxyl groups present on the ZnO-A sample, as assessed on the basis of the XPS studies, leading to a slightly higher upward band bending than for ZnO-D. Additionally, the NBE of both types of hybrid sample, i.e., CNT and PEDOT:PSS, has redshifted, suggesting an increased amount of Zni, further corroborating the Raman spectroscopy results. Finally, the red emission at ~1.75 eV is the least significant component in the PL spectra, associated mainly with VZn-related defects [49].

**Figure 7.** PL emission spectra of (**a**) ZnO-D, (**b**) ZnO-D-CNT, (**c**) ZnO-D-PEDOT:PSS, (**d**) ZnO-A, (**e**) ZnO-A-CNT and (**f**) ZnO-A-PEDOT:PSS. Schematics of upward band bending (**g**) chemisorbed Oxygen species, (**h**) CNT decorated with ZnO through hydroxyl groups, (**i**) ZnO-PEDOT:PSS nanohybrids.

#### **4. Conclusions**

In this study, we successfully synthesized ZnO nanoparticles and their hybrids containing CNT and PEDOT:PSS. The effect of hydroxyl groups on the optical properties of ZnO nanoparticles and their hybrids was investigated. The ZnO nanoparticles display optical properties that are both bandgap and defect related. In addition, the choice of precursor immensely influences the overall properties of the nanoparticles and their hybrids. During the synthesis of the hybrid nanocomposites, hydroxyl groups adhere to the surface of the ZnO nanoparticles, and in turn, intensify the defect related emission. These hydroxyl groups are necessary for the successful decoration of the CNT and the incorporation of ZnO in the PEDOT:PSS matrix. Additionally, the linear transformation of PEDOT from the coil structure implies a more conductive polymer, which would enhance the I-V characteristics of the nanocomposite. However, ZnO-A tends to degrade the PEDOT:PSS macromolecule by removing or degrading the conducting PEDOT, which could prove detrimental to the electrical properties of the nanocomposite. Future research consists of incorporating these nanoparticles and their hybrids in LED applications. This present work aids in understanding the modification of the physical and chemical properties of the ZnO nanoparticles when hybridized to PEDOT:PSS and CNT. On the basis of this work, we conclude that ZnO-D and its hybrid nanocomposites, synthesized with hydrate precursors show higher stability and are likely to offer better electrical conductivity when used in LED.

**Author Contributions:** Conceptualization: K.N., P.R.; methodology: K.N., P.R., E.R.; validation, P.R., E.R. and E.E.; formal analysis, K.N., E.E., M.R.S.; investigation, K.N., P.R., E.R.; resources, E.E., P.R., E.R.; data curation, K.N.; writing—original draft preparation, K.N.; writing—review and editing, K.N., P.R. and E.R.; supervision, P.R. and E.R.; project administration, P.R.; funding acquisition, P.R. and E.R. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research has been supported by the European Regional Development Fund project grant number TK134 "EQUiTANT" and T210013TIBT "PARROT mobility program". We thank EU-H2020 research and innovation program under grant agreement no. 1029 supporting the Transnational Access Activity within the framework NFFA-Europe.

**Data Availability Statement:** Not applicable.

**Conflicts of Interest:** The authors declare no conflict of interest.

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### *Article* **Optimizing the PMMA Electron-Blocking Layer of Quantum Dot Light-Emitting Diodes**

**Mariya Zvaigzne 1,\*, Alexei Alexandrov 2, Anastasia Tkach 1, Dmitriy Lypenko 2, Igor Nabiev 1,3,4 and Pavel Samokhvalov 1,\***


**Abstract:** Quantum dots (QDs) are promising candidates for producing bright, color-pure, costefficient, and long-lasting QD-based light-emitting diodes (QDLEDs). However, one of the significant problems in achieving high efficiency of QDLEDs is the imbalance between the rates of charge-carrier injection into the emissive QD layer and their transport through the device components. Here we investigated the effect of the parameters of the deposition of a poly (methyl methacrylate) (PMMA) electron-blocking layer (EBL), such as PMMA solution concentration, on the characteristics of EBLenhanced QDLEDs. A series of devices was fabricated with the PMMA layer formed from acetone solutions with concentrations ranging from 0.05 to 1.2 mg/mL. The addition of the PMMA layer allowed for an increase of the maximum luminance of QDLED by a factor of four compared to the control device without EBL, that is, to 18,671 cd/m2, with the current efficiency increased by an order of magnitude and the turn-on voltage decreased by ~1 V. At the same time, we have demonstrated that each particular QDLED characteristic has a maximum at a specific PMMA layer thickness; therefore, variation of the EBL deposition conditions could serve as an additional parameter space when other QDLED optimization approaches are being developed or implied in future solid-state lighting and display devices.

**Keywords:** quantum dots; QDLED; electron-blocking layer; PMMA

#### **1. Introduction**

Fluorescent semiconductor nanocrystals (NCs) or quantum dots (QDs) have plenty of advantageous properties, such as the possibility of tuning the luminescence wavelength by varying the physical size of the NCs and the capacity for forming stable colloidal solutions, which makes it possible to obtain coatings by inexpensive solution-process methods, and make QDs promising materials in optoelectronic, bioimaging, lighting, and other applications [1–5]. Today, light-emitting devices (LEDs) based on organic compounds (OLEDs) prevail in commercial lighting and display appliances. Quantum dots outperform traditional organic dyes in terms of the width of the absorption spectrum, molar extinction, and photostability. Thus, quantum dots are expected to be promising candidates to overcome the material stability issues typical of OLEDs, such as drastic efficiency roll-off at high current densities and mediocre operational lifetimes. Moreover, due to their inorganic nature, QDs are much more thermally stable materials, which makes it possible to increase the brightness of QD-based LEDs by increasing the current density in the device.

**Citation:** Zvaigzne, M.; Alexandrov, A.; Tkach, A.; Lypenko, D.; Nabiev, I.; Samokhvalov, P. Optimizing the PMMA Electron-Blocking Layer of Quantum Dot Light-Emitting Diodes. *Nanomaterials* **2021**, *11*, 2014. https:// doi.org/10.3390/nano11082014

Academic Editors: Protima Rauwel and Erwan Rauwel

Received: 12 July 2021 Accepted: 4 August 2021 Published: 6 August 2021

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**Copyright:** © 2021 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https:// creativecommons.org/licenses/by/ 4.0/).

In addition, QDs have quite narrow fluorescence and electroluminescence spectra and, hence, are promising components of displays or illuminators with a wide color gamut. To fully exploit the superior properties of QDs, a number of QD-based LED (QDLED) structures with different device and material configurations have been developed [6]. A typical QDLED-emitting layer represents a thin QD film sandwiched between two chargetransport layers, and its interaction with them may cause luminescence quenching through various nonradiative pathways, such as the Auger recombination [7,8], QD charging [7], and charge and/or energy transfer from QDs to the charge-transport materials [7,9], and so forth.

One of the main shortcomings of QD-based LEDs is imbalance between the rates of charge carrier injection and transport [7], which leads to the formation of excess charges (electrons or holes) in the emitting QD layer and quenching of QD radiation due to the aforementioned nonradiative processes. This phenomenon leads to a significant decrease in radiation efficiency, especially at high current densities, and, hence, to overall low performance of the QDLEDs. In most modern QDLED configurations [10], this imbalance mainly results from a larger potential barrier for the injection of holes into the QD layer than for the injection of electrons, as well as a higher mobility of electrons in the electron transport layer (ETL), usually based on ZnO, compared to the hole mobility in organic hole transport layers (HTL) [11]. One of the approaches to solving this problem is the introduction of an electron-blocking layer (EBL). The EBL materials and methods of its integration into the QDLED structure vary widely. For example, it has been shown that the addition of a 4,4,4 tris(N-carbazolyl)-triphenylamine (TcTa) EBL between the hole-transporting layer and the light-emitting QD layer [12] or a combined hole-transporting and electron-blocking layer of deoxyribonucleic acid (DNA) complexed with cetyltrimetylammonium (CTMA) [13] enhances efficiency due to the reduction of electron overflow and improvement of hole injection. Another approach is to insert an EBL between the ETL and the QD layer to restrict the flow of electrons. In the framework of this approach, Al2O3 [14] and poly(methyl methacrylate) (PMMA) [15] have been demonstrated to be effective materials for the EBL. However, a thin film of poly (methyl methacrylate) (PMMA) is more often used as an EBL [15,16], because the positions of its energy levels provide a high potential barrier for electron injection into the emitting layer for most types of QDs. In addition, PMMA is soluble, such as in acetone, which makes it possible to apply PMMA onto the underlying QD layer without partially dissolving or deforming the latter.

The first introduction of a PMMA electron blocking layer into the QDLED structure was reported by Dai et al. [15], who fabricated QDLEDs with a 6 nm PMMA insulating layer between a CdSe/CdS core/shell QD layer and a ZnO ETL to optimize the charge balance in the device. They compared the hole mobility in an HTL consisting of poly-TPD (1·10−<sup>4</sup> cm2·V−1·s−1) and PVK (2.5·10−<sup>6</sup> cm2·V−1·s−1) with the electron mobility in an ETL based on a ZnO nanocrystal film (∼1.8·10−<sup>3</sup> cm2·V−1·s−1), and concluded that the insertion of the PMMA layer may lead to excess electron injection into the QD emissive layer. To confirm this assumption, they measured and compared the current densities of the electron-only devices (ITO/Al/QDs/ZnO/Al) and hole-only devices (ITO/PEDOT:PSS/poly-TPD/PVK/QDs/Pd). In this case, the addition of the PMMA layer between the ETL and the QD layer to optimize charge balance did not cause any considerable changes in either the turn-on voltage or the brightness in comparison with the control QDLEDs without the PMMA layer. However, for equally bright QDLEDs, the current density in the control device was much greater, which indicated that the efficiency of this device was substantially lowered by the excess electron current. Thus, the efficiency of the EBL-based approach in terms of enhancing the performance of QDLEDs was confirmed. In addition, the stability of the devices without the PMMA layers was relatively poor, and it was improved 20-fold by adding the PMMA EBL.

Rahmati et al. [16] presented a new QDLED architecture with multiple PMMA EBLs sandwiched between a pair or more of QD layers. The authors developed QDLED structures with one, two, and three PMMA layers and showed that a device containing two

PMMA and three QD layers had the best current efficiency of 17.8 cd·A−<sup>1</sup> and a luminance of 194,038 cd·m−2. The substantial improvement of QDLED performance was mainly attributed to the addition of the PMMA EBL, which reduced the backward electron leakage from the active QD region and enhanced electron confinement, leading to an increased electron concentration in the QD active layers and a higher radiative recombination rate. It is worth noting that the aforementioned QDLED configuration where the EBL is sandwiched between a pair of QD EMLs could be employed in the design of white QDLEDs by combining isolated blue/green/red QD layers separated by two PMMA spacers into a complex emissive layer [16].

Although it has been demonstrated that adding a PMMA layer to the QDLED structure may significantly improve the performance of QDLEDs, this approach needs further optimization and detailed study to rationally exploit the PMMA electron-blocking capacity. Here, we studied the correlations between the parameters of the PMMA EBL deposition, such as the PMMA solution concentration, and the most important performance characteristics of QDLEDs.

#### **2. Materials and Methods**

#### *2.1. Synthesis of CdSe/ZnS/CdS/ZnS (CdSe/MS) QDs*

The synthesis of CdSe cores with a diameter of 2.3 nm was carried out by the hot injection technique using cadmium hexadecylphosphonate and trioctylphosphine as precursors at a temperature of 240 ◦C; the procedure is described in more detail in [17]. After the synthesis, the separation and purification of the CdSe cores were carried out by means of reprecipitation of nanocrystals and subsequent gel permeation chromatography, after which their surface was treated with oleylamine in the presence of sodium borohydride to replace the hexadecylphosphonic acid residues with oleylamine, which facilitated further growth of inorganic shells. Then, after additional purification steps, the CdSe cores were placed into a reaction mixture of 1-octadecene and oleylamine (1:1, *v*/*v*) for growing the shells. After accurate quantification of CdSe cores in the reaction solution using the approach described in [18], we calculated the quantities of precursors required for obtaining QDs with the desired shell structure. The shells were grown at 170 ◦C in an argon atmosphere at an average growth rate of 1 monolayer per 30 min. After the synthesis and isolation of QDs from the crude solution, the organic ligands were replaced with hexadecylammonium palmitate (HDA-PA), which reduced the sensitivity of the optical properties of QDs to atmospheric exposure during long-term storage. The luminescence and absorption spectra, as well as the transmission electron microscopy (TEM) image of the obtained QDs are shown in Figures S1 and S2, provided in the Supporting Information (SI) file.

#### *2.2. Fabrication of QDLED Devices*

Glass substrates with an indium tin oxide (ITO) layer were preliminarily cleaned by treatment in an ultrasonic bath and then in oxygen plasma. Then, a hole-injecting layer of PEDOT:PSS (poly(3,4-ethylenedioxythiophene): poly(styrene sulfonate)) was deposited on the substrates by spin-coating at 2000 rpm, followed by annealing at 110 ◦C for 10 min. The film thickness was 40 nm. The substrates coated with a PEDOT: PSS layer were transferred into a glove box containing argon (O2 < 1 ppm, H2O < 1 ppm). Next, hole transport layers of poly-TPD (poly(N,N -bis-4-butylphenyl-N,N -bisphenyl) benzidine, solution in chlorobenzene, 8 mg/mL) and PVK (poly (vinylcarbazole), solution in o-xylene, 1.5 mg/mL) were deposited alternately by spin-coating at 2000 rpm. Layers of poly-TPD (30 nm) and PVK (5 nm) were annealed at 100 ◦C for 10 min before applying the next layer. Then, a QD layer was applied from a solution in n-octane (20 mg/mL) by spin-coating at 2500 rpm and annealed at 100 ◦C for 10 min. The QD film thickness was 40 nm. The following PMMA (Sigma Aldrich, Saint Louis, MO, USA, average Mw ~ 120,000 Da) layer was applied by spin-coating from a solution in acetone at 3500 rpm and then annealed at 100 ◦C for 10 min. The concentration of the acetone solution of PMMA varied from 1.2 mg/mL to 0.05 mg/mL to obtain different blocking-layer thicknesses. A 50 nm electrontransport layer (ETL) was applied from solutions of ZnO nanoparticles in isopropyl alcohol (25 mg/mL) by spin-coating at 1000 rpm followed by annealing at 60 ◦C for 10 min. Finally, an aluminum cathode with a thickness of 80 nm was deposited onto the ETL through a shadow mask by thermal evaporation in vacuum (2·10−<sup>6</sup> mbar).

#### *2.3. Instrumental Methods*

The luminescence spectra were measured using an Agilent Cary Eclipse spectrofluorimeter. The absorption spectra were measured using an Agilent Cary 60 UV-Vis spectrophotometer. Transmission electron microscopy (TEM) images were obtained on a JEOL JEM-2100F (JEOL Ltd., Tokyo, Japan) instrument operated at 200 kV acceleration voltage. TEM specimens were prepared by drop-casting a solution of QDs in hexane onto carbon/Formvar-coated 200 mesh copper TEM grids. The voltage–current and voltage– brightness characteristics were measured with a Keithley 2601 SourceMeter 2601 (Keithley Instruments, Inc., Solon, OH, USA), a Keithley 485 picoampermeter (Keithley Instruments, Inc., Solon, OH, USA), and a TKA-04/3 luxmeter–brightness meter (Scientific Instruments "TKA", St. Petersburg, Russia). The preparation of QDLED samples and measurements of their characteristics were performed at room temperature in an argon atmosphere. The film thicknesses were determined by ellipsometry using an MII-4 interferometer ("LOMO", St. Petersburg, Russia) and by means of a MultiMode V (Bruker Corporation, Billerica, MA, USA) atomic force microscope.

#### **3. Results and Discussion**

Deposition of an ultimately thin PMMA charge-blocking layer on top or within the emissive QD layer by means of solution processes requires the minimum possible distortion of the QD layer to achieve efficient performance of the QDLED. Thus, selection of the appropriate solvent for PMMA is of utmost importance. PMMA can be dissolved in a number of organic nonpolar (toluene, chloroform, etc.) and weakly polar (acetone) solvents. In most QDLED configurations, QDs capped with long-chain aliphatic ligands are deposited from nonpolar solvents, such as octane and toluene. Therefore, acetone becomes the solvent of choice for deposition of PMMA because, being polar, it cannot dissolve the underlying QDs and, at the same time, being only mildly polar, it cannot cause severe wrapping or cracking of the underlying thin film of QDs. On the other hand, acetone is quite volatile and has a low viscosity; therefore, it is hard to control the parameters of its deposition by spin-coating other than the concentration of PMMA in solution.

The structure of the fabricated QDLED devices is illustrated in Figure 1 along with the schematic of the flat-band energy level diagram of the layers in the device.

We used QDs with the core/multishell structure to eliminate the negative effects on QD fluorescence caused by Auger recombination and surface-trapping [19–21], and because our previous studies revealed that this type of QD possessed the optimal characteristics of photostability due to the suppression of charge transfer [4]. In order to increase the efficiency of hole injection into the QD-emitting layer, we added a poly-TPD/PVK bilayer-structured hole-injection layer between QDs and PEDOT:PSS. This configuration creates a gradual step transition between the hole energy levels of QDs and PEDOT:PSS hole-transport layer. A thin layer of ZnO nanoparticles was deposited as an electron transport layer (ETL) because ZnO proved to be the most favorable ETL material due to its high transparency, low work function, and high electron mobility [6,10,15]. To investigate the effect of the PMMA layer preparation routine on the QDLED performance, a series of devices were fabricated. To do this, we varied the concentration of the PMMA solution in acetone from 0.05 to 1.2 mg/mL. Unfortunately, we were unable to measure the exact thickness of the PMMA EBL using the available AFM instrumentation, because the measurement error was higher than the measured value. Otherwise, we may roughly estimate the thickness range of PMMA EBL in our devices as 0.13–3 nm, corresponding to the lower and upper limits of the solution concentrations, respectively. The details of this estimation are given in the SI, the results of the calculation are given in the Table S1. Yet, in

the following sections, we prefer to stick to the known experimental concentration values rather than our rough thickness estimations.

**Figure 1.** Flat-band energy level diagram of the fabricated QDLEDs (**a**) and electroluminescence spectrum of QDLED with the PMMA electron-blocking layer deposited from a solution with a PMMA concentration of 0.2 mg/mL (**b**). The insets in panel (**b**) show photographs of the device under ambient light (left) and operated at 6 V (right).

Figure 2a shows the current density–voltage and luminance–voltage characteristics of the devices under investigation. As can be seen, the addition of a PMMA blocking layer in most cases led to an accelerated rise and an overall increase in the current density at all voltages relative to the structure without a blocking layer. Only when we applied the EBL of PMMA from a solution with the highest concentration (1.2 mg/mL) did we observe a drop in this characteristic. These results might be counterintuitive at first glance, because blocking of electron flux through the device by a potential barrier created by the PMMA layer led to an overall increase in the current flow through the whole device. However, previous studies showed that the imbalance between the electron and hole currents led to the accumulation of excess electrons inside the device and interfacial charging [7,11], which, in turn, acts as a counter-driving force for electron currents and leads to less efficient electron transport and injection. EBL, in this case, diminishes the charge flow imbalance and allows the device to operate in the optimal regime, since the PMMA interlayer provides quite a high energy barrier of around 3 eV against electron flow from the ETL (Figure 1). Thus, PMMA EBL can block excess electron flow from ETL to the QD light-emitting layer by reducing the electron current density, which leads to the improved charge carrier balance inside the emissive QD layer, as it was shown for a number of other EBL materials [22–25].

As can be seen from our data, this optimal charge carrier balance was achieved when EBL was deposited from a solution with a PMMA concentration in the range of 0.1–0.4 mg/mL. In this range, higher PMMA concentrations yield devices with a lower current density but similar luminance. On the other hand, when the PMMA concentration in deposition solution was lowered to 0.05 mg/mL, we observed current leakage even at low voltages (0–2.5 V), which suggested a short circuit due to disturbance of the emissive layer during EBL deposition.

**Figure 2.** Current density (**a**) and luminance (**b**) versus voltage characteristics of QDLED samples employing a PMMA EBL deposited from PMMA solutions in acetone with different concentrations.

A similar trend was observed for the luminance–voltage characteristics (Figure 2b). In this case, the luminance saturation plateau was reached faster, and the brightness values were higher in QDLED structures fabricated with a blocking PMMA layer. As an exception, QDLED samples employing a PMMA EBL deposited from solutions with concentrations of 0.8 and 1.2 mg/mL exhibited only minor, if any, improvement of this characteristic. In the case of a 0.8 mg/mL solution, sharper growth was observed, but the brightness value did not exceed that for the device without an EBL. These effects may arise from hindered injection of electrons into the emitting layer due to an increase in the thickness of the potential barrier and, as a consequence, a decrease in the probability of carrier tunneling.

The performance parameters of all fabricated QDLED devices are summarized in Table 1. The lowest turn-on voltage of 2.1 V was observed for the two lowest PMMA solution concentrations, 0.1 and 0.05 mg/mL. At the same time, the QDLED structure without a PMMA layer had one of the highest turn-on voltage values, 3.3 V. In general, a distinct minimum was observed in the plot of the turn-on voltage versus the PMMA solution concentration (Figure 3).


**Table 1.** Summary of the performance parameters of the fabricated QDLED devices with and without a PMMA layer deposited from PMMA solutions with different concentrations.

In terms of the maximum current efficiency, the QDLED with an EBL deposited from a 0.4 mg/mL PMMA solution turned out to be the optimal one (Table 1). An increase in PMMA concentration led to a sharp drop of the current efficiency, while its decrease also resulted in a 1.5-fold lower current efficiency. For concentrations of 0.2 and 0.1 mg/mL, there were no significant differences in either current efficiency or turn-on voltage. However, the brightness steadily increased with decreasing PMMA concentration in the EBL deposition procedure. Thus, the maximum brightness in our experiment was 18 671 cd/m2, obtained in the case of QDLEDs with an EBL fabricated using a 0.05 mg/mL PMMA solution. This luminance value was four times higher than that for devices without a blocking layer. Figures 4 and 5 show the dependences of the current efficiency and luminance at 9 V on the PMMA solution concentration.

**Figure 3.** Effect of PMMA solution concentration on the turn-on voltage value of the QDLED device.

**Figure 4.** Effect of the PMMA solution concentration on the current efficiency of the QDLED device at 9 V.

**Figure 5.** Effect of the PMMA solution concentration on the luminance of the QDLED device at 9 V.

As can be seen, in the case of current efficiency, the apparent maximum is observed for the device where the EBL had the minimum thickness, when it was deposited from a solution with a PMMA concentration of 0.05 mg/mL, and for the device without EBL. Additionally, a local maximum was observed for the device whose EBL was formed from a 0.4 mg/mL PMMA solution. This result suggests that either excess electron injection or over-blocking of the electron current deteriorates charge balance in the QDLEDs and thereby degrades current efficiency values.

Regarding the luminance (Figure 5), addition of even the thinnest PMMA layer to the QDLED led to a drastic increase in this characteristic, apparently due to reducing the probability of the formation of excess charges in the QD emissive layer and preventing the luminescence quenching via nonradiative processes. However, further increase in the concentration of PMMA in the EBL deposition solution resulted in deterioration of this characteristic. This may have been due to the hindered injection of electrons into the emitting layer as a result of an increased thickness of the potential barrier and, as a consequence, a decreased probability of carrier tunneling.

Our findings show that the most important characteristics of QDLEDs can be substantially improved by careful adjustment of the PMMA EBL deposition parameters, such as PMMA solution concentration. Notably, among the QDLEDs studied here, there was no obvious best device in terms of the turn-on voltage, current efficiency, and luminance. Therefore, the addition of the PMMA as an EBL alone should not be considered as a single treatment to improve all the QDLED characteristics, but it may be quite effective if applied along with other optimization approaches. In this case, the PMMA layer deposition parameters should be adjusted according to the requirements of each specific QDLED structure. Our results may be helpful as guidance for the preparation of a PMMA EBL in order to adjust specific QDLED parameters.

#### **4. Conclusions**

An electron-blocking layer of poly(methyl methacrylate) was added to the standard QDLED structure in order to improve the brightness characteristics and current efficiency. It has been shown that the concentration of the PMMA solution during layer deposition plays a significant role in achieving high QDLED efficiency. Specifically, at a concentration as high as 1.2 mg/mL, the characteristics of the current efficiency and brightness of the QDLEDs dropped significantly relative to a similar device without an EBL. This may be due to the hindered injection of electrons into the emitting layer due to an increase in the thickness of the potential barrier and, as a consequence, a decrease in the probability of carrier-tunneling.

At the same time, a low concentration of the initial PMMA solution leads to a sharp improvement of the characteristics of the QDLEDs, both in terms of brightness and current efficiency and in terms of lowering the turn-on voltage. In terms of current efficiency, the QDLED sample with an EBL deposited from a 0.4 mg/mL PMMA solution turned out to be the optimal one. Apparently, this was why the resultant EBL provided a better balance of the inflow of charge carriers into the QD layer. In the case of the minimum concentration of the PMMA solution, the brightness of the LEDs produced was 18,671 cd/m2, which is four times higher than these values for devices without a blocking layer due to reducing the number of charged QDs and probability of nonradiative processes.

**Supplementary Materials:** The following are available online at https://www.mdpi.com/article/10.339 0/nano11082014/s1. Figure S1: Luminescence and absorption spectra of the synthesized CdSe/ZnS/CdS/ ZnS QDs; Figure S2: TEM image of the synthesized CdSe/ZnS/CdS/ZnS QDs; Formulas regarding the estimation of the thickness of PMMA electron blocking layers; Table S1: Estimated PMMA EBL layer thickness deposited from PMMA solutions in acetone with different concentrations.

**Author Contributions:** Conceptualization, I.N. and P.S.; methodology, M.Z., A.A., A.T. and D.L.; validation, D.L.; writing—original draft preparation, M.Z.; writing—review and editing, P.S. and

I.N.; project administration, P.S. The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript.

**Funding:** This work was supported by the Russian Science Foundation (grant no. 18-19-00588-Π) in its part related to the synthesis of nanomaterials and QDLEDs preparation and by the Ministry of Science and Higher Education of Russian Federation (grant no. FSWU-2020-0035) in its part related to the characterization and validation of created samples.

**Data Availability Statement:** The data presented in this study are available on request from the corresponding authors.

**Acknowledgments:** I.N. acknowledges supports from the Université de Reims Champagne-Ardenne and the Ministry of Higher Education, Research and Innovation of French Republic. We thank Vladimir Ushakov for the help with technical preparation of the manuscript.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**

