Interfacial Engineering of Hydrophobic Montmorillonite for High-Energy-Capability Polypropylene Nanocomposite Dielectrics

Abstract

Polypropylene (PP) dielectric capacitors are crucial for electronics and electric power systems due to their high power density. However, their relatively low energy density limits their practical application in energy storage devices, presenting a long-standing challenge. Montmorillonite (MMT), a natural phyllosilicate mineral abundantly found on earth, features a two-dimensional nanosheet structure and excellent insulating properties. MMT nanosheets have shown promise in enhancing the breakdown strength and energy storage capability of PP dielectric, but compatibility issues with the PP matrix remain a challenge. In this study, we propose a novel surface modification strategy in which polystyrene (PS)-capped MMT (PCM) nanosheets are synthesized through a polymerization–dissolution process. The modified PCM nanosheets demonstrate improved compatibility and are well dispersed within the PP matrix. Optimal loading of the PCM nanosheets effectively dissipate charge energy and hinder the growth of electric trees in the PP matrix. As a result, the PP nanocomposite with 0.2 wt% PCM nanosheets exhibits an enhanced breakdown strength of 619 MV m⁻¹ and a discharged energy density of 4.23 J cm⁻³, with an energy storage efficiency exceeding 90%. These findings provide a promising strategy for the development of high-energy-density dielectric capacitors in an economical manner.

1. Introduction

Polymer dielectric capacitors are essential for advanced pulsed power systems due to their rapid charge–discharge capability, high power density and remarkable reliability [1,2,3,4]. Among various dielectric materials, biaxially oriented polypropylene (BOPP) stands out due to its excellent breakdown strength, minimal dielectric loss and ultrahigh efficiency [5,6]. Nevertheless, the energy storage density (U_e) of polypropylene (PP) is constrained to around 2 J cm⁻³ because of its relatively low dielectric constant (approximately 2.2) [7]. This constraint limits its application in the integration, implementation and miniaturization of cutting-edge electronic devices. Therefore, it is critical to improve the U_e of PP dielectric.

The U_e of a linear dielectric is determined by the equation:

$$U_e=0.5{\epsilon }_r{\epsilon }_0{E}_b^2,$$

(1)

where ε_r is the relative dielectric constant, ε₀ is the vacuum dielectric constant and E_b is the breakdown strength [8,9,10]. In order to effectively improve overall U_e, both ε_r and E_b should be simultaneously enhanced. Considerable efforts have been made to increase ε_r by incorporating inorganic nanofillers [11,12,13]. For example, adding 20 vol% barium titanate (BT) nanoparticles to polyvinylidene difluoride (PVDF) increased ε_r from 10 to 24 [14]. However, substantial increases in ε_r (>20 vol% filler content) often lead to a severe decrease in E_b. For instance, the ε_r of PP was enhanced to 77 with the addition of 55 vol% polyvinylpyrrolidone-modified BT nanoparticles, but the E_b dropped significantly to 140 MV m⁻¹ [15]. As indicated by Equation (1), U_e is directly proportional to ε_r and quadratically dependent on E_b, indicating that enhancing E_b is more efficient than solely increasing ε_r.

Based on the classical theory of electric breakdown in solid dielectrics [16,17,18], a strong electric field can lead to the ionization of dielectric atoms. Electrons acquire energy in these fields and tend to tunnel through the dielectric, gaining enough momentum to collide with other atoms, generating secondary electrons. When the number of free electrons exceeds a threshold, a breakdown occurs. A key strategy to improve E_b is to reduce the kinetic energy of electrons and suppress the formation of secondary electrons. Recent studies have shown that 2-dimensional nanofillers can increase ε_r at low loadings without sacrificing E_b [19,20,21,22,23]. In some cases, low concentrations of 2D nanosheets can even enhance E_b by creating scattering centers and tortuous paths for electrons during the breakdown process. For example, Li et al. [17] introduced negatively charged Ca₂Nb₃O₁₀ (CNO) nanosheets into a PVDF matrix to improve energy storage capability. E_b increased significantly due to the nanosheets’ ability to suppress secondary electrons and block the breakdown path within the polymer matrix. Similarly, Lei et al. [24] prepared a PP nanocomposite with boron nitride (BN) nanosheets. The large aspect ratio of BN nanosheets created effective conduction barriers that restricted the movement of charge carriers and impeded the growth of electric trees during breakdown, thus enhancing both E_b and U_e.

Montmorillonite (MMT) is a widely available phyllosilicate mineral, with a crystal structure consisting of an octahedral aluminum-centered layer sandwiched between two tetrahedral silicon-centered sheets. These layers are connected by tetrahedral oxygen and octahedral hydroxyl groups [25,26,27]. The two-dimensional layered structure of MMT provides a high surface area, which is beneficial for enhancing interfacial polarization in polymer composites [28]. Furthermore, its excellent insulating properties allow MMT to suppress the growth of electric trees within the polymer, thereby improving E_b and the energy density of polymer composites. However, the incompatibility between MMT nanosheets and the PP matrix limits the U_e of the PP composite [29,30].

In this work, a novel polystyrene (PS)-capped MMT (PCM) nanosheet was developed for the first time to enhance the energy storage performance of PP nanocomposites [31]. To achieve this, a polymerization–dissolution approach was employed to covalently attach PS chains onto the MMT nanosheets, which were then blended with a PP matrix. The PCM nanosheets exhibited improved compatibility and were uniformly dispersed in the PP matrix. These well-dispersed PCM nanosheets effectively dissipated the charge energy, suppressed the formation of secondary electrons and impeded the growth of electric trees within the PP matrix, resulting in simultaneous enhancements in ε_r and E_b. Notably, the E_b of the PP nanocomposite with 0.2 wt% PCM nanosheets reached 619 MV m⁻¹, a 33.7% increase compared with pristine PP. An ultrahigh discharged energy density of 4.23 J cm⁻³ was achieved at an electric field of 650 MV m⁻¹, representing a remarkable 125% increase from pristine PP. This study offers an effective and cost-efficient approach to enhancing the U_e of PP-based nanocomposite dielectrics.

2. Materials and Methods

2.1. Materials

Isotactic PP (average Mw ~250,000, average Mn ~67,000) was purchased from Sigma-Aldrich (Shanghai, China). MMT (cation exchange capacity of 130.94 mmol/100 g) was purchased from Guzhang Shan Lin Shi Yu Mineral Co., Ltd. (Xiangxi Autonomous Prefecture, China). Ethanol (99.8%), isopropanol (99.7%), toluene (99%), paraxylene (99%), [3-(methacryloyloxy)propyl]trimethoxysilane (MPS, 97%), sodium dodecyl sulfate (SDS, 99%), styrene, potassium persulfate (K₂S₂O₈, 99%) and sodium hydroxide (NaOH, 97%) were obtained from Shanghai Aladdin Bio-chem Technology Co., Ltd., Shanghai, China.

2.2. Synthesis of MPS-Capped MMT (MCM) Nanosheets

Prior to the synthesis of the PCM nanostructure, the MMT nanosheets were functionalized with MPS, which contains a C=C double bond. In a standard procedure, 50 mg of MMT was initially dispersed in 5 mL of water. Subsequently, 1 M NaOH aqueous solution was added, and the mixture was stirred for 24 h to introduce hydroxyl groups. The resultant MMT was then redispersed in 2 mL of ethanol and combined with 20 mL of isopropanol and 100 μL of MPS. This mixture was refluxed at 80 °C for 2 h. Following this, the MCM nanosheets were separated by centrifugation and redispersed in 2 mL of ethanol in preparation for the subsequent emulsion polymerization process.

2.3. Synthesis of PCM Nanosheets

The MCM nanosheets were first redispersed in a mixture of 15 mL of water and 5 mL of ethanol. Following 30 min of degassing with nitrogen, 1 mL of SDS aqueous solution (3 mg mL⁻¹) and 0.3 mL of styrene were added to the mixture. Once the reaction temperature reached 75 °C, 0.5 mL of K₂S₂O₈ aqueous solution (20 mg mL⁻¹) was introduced into the mixture. The polymerization process continued for 7 h. The resulting PS-coated MMT (PCOM) nanosheets were isolated through centrifugation, washed three times with ethanol and dried at 80 °C. Subsequently, 10 mL of toluene was added to the dried samples to dissolve the PS layer. A well-dispersed suspension was formed rapidly within 5 min, indicating that all MMT nanosheets were successfully redispersed in toluene, which suggests a transfer efficiency of 100%. Finally, the PCM nanosheets were obtained by centrifugation to remove free PS chains in the solution.

2.4. Preparation of PP Nanocomposites

PP nanocomposites were synthesized using a solution blending method. In this process, 1 g of PP was dissolved in 15 mL of paraxylene at 120 °C and stirred for 2 h. Subsequently, a suspension of the PCM in toluene was incorporated into the solution. The mixture was subjected to ultrasonic treatment for 30 min, followed by stirring for an additional 4 h. Afterward, the solvent was removed through freeze-drying, and the resulting nanocomposites were further dried under vacuum at 80 °C overnight. The PP nanocomposites were then hot-pressed into films at 190 °C under a pressure of 25 MPa. The PCM content within the nanocomposites varied between 0.1 wt% and 0.5 wt%, and the resulting films exhibited a uniform thickness of approximately 10 μm.

2.5. Characterization

To study the surface chemical bond of the MMT nanosheets, Fourier transform infrared (FTIR) spectroscopy measurements were carried out by Bruker Tensor α over the range of 400–4000 cm⁻¹. X-ray photoelectron spectroscopy (XPS, ESCALAB 250XI, Thermo Scientific, Waltham, MA, USA) was employed to verify the chemical element composition on the surface of MMT nanosheets. The morphologies and structural details of the MMT nanosheets were examined using scanning electron microscopy (SEM, JSM-7610, JEOL, Tokyo, Japan). X-ray diffraction (XRD, ultima χ, Rigaku, Tokyo, Japan) was performed to examine the crystalline structure and interlayer information of the MMT.

Gold electrodes, approximately 70 nm in thickness and 3 mm in diameter, were magnetron sputtered on both sides of the PP nanocomposite films for electrical measurements. The dielectric properties were assessed using an impedance analyzer (WK6500B, Wayne Kerr, London, UK) at room temperature over a frequency range of 10² Hz to 10⁶ Hz. DC dielectric breakdown strength tests were performed at room temperature on a high-voltage amplifier using the electrostatic pull-down method under a voltage ramp of 500 V/s. High-field displacement–electric (D-E) loops were obtained using a ferroelectric analyzer with a high-voltage source of 10 KV (PK-CPE1801, Polyk, Philipsburg, PA, USA) at a frequency of 100 Hz. All dielectric breakdown and energy storage tests were conducted in silicone oil at room temperature.

2.6. Phase-Field Simulation

The penetration pathway in PP nanocomposite was simulated utilizing a phase-field breakdown model. A scalar phase-field variable s (x, t) was employed to represent the dielectric breakdown characteristic of the PP nanocomposite. The value of s varied from 0 to 1, indicating a completely damaged state and an undamaged state, respectively. For values in between, the permittivity was interpolated as follows:

$$\epsilon \left(s\right)={\displaystyle \frac{{\epsilon }_{ini}}{f\left(s\right)+\eta }},$$

(2)

where $f\left(s\right)=4s^3-3s^4$, ε_ini denotes the initial permittivity, and η is a sufficiently small constant. The overall potential energy of the polypropylene nanocomposite is expressed as:

$$\Pi \left[s,\varphi \right]={\int }_{\Omega }\left[W_{es}\left(E,s\right)+W_d\left(s\right)+W_i\left(𝛻s\right)\right]dV,$$

(3)

In this equation, $W_{es}\left(E,s\right)$ corresponds to the complementary electrostatic energy per unit volume, $W_d\left(s\right)$ represents the damage energy, and $W_i\left(𝛻s\right)$ accounts for the gradient energy term that manages the sharpness of phase boundaries. The evolution of the damage variable s is governed by the linear kinetic law:

$$𝜕s/𝜕t=-m\delta \Pi /\delta s,$$

(4)

where mobility m serves as a material parameter that influences the rate of breakdown propagation within the PP nanocomposite.

3. Results

3.1. Characterization of Microstructure and Surface Composition of MMT

MMT is hydrophilic, while PP is hydrophobic, leading to poor compatibility between them. To improve the dispersion and compatibility of MMT within the PP matrix, a polymerization–dissolution strategy was employed to facilitate the phase transfer of hydrophilic MMT into nonpolar solvents [32,33]. As illustrated in Figure 1, this process involves two surface modification steps: first, the surface is capped with MPS, followed by PS capping. In the experimental procedure, the MMT suspension was treated with an alkaline solution to introduce hydroxyl groups on the MMT surface, which is the key to the silane coupling reaction. MPS was then capped onto the MMT surface to introduce C=C bonds. The density of MPS molecules on the MMT surface influences surface wettability. After MPS modification, a PS layer was coated onto the MMT surface via seeded emulsion polymerization. Upon adding K₂S₂O₈, styrene monomers began to polymerize and anchor onto the MMT nanosheet surfaces via copolymerization with the C=C bonds of MPS. Styrene monomers can also self-nucleate and form PS spheres. Both the PS layer on the MMT surface and the free PS spheres can be dissolved by toluene, following the “like dissolve like” principle [34]. After dissolution, the PCM nanosheets were obtained, as one end of the PS chain is covalently bonded to the C=C bonds on the MMT surface.

SEM was performed to examine the morphological evolution of MMT during the two surface modification steps. Pristine MMT (PM) has a layered structure, with lateral dimensions ranging from several hundred nanometers to several micrometers and a thickness of only a few nanometers. MMT layers tend to aggregate due to electrostatic attraction. As shown in Figure 2a and Figure S1a, bulk MMT exhibits significant aggregation, where the MMT nanosheets are in contact with each other. To obtain few-layered MMT nanosheets, the bulk MMT was treated with ultrasonication to exfoliate the nanosheets. The exfoliated MMT nanosheets were then soaked in an alkaline solution, introducing hydroxyl groups onto the edges and in-plane surfaces of the MMT. Subsequently, MPS was added to the MMT suspension, where the MPS molecules coupled with the hydroxyl groups on the MMT surface via silane coupling, forming MCM nanosheets. Figure 2b and Figure S1b show the SEM images of the MCM, where the exfoliated MMT nanosheets maintain their few-layered structure after MPS treatment. Figure 2c and Figure S1c present SEM images of MMT nanosheets after PS coating, showing the formation of a multi-layered MMT structure and the presence of some free PS spheres. After treatment with toluene, the PS layer and the free PS spheres dissolved, leaving only the few-layered MMT nanosheets, as shown in Figure 2d and Figure S1d.

FTIR was conducted to further investigate the surface modification process of MMT. As shown in Figure 3a, the absorption peak at 2960 cm⁻¹ corresponds to the stretching vibration of the CH₃ group, while the peak at 1419 cm⁻¹ is attributed to the symmetric stretching vibration of the C=O-O group. These peaks confirm that MPS has been successfully capped onto the MMT surface, albeit with weak peaks due to the small quantity of MPS molecules. After being capped with the PS layer (Figure 3b), new peaks appear at 2923 cm⁻¹, 2870 cm⁻¹ and 1450 cm⁻¹, corresponding to the stretching vibrations of CH₃ and CH₂. Peaks at 3061 cm⁻¹ and 3024 cm⁻¹ correspond to the stretching vibration of CH in the aromatic rings, and the peaks at 1604 cm⁻¹, 1495 cm⁻¹, 755 cm⁻¹ and 699 cm⁻¹ are assigned to the symmetric stretching vibration of the aromatic rings [31]. These results indicate that the PS layer has been successfully capped onto the surface of MMT nanosheets. Figure 3c shows the XPS survey spectrum of the PM, MCM and PCM nanosheets. After calibrating the C1s peak, Figure 3d reveals a peak at 284.6 eV, which corresponds to the C-C bond. Figure 3e shows the C1s spectrum of MCM nanosheets, where peaks at 286.6 eV and 289.2 eV are attributed to the C-O and C=O-O bonds, respectively. This indicates that MPS has been successfully grafted onto the MMT surface. Figure 3f shows the C1s spectrum of PCM, where the peaks at 284.4 eV and 291.1 eV can be assigned to the π-π * transitions in the aromatic rings and the π-π * satellite. This provides evidence for the presence of the PS layer. These XPS and FTIR results collectively demonstrate the successfully capping of MPS and PS layers on the MMT nanosheet surface. In addition to the morphological structure and chemical composition, the surface electrical properties are another key factor influencing the energy storage performance of PP nanocomposites [8]. Figure 3g–i show the zeta potential of pristine and modified MMT nanosheets. The PM nanosheets exhibit an average potential of −10.3 mV, indicating a negative surface charge of PM nanosheets. After capping with the MPS molecule, the average potential decreased to −5.86 mV due to the electric neutral nature of the MPS. After capping with the PS layer, the average potential further decreased to −1.7 mV, demonstrating successful surface modification with MPS and PS. The XRD patterns in Figure S2 show distinct diffraction peaks corresponding to MMT (PDF#46-1045) and a relatively weak diffraction peak of SiO₂ (PDF#13-0259), which originates from impurities in the PM sample. Additionally, the d₀₀₁ diffraction peak shifted to a small angle after surface modification, indicating an increase in the interlayer spacing [35]. This suggests that the MPS molecules are not only capped on the surface of MMT nanosheets but are also intercalated into the interlayer of the MMT nanosheets.

3.2. Dielectric Energy Storage Properties of PP Nanocomposites

After surface modification with the PS layer, the PCM nanosheets show improved compatibility with the PP matrix. The PP nanocomposites were prepared by solution blending. Initially, PP pellets were dissolved in paraxylene at high temperature. Then, the suspension of PCM nanosheets in toluene was added. After thorough mixing, the PP nanocomposite was obtained after freeze-drying and hot-pressing. As aforementioned, the E_b is a crucial parameter in determining the overall U_e of the dielectric materials. A two-parameter Weibull distribution function:

$$P=1-exp(-E_b/\alpha {)}^{\beta },$$

(5)

was employed to evaluate the characteristic E_b of the PP nanocomposites with PCM nanosheets. In Equation (5), P is the cumulative failure probability, α is the characteristic E_b that corresponds to a 63.2% probability of failure, and β is the slope parameter that characterizes the scatter of the result [36,37]. In this work, E_b was calculated using linear fitting of the Weibull distribution function for 15 specimens per sample. As shown in Figure 4a,b, the E_b value of pristine PP is 463 MV m⁻¹. Compared with pristine PP, the E_b values of PP nanocomposites with increasing PCM loading fractions initially increase and then decrease. At low filler concentrations, the nanosheets are well dispersed within the matrix. The 2-dimensional nanosheets can effectively suppress charge transfer through the matrix and delay space charge accumulation, thereby enhancing E_b. Additionally, the well-dispersed fillers can obstruct the expansion of electrical treeing paths, thus prolonging the time required for breakdown. However, at high filler concentrations, the nanosheets show poorer dispersion, and the interface regions begin to overlap, reducing their ability to trap free charge carriers. Once breakdown initiates within the dielectric, the electrical treeing path forms more easily through the overlapped interface regions. The maximum E_b value of 619 MV m⁻¹ is obtained from the PP nanocomposite with 0.2 wt% PCM nanosheets, which is 33.7% higher than that of pristine PP. To investigate the effect of PCM nanosheets on the E_b of PP nanocomposites, we employed the phase-field breakdown model to simulate the dielectric breakdown process. Notably, a conductive field concentrator was introduced to account for the effect of random defects, in line with standard fracture mechanics testing procedures [38,39]. Figure 4c,d display the breakdown paths of pristine PP and PP nanocomposites with PCM nanosheets. For pristine PP, breakdown initiates in regions with significant electric field concentrations, penetrating through the entire matrix and forming a breakdown path. In contrast, when the PCM nanosheets are oriented perpendicular to the electric field in the PP matrix, they dissipate energy and create a more tortuous path, thus increasing the E_b of the PP nanocomposite, as shown in Figure S3. This increase in E_b can be ascribed to the improved compatibility between PCM nanosheets and the PP matrix. The optimized dispersion of PCM nanosheets in the PP matrix serves as energy dissipation sites and provides a more tortuous path for charge carriers, enhancing E_b.

In addition to E_b, ε_r is another key factor influencing U_e. The ε_r of the PP nanocomposites was analyzed, as shown in Figure 5a. Both ε_r and tanδ exhibit stable frequency-dependent behavior in the range of 10² to10⁶ Hz. The ε_r of the PP nanocomposites increases gradually with the increasing PCM loading fraction. The maximum ε_r value of 2.76 at 10³ Hz is obtained for the PP nanocomposite with 0.5 wt% PCM nanosheets, which is 23.2% higher than the pristine PP (2.24 at 10³ Hz). For PP composites, ε_r is influenced not only by the inherent properties of the polymer matrix and PCM nanosheets, but also by the interface characteristic between the two phases. When an electric field is applied, charges accumulate at the interface between the polymer matrix and the fillers, resulting in local polarization effects, a phenomenon known as interfacial polarization. This effect is particularly pronounced at the interface, especially when the filler exhibits a large surface area. The increase in the ε_r of the PP nanocomposite containing PCM nanosheets can be attributed to the interfacial polarization induced by the large surface area of the nanosheet fillers. Notably, the ε_r of the polymer nanocomposites increases substantially even at very low filler loading (<0.3%). Moreover, while tanδ increases with filler loading, it remains at very low levels. For example, the tanδ of pristine PP is 0.0003 at 10³ Hz, while that of the PP nanocomposite with 0.5 wt% PCM is 0.0017 at 10³ Hz. The tanδ of the PP nanocomposite with 0.2 wt% PCM is 0.0006, which is extremely low and comparable to that of pristine PP. This suggests that the energy loss in the PP composite dielectric is minimal, favoring high efficiency and U_e.

Polarization, which is related to ε_r, is another critical parameter for determining the U_e of dielectric [40]. The unipolar displacement–electric field (D-E) loops under varying electric fields for PP and PP nanocomposites are shown in Figure 5b and Figure S4. All the loops exhibit linear characteristics. The maximum polarization (P_m) derived from these loops is shown in Figure 5c. Compared with pristine PP, the P_m of PP nanocomposites improves significantly with the incorporation of PCM nanosheets. Notably, P_m increases initially and then decreases as the PCM loading ratio increases. The P_m value of pristine PP is 0.91 μC cm⁻². With 0.2 wt% PCM nanosheets, the P_m value of the PP nanocomposite increases to 1.47 μC cm⁻². However, further increasing the PCM loading beyond 0.2 wt% leads to a decrease in P_m. The initial increase in P_m is attributed to enhanced interfacial polarization due to the large surface area of the PCM nanosheets. The subsequent decrease in P_m beyond 0.2 wt% PCM nanosheets can be ascribed to the reduced P_m of these composites. As the key parameters in evaluating the energy storage capability, the U_e and η of the PP nanocomposites were calculated from the D-E loops. As shown in Figure 5d, U_e increases with PCM loading at the same electric field due to the increase in polarization. However, the maximum U_e of the PP nanocomposites shows a tendency to increase first and then decrease with increasing PCM loading. For instance, the maximum U_e of 4.23 J cm⁻³ is achieved for the PP nanocomposite with 0.2 wt% PCM nanosheets, which is 225% that of pristine PP (U_e value of 1.88 J cm⁻³). In contrast, the U_e of the PP nanocomposite with 0.5 wt% PCM nanosheets is 1.93 J cm⁻³, much lower than that of the 0.2 wt% PCM nanosheets. This decreased U_e could be ascribed to the reduced E_b of the PP nanocomposite with 0.5 wt% PCM nanosheets, despite its enhanced ε_r. Although the η value of all samples decreases with an enhanced applied electric field, it remains above 90% until breakdown occurs.

4. Discussions

This work demonstrates that grafting a PS layer onto the surface of MMT significantly improves compatibility between the PP matrix and MMT nanosheets, thereby enhancing the energy storage capability of the PP nanocomposites.

Table 1. Comparison of dielectric properties of PP nanocomposites with 2-dimensional nanofillers.

Sample	εr (1000 Hz)	Eb (MV m−1)	Ue (J cm−3)	Reference
BN/BT	2.92	469	2.82	[41]
PP-g-MAH/BN/BT	4.68	324	2.45	[42]
PP-g-MAH/BN	2.27	437	4.11	[43]
MMT	2.38	370	1.70	[44]
PP-g-MAH/MMT	3.35	520	5.20	[45]
PP-g-MAH/MMT	3.75	530	5.21	[29]
PCM	2.47	619	4.23	This work

Table 1. Comparison of dielectric properties of PP nanocomposites with 2-dimensional nanofillers.

Sample	εr (1000 Hz)	Eb (MV m−1)	Ue (J cm−3)	Reference
BN/BT	2.92	469	2.82	[41]
PP-g-MAH/BN/BT	4.68	324	2.45	[42]
PP-g-MAH/BN	2.27	437	4.11	[43]
MMT	2.38	370	1.70	[44]
PP-g-MAH/MMT	3.35	520	5.20	[45]
PP-g-MAH/MMT	3.75	530	5.21	[29]
PCM	2.47	619	4.23	This work

Table 1 compares the dielectric properties and energy storage performance of the present study with other PP nanocomposites that incorporate various two-dimensional nanofillers. The PP nanocomposites featuring PCM nanosheets, as presented in this work, demonstrate a competitive U_e. Although the PP nanocomposites referenced in [29,45] show higher U_e due to the increased polarization associated with PP-g-MAH, the PP nanocomposites with PCM nanosheets exhibit a superior E_b, indicating higher electrical resistance.

This innovative polymerization–dissolution strategy holds promise for industrial applications in the manufacturing of PP nanocomposites. However, the use of toluene in this process presents a challenge for scaling up due to its environmental impact and the operational costs associated with recovery. While the results indicate an enhancement in dielectric constant, breakdown strength and energy storage density, several limitations should be addressed in future work. First, interfacial properties, including the grafting density of the PS layer, molecular architecture and the interfacial electronic structure, could significantly influence the energy storage performance of the PP nanocomposites. Second, variations in MMT loading may affect the crystallinity of the PP matrix, which is another critical factor impacting the energy storage capability of these nanocomposites. Therefore, further research should focus on the crystallinity of PP to comprehensively explore the effects of MMT loading from a deeper perspective. Third, the reliability of dielectric capacitors, particularly concerning thermal stability, cycling endurance and long-term performance, remains crucial for practical applications. A thorough investigation into the lifespan and failure mechanisms of the PP nanocomposites incorporating PCM nanosheets under elevated temperatures, along with continuous cycling performance tests, is essential for practical implementation.

5. Conclusions

In summary, we have developed a novel polymerization–dissolution strategy for modifying hydrophilic MMT nanosheets. Initially, the surface is capped with MPS molecules, followed by a PS capping. This capping with nonpolar molecules enhances the compatibility of the MMT nanosheets with the PP matrix. Through phase-field simulations, the underlying mechanism of charge transfer has been revealed. The well-dispersed PCM nanosheets serve as scattering centers, reducing the likelihood of secondary impact-ionized charge generation. Additionally, they effectively dissipate energy and hinder charge acceleration within the PP matrix, thus preventing the growth of electrical trees. As a result, the PP nanocomposite with 0.2 wt% PCM nanosheets exhibits a significantly enhanced E_b of 619 MV m⁻¹ and a discharged U_e of 4.23 J cm⁻³. This work provides an effective strategy to enhance the energy storage capability of PP-based nanocomposite dielectrics.