Synthesis of poly(methyl methacrylate–methallyl alcohol) via controllable partial hydrogenation of poly(methyl methacrylate) towards high pulse energy storage capacitor application

Abstract

Methallyl alcohol has never been reported to be homo-polymerized or copolymerized directly with other monomers. In this manuscript, we report the first synthesis of poly(methyl methacrylate–methallyl alcohol) (P(MMA–MAA)) copolymers via an indirect polymerization process involving the partial hydrogenation of PMMA. The copolymers with varied monomer molar ratios have been carefully characterized with nuclear magnetic resonance (NMR), differential scanning calorimetry (DSC), thermogravimetry analysis (TGA) and dynamic mechanical analysis (DMA). The introduction of –OH groups leads to the formation of H-bonds among –OH and ester groups, which is responsible for the enhanced glass transition temperature and Young's modulus. As a result, the permittivity of P(MMA–MAA) is increased at low MAA content and reduced quickly as more MAA introduced. The breakdown strength (E_b) of P(MMA–MAA)s is improved significantly from about 400 MV m⁻¹ of PMMA to over 550 MV m⁻¹ of P(MMA–MAA) bearing 19 mol% MAA units. The highest discharged energy density (U_e) is observed as 13 J cm⁻³ at 550 MV m⁻¹ electric field which is 2–3 times larger than BOPP and 50% higher than that of PMMA. Most interestingly, energy loss (U_l) is well maintained at about 8%@550 MV m⁻¹, which is rather close to biaxially oriented polypropylene (BOPP). The promising energy storage capability and excellent energy discharging efficiency of the P(MMA–MAA) copolymer could finally meet the desperate need in high pulse energy storage capacitors. Constructing strong H-bonds in glassy dipolar polymers might offer a great option for designing and fabricating polymeric dielectrics with high U_e and low U_l.

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Introduction

As static electric energy storing and releasing devices, metalized film capacitors possess many advantages including low cost, high breakdown strength, good processability, high resistivity and low dielectric loss, which allow them to be widely utilized in electric and electronic systems.^1–5 With the quick development and requirement for capacitors with high electric energy storage capability such as hybrid electric vehicles and electrically driven weapon systems, polymeric dielectrics with high energy density, high temperature capability as well as low dielectric loss are desperately needed.^6–8 As indicated in U_e = 1/2ε_rε₀E_b², where ε_r is the relative dielectric constant, ε₀ is the vacuum permittivity, and E_b is the electric field allowed to be applied,⁵ high U_e could be achieved either by enhancing ε_r of the polymeric films or improving their breakdown strength (E_b).

During past decades, a great deal of effort has been devoted to discover and fabricate neat polymers and polymer based composites with high permittivity. Among all the polymeric materials, poly(vinylidene fluoride) (PVDF) based ferroelectrics have attracted most attractions for their high ε_r.^9–18 Several classes of novel PVDF based ferroelectric polymers with varied chemical and aggregation structures and tunable ferroelectric behaviors have been developed chemically or physically to achieve high U_e. The highest U_e of 25 J cm⁻³ under an electric field of 700 MV m⁻¹ has been reported in PVDF based fluoropolymers, which is 6–8 times of current state-of-the-art polymer dielectric material (ca. biaxially oriented polypropylene (BOPP)).¹⁸ However, due to the ferroelectric relaxation nature and relatively low electrical resistivity, the charge and discharge curves on displacement–electric field (D–E) loops are poorly coincided. As a result, relatively large hysteresis together with high energy loss (U_l) (about 30–70% of the charged energy) has been observed. That would lead to the quick temperature rising and early failure of the capacitors during working under elevated electric field. By introducing high glass transition temperature (T_g) side polymer chains (ca. polystyrene and poly(methacrylate esters)), U_l could be effectively reduced to about 20% under 650 MV m⁻¹ in poly(ethyl methacrylate) (PEMA) grafted PVDF based copolymers, which is still 2–3 times of BOPP under the consistent field. Further reducing U_l of the PVDF based ferroelectric polymers becomes an impossible mission for their high polar nature.^19–22

Modifying the state-of-the-art polymer dielectric materials such as BOPP to improve their permittivity seems to be another option to achieve high U_e in these low polar materials since very high E_b (ca. 730 MV m⁻¹) has already been reached in BOPP and further improving E_b in a large scale is barely feasible in industry.⁷ Apparently, U_e of BOPP based low polar dielectric materials could hardly be further improved with respect to their limited electronic and atomic polarization^23,24 as well as the ultimate E_b values under practical processing conditions. However, chemically introducing small amount of flexible –OH and –NH₂ polar groups (less than 4.2 mol%) into PP has been finely reported in Chung's group to increase ε_r of pristine PP from 2.2 to PP–OH of 4.6. U_e is improved from BOPP of 4.5 to PP–OH of 7 J cm⁻³ at 600 MV m⁻¹ as well.²⁵ Whereas, typical slim D–E loops of pristine BOPP start to widen and lose the linear dielectric behavior as more hydroxyl groups introduced, which means U_l is increasing as well. Meanwhile, introducing more –OH groups would hinder the crystallization of PP segments and reduce the Young's modulus of the films, which is not favorable for high E_b as indicated in E_b = 1.6(Y/ε₀ε_r)^1/2, where Y is referring to the Young's modulus.^26–28 That means functionalizing PP films by introducing polar groups into them could only enhance their U_e in a limited scale with respect to the increased U_l as well. Interestingly, the authors found that a unique network structure via inter-chain –OH groups dimerization (H-bonding) might be formed among –OH groups, offering not only higher polarity but also good reversibility over a wide range of temperatures (−20–100 °C) and frequencies (100 Hz to 1 MHz).

Considering the limited U_e of BOPP and high U_l of PVDF based ferroelectric polymers bearing both crystalline and amorphous phases, the dipolar glassy polymers have been put forward as the third types of polymer candidates to realize high U_e, low U_l and high operative temperature more recently. Polythioureas (PUs) with tunable ε_r from 4.2–5.6 have been reported by Zhang's group to be capable of maintaining its low U_l up to very high operation fields (>600 MV m⁻¹) with a very high U_e (13–20 J cm⁻³).^29–31 These glassy polymers possess high T_g and weak coupling forces among dipoles, which allow them to be operated at elevated temperature without polar hysteresis loss. It seems the freestanding thin films could hardly been fabricated in these polymers for their brittle nature, which might limit their application as foldable dielectric films. More recently, Zhu's group presented another glassy polymer dielectric poly(2-(methylsulfonyl)-ethyl methacrylate) (PMSEMA) with a high dielectric constant of 11.4 and a low dissipation factor (tan δ) of 0.02 at 25 °C and 1 Hz, which is nearly 3 times that of its analogue polymer poly(methyl methacrylate) (PMMA).³² According to the previous work, the promising glassy polymer should have a high T_g and a significant dipolar polarization under high electric field. High T_g is essential for high Young's modulus to prevent U_l from electronic and ionic conductions, and isolated dipole rotation under high field is responsible for the elevated ε_r, thus high energy storage capability. Apparently, syndiotaxy PMMA (s-PMMA) should be a very promising glassy polymer dielectric for its relatively high ε_r (ε_r = 4.3 at 100 Hz and 25 °C), high T_g of about 115 °C and massive productivity with low cost. However, its D–E loops sunder elevated electric field are rather broad and U_l is relatively high, which could be attributed to its ester dipole relaxation around room temperature or β-relaxation.³³ Therefore, minimizing the U_l of PMMA by confining its β-relaxation is expected to be an effective way to realize high U_e and low U_l in PMMA and make it acceptable in high pulse energy storage capacitors.

Inspired by the idea that constructing H-bonds in functionalized PP such as (PP–OH and PP–NH₂), in this paper, a family of poly(methyl methacrylate–methallyl alcohol)s (P(MMA–MAA)s) with tunable MMA/MAA molar ratios has firstly been synthesized by controlled hydrogenating s-PMMA with LiAlH₄ under mild conditions to convert desirable content of MMA into methallyl alcohol (MAA) units. The inter-chain H-bonds constructed among –OH and ester groups in P(MMA–MAA)s are designed to confine β-relaxation of PMMA, thus reduce the U_l under elevated electric field and improve E_b. The structure and the dielectric performance of copolymers with varied composition have been finely characterized. As expected, the dielectric properties of resultant copolymers could be facilely tuned by varying the molar content of MAA units. U_e of P(MMA–MAA)s (bearing 19 mol% MAA) was significantly improved (ca. 2–3 times larger than that of BOPP), and U_l was maintained at as low as 8%@550 MV m⁻¹, which is rather close to BOPP under the consistent electric field.

Experimental

Materials

Polymethyl methacrylate (M_w = 550 000), lithium aluminium hydride (LiAlH₄) were purchased from Alfa-Aesar. Tetrahydrofuran (THF, Tianjin Reagents Co. Ltd, AR grade) was dried and distilled from sodium/benzophenone. The other solvents are purchased from Tianjin Reagents Co. Ltd and used as received.

Synthesis of P(MMA–MAA) copolymer

P(MMA–MAA) copolymer was synthesized via a hydrogenation process of PMMA using LiAlH₄ as reductive agent as described in Scheme 1. 2 g PMMA was dissolved in 50 mL dehydrated THF in a two-necked round-bottomed flask equipped with a magnetic stirrer. The required amount of LiAlH₄ was introduced into the solution and the hydrogenation reaction was carried out at 25 °C for 24 h. The molar ratio of MMA and MAA were varied by altering the amount of LiAlH₄ introduced. Acetone was injected to the solution to terminate the reaction before it was precipitated in a hydrochloride acid solution with a concentration of 5 vol%. The precipitant was dissolved in acetone and precipitated into deionized water repeatedly for 3 times. The obtained P(MMA–MAA) copolymer was soaked in deionized water for 3 days (the water was changed twice everyday) in order to completely remove residue impurities before completely dried under reduced pressure at 50 °C for 3 days.

Scheme 1 Synthesis of P(MMA–MAA) copolymers from hydrogenation of PMMA.

Films fabrication and processing

All the polymer films with a thickness about 20 μm were prepared following a solution-cast process by casting the copolymer solution (about 3 wt% in DMF) onto glass plates. After the solvent was evaporated completely at 70 °C, the films together with the plates were annealed at 150 °C for 4 h before peeled off from glass plates for performance characterization.

Instruments and characterization

¹H NMR spectra was recorded on a Bruker spectrometer (400 MHz, Advance III) in dimethyl sulfoxide-d6 (DMSO-d6) with tetramethylsilane (TMS) as internal standard. Differential scanning calorimetric (DSC) analysis was conducted on a Netzsch DSC 200 PC (Netzsch, Germany) in a nitrogen atmosphere at a scanning rate of 10 °C min⁻¹ after a circle of quick heating (20 °C min⁻¹) and cooling (20 °C min⁻¹) to remove the thermal history. Dynamic thermomechanical analysis (DMA) was carried out by TA DMA Q800 (TA, USA) was performed in tension mode using an oscillating force ramp. All samples were prepared by the solution casting on the Teflon plate, the sample size was 26 mm (L) × 0.2 mm (T) × 5 mm (W). The initiated force was 0.01 N at a frequency of 10 Hz and a heating rate of 10 °C min⁻¹. Thermogravimetric analysis (TGA) was performed on a thermal analyzer (Mettler Toledo, TGA 1) at a heating rate of 10 °C min⁻¹ under inert gases. Fourier transform infrared (FTIR) spectroscopy of the films was recorded on a Tensor 27 (Bruker, Germany) with a resolution of 1–0.4 cm⁻¹ in transmission mode. Scanning electron microscope (SEM) images were recorded on a field emission scanning electron microscope (Hitachi FESEM SU6600) at an accelerating voltage of 4.0 kV. For electrical performance characterization, gold electrodes with a thickness about 80 nm were sputtered on both surfaces of the polymer films with a JEOL JFC-1600 auto fine coater (Japan). The dielectric properties were measured in a frequency range from 100 Hz to 10 MHz at 1 V on an Agilent (4294A) precision impedance analyzer. The electric D–E hysteresis loops at room temperature were obtained on a Premiere II ferroelectric tester from Radiant Technologies, Inc., where AC electric fields ranging from 25–550 MV m⁻¹ were applied across polymer films with a triangular waveform at a frequency of 10 Hz. The breakdown electric field for all the sample films was determined by the breakdown voltage tester (Beijing Beiguangjingyi Instrument Equipment Co., LTD) and the electric field applied rate of 1 kV min⁻¹. The diameter of the cylindrical electrode was 10 mm. The leakage current was obtained as less than 10 mA until the film was electrically broken down.

Results and discussion

Synthesis of P(MMA–MAA) copolymer

Scheme 1 illustrates the synthesis of P(MMA–MAA) from PMMA using conventional hydrogenation reaction via converting MMA into MAA units by LiAlH₄ in THF at 25 °C. The structure of the purified P(MMA–MAA) was confirmed by ¹H NMR and FTIR (Fig. 1, S1 and S2†). The assignment of ¹H NMR is as follows, the peaks at 0.3–0.8 ppm (m, 3H, –CH₂–CH(COOCH₃)–CH₃) (I₁), 0.8–1.1 ppm (m, 2H, –CH₂–CH(COOCH₃)–CH₃–) (I₂) and 4.0–4.2 ppm (m, 3H, –CH₂–CH(COOCH₃)–CH₃–) (I₃) are consistent with the pristine PMMA. The new signals emerging at 2.7–3.2 ppm (m, 2H) and 4.0–4.5 ppm (m, 1H) are assigned to the protons on –CH₂–CH(CH₃)–CH₂OH– (I₄) and –CH₂–CH(CH₃)–CH₂OH– (I₅) of MAA, respectively. The peaks at 2.54 ppm and 3.33 ppm could be assigned to the trace amount of non-deuterized solvent and water. The chemical compositions of the as-formed copolymer could be calculated from ¹H NMR results using the following equation,

Fig. 1 (a) ¹H NMR spectra of pristine PMMA and as-formed P(MMA–MAA) copolymer with a molar ratio of MMA/MAA = 81/19. DMSO-d6 was used as solvent with TMS as internal standards. (b) ¹³C NMR spectra of pristine PMMA and as-formed P(MMA–MAA) copolymer with a molar ratio of MMA/MAA = 81/19. Chloroform-d was used as solvent.

Herein, P(MMA–MAA)s with varied MAA contents ranging from 3 to 19 mol% were successfully obtained by facilely altering the addition of LiAlH₄, and the results were listed in Table 1.

Table 1 P(MMA–MAA)s with varied composition synthesized under different conditions

Entry^a	LiAlH4 [g]	MMA/MAA in molar ratio^b	Tg^c [°C]	Storage modulus^d [MPa]	Loss modulus^d [MPa]
a 2 g PMMA and 50 mL dried THF were utilized in all the reactions and the reactions were conducted at 25 °C.b The molar ratio of MMA/MAA was determined by ^1H NMR.c Determined by DSC.d Determined by DMA at 25 °C.
Pristine PMMA	0.025	100/0	116	1464	111
P(MMA–MAA)-1	0.075	97/3	117	1501	112
P(MMA–MAA)-2	0.1	92/8	119	1539	121
P(MMA–MAA)-3	0.15	86/14	121	1749	131
P(MMA–MAA)-4	0.2	81/19	128	2154	127

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The structure of as-formed P(MMA–MAA) copolymer was further confirmed by ¹³C NMR. Fig. 1b shows the completely assigned ¹³C NMR spectrum of pristine PMMA and as-formed P(MMA–MAA) copolymer with 19 mol% MAA unit. The signals centered at 178.0 (a), 54.0 (b), 51.1 (c), 41.0 (d) and 20.0 (e) ppm have been assigned to carbonyl, methylene, methoxy, quaternary carbon and α-methyl of PMMA, respectively. The new signal peaks situated at 60.1 (f) and 30.0 (g) can be ascribed to methylene neighboring to hydroxyl groups and quaternary carbon of MAA unit. The new peaks around peak (e) could be ascribed to the methyl of MAA unit.

Thermal and rheology properties

As an amorphous polymer, pristine PMMA shows a T_g around 116 °C at a heating rate of 10 °C min⁻¹ (see Fig. 2a) determined by DSC due to the strong dipolar interactions among ester side groups. After MMA units were converted into MAA, T_g of P(MMA–MAA) copolymers were continuously elevated from 116 °C of PMMA bearing no MAA units to 128 °C of P(MMA–MAA) containing 19 mol% MAA units as shown in Fig. 2a and listed in Table 1. That might be ascribed to the weak H-bonds formed between –OH on MAA and ester groups on MMA and strong H-bonds between two –OH groups when MAA content is sufficient high, which would be discussed in detail later on. As shown in Fig. 2b, the dynamic mechanical properties of the pristine PMMA and as-formed P(MMA–MAA) show strong dependence on the chemical composition thus the interaction among chains induced. For the glassy polymers, the polymer with strong interaction was usually stiffer at room temperature and more rigid than that with weak interaction. From DMA analysis, the storage modulus E′ of as-formed copolymer with lower MAA unit content (0–8 mol%) is slightly enhanced (less than 100 MPa) comparing with pristine PMMA. However, when MAA content is elevated to 14 and 19 mol%, E′ of P(MMA–MAA)s is improved to 1749 and 2154 MPa at 10 Hz frequency, which is 20% and 50% larger than that of PMMA, respectively. Although the dipole moment of –OH groups (about 1.68 D) is even slightly smaller than 1.75 D of ester groups, it seems the introduction of –OH groups improved the molecular interaction in P(MMA–MAA). Again, that might be attributed two types of H-bonding formed in the copolymers among –OH and ester groups. When MAA content is low, –OH groups would be randomly scattered in the copolymers and H-bonds are expected to be formed mostly between –OH and ester groups as illustrated in Scheme 2. The intermolecular force is slightly larger than that between ester groups in pristine PMMA, which may address for both the increased (E′) and T_g in small scale. As MAA content increases, the density of –OH groups would be quickly increased and the distance between two adjacent –OH groups are sharply reduced as well. When adjacent –OH groups are sufficient close, H-bonds among them are very likely to be generated. As a result, a three-dimensional network connected with two types H-bonds would be constructed in P(MMA–MAA). That is responsible for not only the significantly improved T_g and storage modulus but also the toughness in the P(MMA–MAA) bearing higher content of MAA units as indicated in Scheme 2. Most attractively, different from the high brittleness of PMMA, the three-dimensional networks constructed from the H-bonds with high mobility and flexibility is responsible for the improved TGA was employed to examine the thermal stability of the modified copolymers, as shown in Fig. 2c, where TGA thermograms of commercially available PMMA and P(MMA–MAA) containing 3 mol%, 8 mol%, 14 mol%, and 19 mol% of MAA unit were presented. Pristine PMMA possess two decomposition platforms at 250 °C and 300 °C, respectively. About 5% weight loss detected at 180–200 °C might be caused by the unpolymerized MMA monomer and trace content of moisture. That could be confirmed by the fact that no weight loss was detected in P(MMA–MAA) below 300 °C, which might be attributed to the removal of small molecular impurities including MMA monomer and moisture from the resultant copolymers during the repeatedly resolving and precipitation processing. It has been well investigated that the thermal decomposition of PMMA under inert gas atmosphere is mostly in the de-polymerization way, namely releasing MMA units.³⁴ Interestingly, the introduction of lower content of MAA units leads to a significantly elevated initial decomposition temperature (ca. 300 °C) and quick weight loss starting from about 350 °C, which are about 50 °C higher than that of pristine PMMA. That might be attributed to the different decomposition method of MAA from MMA units and the weak H-bonds constructed between –OH groups on MAA and ester groups on MMA. Meanwhile, chemical crosslinking is likely to take place based on the interesterification³⁵ between the –OH groups on MAA units and –COOCH₃ groups on MMA at elevated temperature. Apparently, these effects may improve the intermolecular interaction, lower down the sensitivity of MMA units to heating and increase the degradation temperature of P(MMA–MAA). Further increasing the MAA content would result into the formation of H-bonding between two adjacent –OH groups as discussed above. Different from copolymers with lower MAA content, H₂O produced from the etherification of two –OH groups at high temperature would take place,³⁶ which could be identified by the initial decomposition temperature at 250 °C on the TGA curves of P(MMA–MAA)-3 and P(MMA–MAA)-4 as illustrated in Scheme 3. The ether structure formed is also responsible for the further improved quick weight loss temperature to 370 and 380 °C in P(MMA–MAA)-3 and P(MMA–MAA)-4, respectively. In general, the introduction of –OH groups may help to construct flexible crosslinking by H-bonds in the resultant P(MMA–MAA) copolymers, which is responsible for the improved thermal stability, Young's modulus, and ductility (Fig. 3).

Fig. 2 (a) DSC thermograms of pristine PMMA and as-formed P(MMA–MAA) copolymers with varied MAA contents during the second heating circle (heating rate: 10 °C min⁻¹ under nitrogen gas). (b) Storage modulus and loss modulus measured from DMA of pristine PMMA and as-formed P(MMA–MAA) copolymers (initiated force: 0.05 N at frequency 10 Hz; heating rate: 10 °C min⁻¹ under argon gas). (c) TGA thermograms of pristine PMMA and as-formed P(MMA–MAA) copolymers with varied MAA contents during the heating process (heating rate: 10 °C min⁻¹ under argon gas).

Scheme 2 Formation of weak and strong H-bonds networks in as-formed P(MMA–MAA) with various MAA unit content.

Scheme 3 Proposed interesterification and etherification mechanisms of P(MMA–MAA) copolymer at elevated temperature.

Fig. 3 Photographs of (a) a bent 15 μm-thick P(MMA–MAA)-4 film and (b) 15 μm-thick P(MMA–MAA)-4 film wrapped around a glass tube with diameter 5 mm.

Dielectric properties

To investigate the dielectric properties of the copolymers, all the samples were prepared in a typical solution-casting procedure at 70 °C followed by a subsequent annealing treatment at 150 °C aiming to remove the solvent completely and improve the quality of films. Measured at 1 V bias and at elevated temperature of 25 °C, 65 °C, 90 °C, and 115 °C respectively, ε_r and dielectric loss (tan δ) of all the samples as a function of frequency were presented in Fig. 4. As testing frequency increases, ε_r of all the copolymers decreases slowly in the frequency range from 100 Hz to 1 MHz at 25 °C (Fig. 4a, c, e and g). Intriguingly, the permittivity at 100 Hz is significantly improved from 4.4 of PMMA to 6.0 of P(MMA–MAA)-1 bearing 3 mol% MAA. Further increasing MAA units from 3 mol% to 19 mol% leads to quickly dropped permittivity to 4.0 in P(MMA–MAA)-4. The composition dependence relationship of permittivity measured at elevated temperature is rather close to that detected at 25 °C, namely P(MMA–MAA)-1 > P(MMA–MAA)-2 > PMMA > P(MMA–MAA)-3 > P(MMA–MAA)-4. Although it has been well reported that the polarity of hydroxyl group (dipole moment of 1.68 D) is slightly smaller than that of ester groups (dipole moment of 1.75 D), in glassy state, the contribution of ester groups to the permittivity is rather small since all the high polar side groups are frozen in PMMA. However, the steric bulk of –CH₂OH is smaller than that of –COOCH₃, which means the swing of –CH₂OH requires less free space than that of –COOCH₃ groups driven by external electric field. Therefore, the higher mobility of –CH₂OH should be responsible for the elevated ε_r in P(MMA–MAA) bearing low content of MAA. As discussed above, as –OH group content increases, stronger H-bonds between two –OH groups would form, which has been well investigated in PP–OH polymers by Chung's group.^25,37,38 They demonstrated that these H-bonds are apt to form tight H-bonds clusters, which may prevent the –OH groups from relaxing and contributing to the dielectric loss, except possibly at time scales longer than the experimental measurement. Meanwhile, the formation of H-bonds and three dimensional network as well as increased modulus at higher content of hydroxyl group, may effectively confine the relaxation motion and switching of the dipole, which is responsible for the decreased ε_r in the copolymer containing higher MAA units. As the testing temperature increases, ε_r of all the samples is enhanced, which can be mostly ascribed to the increase of the free volume. That allows the dipoles to orientate more easily and higher ε_r is consequently observed. As testing temperature increases, the decreasing speed of permittivity against frequency is improved as well. That indicates the polarity and relaxation of the high polar groups requires varied time at different temperature, which would be more clearly illustrated by the dielectric loss later on.

Fig. 4 Dielectric constant ((a), (c), (e), and (g)) and dielectric loss ((b), (d), (f), and (h)) of pristine PMMA and as-formed P(MMA–MAA) copolymers with varied MAA contents, measured at 1 V bias and 25 °C, 65 °C, 90 °C, and 115 °C, respectively.

Dielectric loss of pristine PMMA and as-formed P(MMA–MAA)s measured at elevated temperature were presented in Fig. 4(b, d, f and h) respectively. At 25 °C, tan δ of copolymers is rather close to pristine PMMA since most of the dipole moments exhibit rather low mobility in glassy state. At 1 kHz, it could be identified that tan δ is in the same order of dielectric constant, namely, P(MMA–MAA)-1 > P(MMA–MAA)-2 > PMMA > P(MMA–MAA)-3 > P(MMA–MAA)-4. Meanwhile, tan δ of all the polymers show rather low dependence on the testing frequency since no relaxation would be tested in the range of 100 Hz to 1 MHz. As testing temperature increases, tan δ of all the samples at the consistent frequency is enhanced. As discussed above, elevated temperature favors the formation of larger free volume and thus the orientation of polar groups. As a result, the enhanced relaxation of the oriented polar groups is responsible for the enlarged tan δ. Meanwhile, as testing temperature increases, a broad peak appears in the testing range from 100 Hz to 100 kHz. The maximum tan δ is observed at about 1 kHz measured at 65 °C, at 3 kHz measured at 90 °C, and at 10 kHz on the curves detected at 115 °C, namely, shifting to the higher frequency as testing temperature increases. It is accepted that β relaxation observed in polymethacrylates (PMA) and PMMA arises from the rotational motion of the side chain about the C–C bond which links to the main chain.³⁹ β transition for PMA is reported to be between 10 and 60 °C, regardless of the length of the alkyl side chains (e.g., methyl, ethyl, propyl, and butyl).^33,40 Therefore, the broad peaks observed on tan δ curves could be assigned as the β relaxation considering the gradually increased ε_r and shifting to higher frequency as testing temperature increases.⁴¹ The characteristic relaxation time of the β process shows an Arrhenius-like temperature dependence. That may address for the moving towards high frequency zone with the elevation of testing temperature. As discussed above, the introduction of lower content of –OH groups favors the ε_r for the higher mobility of –OH groups with smaller steric bulk, which is responsible for the higher tan δ peak value obtained as well. The crosslinking networks formed among higher content of –OH groups is responsible for the reduced tan δ peak value as well. Meanwhile, tested at 25 °C and 65 °C, tan δ peaks of the copolymers exhibits slight frequency dependence, namely, the maximum tan δ of copolymer bearing higher content of MAA is observed at lower frequency. That strongly confirms the confinement effect of H-bond networks onto the orientation of dipole moments in P(MMA–MAA).

D–E hysteresis behavior under high electric fields

The D–E hysteresis behaviors of the pristine PMMA and as-formed P(MMA–MAA) films were studied using D–E hysteresis loop measurements under AC and DC electric fields at a frequency of 10 Hz as shown in Fig. 5. The applied fields were increased with a step of 25 MV m⁻¹ until the films were electrically broken down (Fig. 5b and c). As expected, PMMA film displays linear D–E loops with relatively low hysteresis slope, indicating a consistent permittivity and completely reversible polarization–depolarization over the whole applied electric fields. At ambient temperature, the high polar ester groups are completely frozen and the flipping of the side groups is frozen by the strong intermolecular force. No relaxation would be detected at 10 Hz even under elevated electric field. As a result, a rather low slope of the curve and slim D–E loops are observed. All the P(MMA–MAA) copolymers exhibit similar linear and slim D–E loops as well as shown in Fig. 5a. Interestingly, the slope of the D–E loop increases with the hydroxyl group content from 0 to 3 mol% and decreases with hydroxyl group content from 3 to 19 mol%. The trend is consistent with the dielectric performance under low electric field as shown in Fig. 4, namely, the dielectric constant (slope) increases with the hydroxyl group content from 0 to 3 mol% and decreases with the hydroxyl group content from 3 to 19 mol% in P(MMA–MAA) copolymers. As an instance, the displacement of P(MMA–MAA)-1 (3 mol%) reaches 3.57 μC cm⁻² at 250 MV m⁻¹, which is higher than that of PMMA (3.16 μC cm⁻²) under the same applied electric field. Meanwhile, a slightly broader D–E loops than PMMA are observed in P(MMA–MAA)-1 for the relaxation of oriented dipole moment from –OH groups. Further increasing MAA molar content leads to continuously reduced displacement of 3.26 μC cm⁻² in P(MMA–MAA)-2, 2.75 μC cm⁻² in P(MMA–MAA)-3 and 2.44 μC cm⁻² in P(MMA–MAA)-4. Slimmer D–E loops than PMMA are detected as well. That could be finely illustrated by the weak and strong H-bonds networks constructed among –OH and ester groups as well. At low MAA content, the higher mobility of –OH groups and weak H-bonds between –OH and ester groups are responsible for the elevated displacement under consistent applied field. The formation of strong H-bonds among –OH groups and its confinement onto the mobility of ester groups are the major reason for the depressed slope of D–E curves.

Fig. 5 (a) Comparison of D–E loops of pristine PMMA and a series of P(MMA–MAA) with different MAA mol% content at an electric field of 250 MV m⁻¹. D–E loops of pristine PMMA (b) and P(MMA–MAA)-4 (c) at elevated electric fields.

Therefore, the dielectric performance dependence onto temperature, frequency and electric field results clearly indicates the unique contribution of –OH polar groups in P(MMA–MAA) dielectric thin film. At lower hydroxyl group content before the strong H-bond clusters formed, the smaller –OH side groups will contribute directly to the enhanced dielectric constant and charge displacement for their better mobility than ester groups. As soon as the strong H-bonds state to be formed with the increasing of hydroxyl group content, the physical crosslinking network structure formed may confine the orientation of the dipole moment both on side groups and chain segment.

Breakdown strength

According to the formula U_e = 1/2ε_rε₀E_b², the electric field allowed to be applied E_b plays a more important role than ε_r in achieving high U_e in linear dielectrics. To fully evaluate and understand the electric performance, E_b of PMMA and P(MMA–MAA) copolymer samples were examined for over 20 times, and the experimental results were analyzed by the two-parameter Weibull analysis method illustrated by the following equation,⁴²wherein x is the measured E_b, α is the characteristic breakdown field, at which at least 63.2% of the samples are observed to fail, and β is a shape parameter. Fig. 6 and S3† present E_b together with the estimated α and β values of pristine PMMA and as-formed P(MMA–MAA) copolymers.

Fig. 6 Weibull breakdown strength and shape parameter as functions of MAA mol% content in pristine PMMA and as-formed P(MMA–MAA) copolymers.

As shown in Fig. 6, pristine PMMA possesses a mediate breakdown strength (α value) of 436 MV m⁻¹. P(MMA–MAA)-1 with 3 mol% MAA exhibits a rather close α value (about 420 MV m⁻¹) to PMMA. As more MAA introduced, α value is significantly enhanced to over 520 MV m⁻¹ and reaches the maximum of 561 MV m⁻¹ in P(MMA–MAA)-4 with 19 mol% MAA units, which is about 30% larger than that of PMMA and rather close to the BOPP dielectric films. Meanwhile, a remarkable increase of β value from 7 of pristine PMMA to 16 in the as-formed P(MMA–MAA)-4 with 19 mol% MAA unit is found, suggesting vastly improved dielectric reliability of the as-formed P(MMA–MAA). As for electromechanical breakdown is well accepted breakdown method of polymers at room temperature, and many polymers have been found to follow the Stark Garton model,²⁶ whose E_b can be expressed as following equation,

where Y is Young's modulus of dielectrics, ε₀ is the permittivity of vacuum, and ε_r is the relative dielectric constant of dielectrics. Therefore, either lowering ε_r or improving Young's modulus favors the enhancement of E_b. Especially, Young's modulus of the materials is crucial for high E_b because it dominates the electromechanical failure caused by mutual Coulomb force from the opposite electrodes under high electric field. Apparently, the higher the modulus is, the larger Coulomb force the material can withstand.⁴³ On the other hand, the higher ε_r would lead to more static charges stored on the surface of the dielectric materials under the consistent electric field, which would generate larger Coulomb force and reduce the electric field withstanding ability of the films. From the measured Young's modulus and permittivity discussed above, the theoretical E_b of the samples based on the electromechanical was estimated.⁴⁴ As summarized in Fig. S4,† the theoretical E_b shows a rather similar trend as the experimental results presented in Fig. 6. That could be mostly attributed to the effects of MAA onto ε_r and Young's modulus as well discussed above. It is thus reasonable to conclude that H-bonds network within P(MMA–MAA) creates a robust scaffold to hamper the onset of electromechanical failure. Analysis within the framework of the electromechanical model highlights the significance of the excellent mechanical strength of H-bonds for improving the E_b of as-formed copolymers.

Note that the experimental data of Weibull E_b are much lower than the theoretically predicted values from the electro-mechanical model. This deviation is apparently due to the limitation in the model. Many polymer materials do not exactly follow the elastic stress–strain relationship interpreted by the above equation, and would experience plastic deformation and mechanical failure at an electric field far below E_b.²² Furthermore, dielectric breakdown of materials is usually governed by multiple factors, although the electromechanical effect⁴⁵ has been regarded as the main breakdown mechanism for many polymers. Dielectric materials studied by using the electromechanical model are in principle considered ideal insulators with the assumption of no electrical resistivity under the applied fields. Electrical resistivity,⁴⁵ however, cannot be ignored in a practical dielectric material, especially at high fields, and may dramatically increase the breakdown strength of materials. As shown in Fig. S5,† high-field (50 MV m⁻¹) DC resistivity measurement reveals that the as-formed P(MMA–MAA) with high MAA content ratios, i.e. >8 mol%, shows an order of magnitude improvement in electrical resistivity over that of the pristine PMMA polymer. These results indicate that the insulating network of H-bonds embedded in the polymer functions as an efficient barrier against the leakage current as well as the reduction of free volume space which reduce space-charge conduction.

Fundamentally, the enhanced E_b may be originated from the control of thin film morphology (including chain and/or side group motion) under high applied electric fields. SEM images (Fig. 7) of cross sections for the pristine PMMA and as-formed P(MMA–MAA)-4 copolymer films with 19 mol% MAA content were recorded to understand whether micro-phase separation occurred in the as-formed copolymer. According to the SEM images of section of P(MMA–MAA)-4 (19 mol%) copolymers, the two types of units (MMA and MAA) show fairly good miscibility with each other and no visible phase separation is observed. Meanwhile, the cross section of P(MMA–MAA)-4 is very dense thanks to the strong H-bonds constructed between –OH groups.^37,38 All these effects are crucial to achieve high E_b in dielectric films.

Fig. 7 Cross-section SEM images of the (a) and (b) pristine PMMA polymer film and (c) and (d) P(MMA–MAA)-4 copolymer film.

Energy storage and loss

The discharged energy densities (U_e) and the energy releasing efficiency of neat PMMA and as-formed P(MMA–MAA)s were calculated from the integral of the unipolar D–E loops as described in ref. 15. As shown in Fig. 8a, U_e of all the polymers obtained under the consistent electric field is in the same order as that of dielectric constant, namely, P(MMA–MAA)-1 > P(MMA–MAA)-2 > PMMA > P(MMA–MAA)-3 > P(MMA–MAA)-4. That could be well addressed by the formula U_e = 1/2ε_rε₀E_b², since U_e is in proportion to ε_r under the same electric field for all the linear dielectrics. However, the maximum U_e is increasing as more MAA unit is incorporated because E_b is significantly improved. For instance, the maximum U_e of the P(MMA–MAA)-4 (19 mol%) film is obtained to be 13.0 J cm⁻³ at E_b of 550 MV m⁻¹, which is about 80% larger than 7.3 J cm⁻³ of the pristine PMMA film at E_b of 350 MV m⁻¹. It is about 3 times of the currently widely utilized dielectric polymers BOPP under the consistent electric field for its higher permittivity.

Fig. 8 Comparison of (a) discharged energy density and (b) charge–discharge efficiency of pristine PMMA and as-formed P(MMA–MAA) copolymers with varied MAA content at different electric fields.

Besides U_e, the charge–discharge efficiency is another important metric of dielectric materials for energy storage capacitors since the unreleased energy generates heat. Consequently, unreleased energy densities (U_l) is detrimental to the performance and reliability of capacitors. As suggested in Fig. 8b, all the PMMA and P(MMA–MAA) films exhibit a linearly reduced releasing efficiency against elevated electric field like most of the glassy high polar polymer dielectrics. On the contrary, U_l is increased from about 4% at 50 MV m⁻¹ to about 8% at 350 MV m⁻¹ in PMMA. Under the consistent electric field, U_l exhibits the similar trend as that of U_e, P(MMA–MAA)-1 > P(MMA–MAA)-2 > PMMA > P(MMA–MAA)-3 > P(MMA–MAA)-4. For the linear dielectrics, the higher permittivity and electric energy achieved are always accompanied by the elevated dielectric loss from the oriented dipole relaxation and current leakage. Apparently, the introduction of MAA units at higher content could effectively reduce the energy loss due to the strong net-works constructed by H-bonds. Interestingly, U_l of P(MMA–MAA)-4 is obtained to be about 8% at 550 MV m⁻¹, which is slightly larger than 6.5% of BOPP@550 MV m⁻¹ and rather close to 7.6% of BOPP at its average breakdown strength at 600 MV m⁻¹. Considering to its about 3 times U_e and rather close U_l comparing with BOPP, P(MMA–MAA) copolymers with proper chemical composition and well-tuned flexible crosslinking networks might be promising polymeric dielectrics for high pulse energy storage capacitors.

Conclusions

By controlled converting MMA into MAA units through a hydrogenation process, a series of P(MMA–MAA) copolymers have been synthesized from PMMA and well characterized in this work. The introduction of MAA units may help to construct two possible H-bonds in P(MMA–MAA) copolymers. At low MAA content e.g. <3 mol%, weak H-bonds mostly formed between –OH and ester groups leads to slightly improved T_g and modulus. The weak H-bonds and better mobility of –CH₂OH side chain than that of –COOCH₃ results into the improved dielectric permittivity. As MAA content increases, strong H-bonds would be more likely formed between adjacent –OH groups, which are responsible for the significantly enhanced thermal and mechanical properties, and lowered dielectric permittivity. The breakdown strength, energy density, and energy releasing efficiency are thus remarkably improved. The highest U_e of P(MMA–MAA)-4 (bearing 19 mol% MAA) was measured to be 13 J cm⁻³@550 MV m⁻¹, which is 2–3 times larger than that of BOPP (about 4 J cm⁻³) at average breakdown strength of 600 MV m⁻¹. Whereas, U_l was maintained at as low as 8%@550 MV m⁻¹, which is rather close to BOPP under the consistent electric field. That allows P(MMA–MAA)s with optimized composition to be promising candidates for high pulse energy storage capacitors.

Acknowledgements

This work was financially supported by the National Nature Science Foundation of China (Grant No. 51573146, 51103115, 50903065), Fundamental Research Funds for the Central Universities (Grant No. XJJ2013075, cxtd2015003), International Science & Technology Cooperation Program of China (Grant No. 2013DFR50470), Natural Science Basic Research Plan in Shanxi Province of China (Grant No. 2013JZ003), and National Basic Research Program of China 973 (Grant No. 6132620101).

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Footnote

† Electronic supplementary information (ESI) available. See DOI: 10.1039/c6ra03757g

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