Multiphase Structure and Electromechanical Behaviors of Aliphatic

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Multiphase Structure and Electromechanical Behaviors of Aliphatic Polyurethane Elastomers Dong Xiang,† Jingjing He,† Tongtong Cui,§ Li Liu,‡ Qi Song Shi,† Lan Chao Ma,† and Yongri Liang*,† †

College of Materials Science and Engineering, Beijing Institute of Petrochemical Technology, Beijing 102617, P. R. China State Key Laboratory of Chemical Resource Engineering, and §Key Laboratory of Beijing City on Preparation and Processing of Novel Polymer Materials, Beijing University of Chemical Technology, Beijing 100029, China

Macromolecules Downloaded from pubs.acs.org by EASTERN KENTUCKY UNIV on 08/13/18. For personal use only.

‡

ABSTRACT: Understanding the relationship between multiphase structure and electromechanical property of thermoplastic dielectric elastomers is significantly important in the developments of high-performance and novel dielectric elastomers. In this work, we fabricated a series of aliphatic polyurethane elastomers (PUEs) based on hydroxyl-terminated butadiene-acrylonitrile copolymer (HTBN), hexamethylene diisocyanate (HDI), and various lengths of linear aliphatic diols and investigated effect of their microstructure and morphology on dielectric and electromechanical properties. The FTIR, WAXS, SAXS, and viscoelastic AFM results showed that the PUEs existed in crystalline phase, hard domain (HD) and soft domain (SD) phases composed HD-rich region and few HDs and SD composted SD-rich region by crystallization and microphase separation. Also, the crystal morphology and crystallinity of PUEs are strongly influenced by the length of chain extender due to the chain extender adopting various conformations by hydrogen bonding. The mechanical and electric fields induced responses of segment motions in PUEs at below room temperature were relative to the constrained SS motions from HD-rich and SD-rich regions. The electric field induced strain of PUEs was actuated by both Maxwell stress and electrostriction effect, of which contribution of electrostriction effect was above 64% in total actuation strain. In addition, we found that the effect of electrostriction on the actuation strain played an important role in improving the actuation strain of PUEs at lower electric field. Our results showed that the dielectric and electromechanical properties of PUEs can be adjusted by controlling the crystallization and microphase separation. polymer, z is the thickness of the elastomer film, and E is the electric field.11 The contribution of Maxwell stress on electric field induced thickness actuation strain, sZ,M, is given by

1. INTRODUCTION Dielectric elastomers (DEs) are one electronic type of electroactive polymer material that have attracted much interest due to their outstanding performances, including short time response, high energy density, light weight, high coupling efficiency, large strain, processability, as well as flexibility.1−3 Those various excellent properties give them potential applications in actuators,4 power generators,5,6 artificial muscles,7,8 and energy harvesters.9,10 In general, dielectric elastomers electrically actuate through the two actuation mechanisms of Maxwell stress and true electrostriction effect. The Maxwell stress (σM) is due to the interaction between the free charges on the electrodes (Coulomb interaction) and to electrostatic forces that arise from dielectric inhomogeneity. Maxwell stress is given by σM = ε0εrE2 = ε0εr

2

( Uz ) , where ε

0

εε E 2

sZ,M = − Y0 , where the Y is Young’s modulus. The electrostriction effect (σE) is known as direct coupling between the polarization and mechanical response in the material. The contribution of electrostriction effect on electric field induced thickness actuation strain, sZ,E, is given by sZ,E = Qε20(ε − 1)2E2, where the Q is coefficient of electrostriction effect.12 It has been recognized that the electric field induced actuation behavior of thermoplastic dielectric elastomers such as polyurethanes (PUs),13 poly(styrene-b-ethylene-co-butylene-b-styrene) (SEBS),14 and poly(styrene-b-butyl acrylate-bReceived: June 7, 2018 Revised: July 31, 2018

(8.85 × 10−12 F/m) is the

vacuum permittivity, εr is the dielectric constant of the © XXXX American Chemical Society

A

DOI: 10.1021/acs.macromol.8b01171 Macromolecules XXXX, XXX, XXX−XXX

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Macromolecules styrene) (SBAS)15 is considerably different from that of a conventional chemical cross-linked dielectric elastomer such as acrylic rubbers4 and silicone rubbers.16−19 The Maxwell stress contribution is dominant in the conventional dielectric elastomers,11,20 whereas the strong true electrostriction effect contribution might be dominant in the thermoplastic dielectric elastomers. In addition, the thermoplastic elastomers can be easily processed and changed into various shapes and domain orientations repeatedly at temperature above glass transition (Tg) or melting of hard block, unlike the chemically crosslinked ones which cannot change the shape or orientation once chemically cross-linked. However, the mechanisms of electric field induced actuation behavior have not been fully understood in the nanostructured thermoplastic dielectric elastomers including the aspects of the relationships between multiphase structure and dielectric property, electrostriction, and electromechanical coupling. Up to now, some research works have been focused on understanding the relationship between structure and electromechanical properties of thermoplastic dielectric elastomers. For example, Zhang et al.12 investigated the electromechanical responses of polyurethane elastomer (PUE) at room temperature and in the temperature range near its glass transition. They found that the Maxwell stress contribution to the strain response can be significant at temperatures higher than the Tg. In addition, they found that the chain segment motions can be divided into those related to the polarization response and those related to mechanical response, and the overlap between the two yields the electromechanical response of the material. In their work, the longitudinal electrostrictive coefficient Q11 of PU was −150 to −450 m4/C2. Cho et al.21 fabricated methacrylate-based triblock copolymers (PTMDMT) using poly(dodecyl methacrylate) (PDMA) as soft segment, and three different hard block segments with poly(methyl methacrylate) (PMMA), poly(tert-butyl methacrylate) (PtBMA), and their random copolymers (PMTDTMTs) as hard segment. Their results showed that these triblock copolymers had a variety of morphologies affecting their mechanical (elastic modulus) and electrical (dielectric constant) properties, leading to a tuning of their electromechanical properties. The neutral PTMDMT showed the largest transverse strain (0.658% at 50 V μm−1) under lower electric field because of its low modulus originating from its small domain size in spite of minor changes in its dielectric constant, while the acidified PTMDMT series demonstrated poor electromechanical performance due to their high modulus originating from hydrogen-bonding interactions of their acid groups, although their dielectric constants increased. Kim et al.22,23 investigated the electrostriction effect of SEBS gels and its modified copolymers by synchrotron radiation small-angle X-ray scattering technique. The results indicated that the introduction of nanostructured microdomains with mismatched dielectric constant will result in an unexpected superlarge electrostrictive effect (−7.4 × 105 to −8.4 × 104 m4/C2), causing a large electric field actuation strain response at low electric field. Also, they found that the strong electrostriction of SEBS gels was mainly caused by a strong true electrostrictive effect from the high density of the interfaces between dielectric mismatched nanostructured domains. As is well-known, polyurethane elastomer (PUE) is one kind of thermoplastic elastomers which is composed of soft and hard segments (HS and SS). PUE is a promising elastomer

material for DE applications due to its many advantages such as easily controllable structure, good biocompatibility, mechanical and dielectric properties and processability, etc.24 More recently, we investigated the relationship between structure and dielectric property of hydroxyl-terminated butadiene-acrylonitrile copolymer (HTBN) based PUEs.25−27 Our previous results indicated the dielectric constant and dissipation factor of polyurethanes were not only dependent on the dipole orientation of hard and soft segments but also strongly dependent on their microstructure including degree of microphase separation and crystallinity.28 However, the contributions of Maxwell stress and electrostriction effect on electric field induced strain response in crystallizable aliphatic PUEs have not been fully understood. Since motions of HSs and SSs in PUEs are strongly constrained by their multiphase and multilength scale structures, the mechanical and electric fields induced strain responses should be more complex. The crystal structure and degree of hydrogen bonding of HSs of aliphatic PUs can be adjusted by HS chemical structure including both chemical structure of diisocyanate and chain extender. Saito et al.29 investigated the crystal structure and conformation of aliphatic polyurethanes with various chain length of diols (m,6-polyurethanes, m = 2−6). The m,6polyurethanes can be descripted as the formula where m and n

represent the number of methylene groups in diols and diisocyanate, respectively, and R is number of repeat unit. The results showed that the crystal structures of m,6-polyurethanes with even number of m (i.e., m = 4 and 6) have planar zigzag conformation with perfect hydrogen bonding, whereas the crystal structures of m,6-polyurethane with odd number of m (i.e., m = 3) have twisted conformation (tilting the urethane group planes approximately 30° with respect to the fiber axis). They also reported that the 2,6-polyurethane has two types of crystal structure including type I crystal structure with planner zigzag conformation and type II crystal structure with twisted conformation. In this work, we adjusted the crystallized and microphaseseparated structure of PUEs via change in the length of chain extender (i.e., methylene group members (m) of aliphatic diol, m = 2, 4, 6) and investigated the effect of chain extender length on crystallization, microstructure, static and dynamic mechanical properties, and dielectric and electromechanical properties. The correlations between multiphase structure and segmental dynamic in PUEs were discussed, and the contributions of Maxwell stress and electrostriction effect on electromechanical behaviors of PUEs were also investigated.

2. EXPERIMENTAL SECTION 2.1. Materials. Hydroxyl-terminated polybutadiene-acrylonitrile copolymer (HTBN) which has 13.9 wt % cyano content and 0.6149 mmol g−1 hydroxyl value was supplied by Zibo Qilong Chemical Industry Co. Ltd., China. The average molecular weight of HTBN is 3500 g/mol. Hexamethylene diisocyanate (HDI) with 99% purity was purchased from Aladdin Industrial Co. Ltd. The various chain extenders of ethylene glycol (EG) (99% purity, Acros Co. Ltd.), butane diol (BDO) (99% purity, Ark Co. Ltd.), and hexane diol (HDO) (98% purity, Adamas Reagent Co. Ltd.) were removed in water using the 4A type molecular sieve (Tianjin Fuchen Chemical Reagents Factory) before use. The chemical structures of raw materials are shown in Table 1. B

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Macromolecules Table 1. Chemical Structure of Raw Materials for Synthesis of PUEs

Table 2. Compositions of Synthesized PUEs and Corresponding Sample Codes sample code

molar ratio of n(NCO)HDI/n(OH)HTBN/ n(OH)chain extender

HS wt %a

EG-PUE BDO-PUE HDO-PUE

8.89/1/7.89 8.89/1/7.89 8.89/1/7.89

37.6 40.0 42.2

a

HS wt % =

mHDI + mchain extender mHTBN + mHDI + mchain extender

× 100%, where HS wt % is

weight fraction of hard segment in PU and mHDI, mchain extender, and mHTBN are weight of HDI, chain extender, and HTBN, respectively. 58.380 Å) powder was used as standard material to calibrate the scattering angle. 2.5. Thermal Analysis. Differential scanning calorimetry (DSC) thermograms were obtained by Q2000 thermoanalyzer system (TA Corp. USA) with 10 °C min−1 heating and cooling rate under nitrogen (N2) atmosphere. Each sample for DSC measurements weighed about 5−10 mg. The thermogravimetric analysis (TGA) was conducted on a Pyris 1 thermoanalyzer system (PerkinElmer Corp. USA) with 10 °C min−1 heating rate under N2 atmosphere. Each sample weighed about 3−5 mg for TGA measurement. 2.6. Atomic Force Microscope (AFM). The surface morphology of PUEs was observed by the MFP-3D Origin AFM system (Oxford Instruments Co. USA) with AM−FM viscoelastic mapping mode. The AFM tips (AC160TS-R3, Asylum Research) with 26 N/m spring constant and 300 kHz resonance frequency were used for measurements. 2.7. Measurement of Temperature−Frequency Dependent Dielectric Property. The temperature and frequency dependent dielectric properties of PUEs were measured by Agilent 4294A precision impedance analyzer (Agilent Technologies Co.) at the temperature range from −40 to 120 °C with 102−107 Hz frequency range. The samples were prepared as square shape with 4.5 mm side length and 100 um thickness. The silver pastes were coated on both sides of sample films as electrodes. 2.8. Measurements of Strain−Stress Curves. Tensile testing of PUEs was completed by a universal testing machine (Instron 3366, Instron Cor. USA) with 20 mm min−1 stretching rate at 25 °C. The tensile specimens were prepared as dumbbell shape. The thickness, length, and width of narrow parallel portions of samples were 1, 12, and 2 mm, respectively.

2.2. Synthesis of PUEs. The PUEs were synthesized in two steps of polymerization method as schematically shown in Figure 1. In brief, the excess HDI and HTBN reacted to form isocyanate-terminated prepolymer in the first step at 70 °C for 2 h, and then the prepolymer reacted with diol to form PUE in the second step at 60 °C for 2 h as shown in Figure 1. The final products were then cured in the vacuum oven at 80 °C for 12 h. The PUEs were pressed at 200 °C to prepare films for testing. The compositions of PUEs were denoted as listed in Table 2. 2.3. Fourier Transform Infrared Spectroscopy (FTIR). The FTIR spectra of PUEs were collected by Fourier transformation infrared spectrometer (TENSOR 27, Bruker Co.) using attenuation total reflection (ATR) mode (Golden Gate single reflection diamond ATR, Specac.com) with 32 scans and 2 cm−1 resolution. The GRAMS/AI program (Thermo Galactic Inc.) was used for deconvolution of FTIR spectra peaks assuming infrared peaks with Gaussian distribution. 2.4. Synchrotron Small/Wide-Angle X-ray Scattering (SAXS/ WAXS). The SAXS measurements were carried out at the BL16B beamline in the Shanghai Synchrotron Radiation Facility (SSRF), China. The X-ray wavenumber was 0.124 nm, and a Mar165 CCD detector (2048 × 2048 pixels with a pixel size of 80 μm) was employed to collect the SAXS data. The ox tendon (D = 63.6 nm) was used as standard material for calibration of the scattering vector. The WAXS experiments of PUEs were carried out at the 1W1A Beamline in Beijing Synchrotron Radiation Facility (BSRF) with 1.5473 Å wavelength and Mar345 detector. Silver behenate (d001 =

Figure 1. Schematic illustration of the two-step polymerization process used in this work. HS: hard segment, SS: soft segment, HTBN: hydroxylterminated polybutadiene-acrylonitrile copolymer, HDI: hexamethylene diisocyanate. C

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Macromolecules 2.9. Measurements of Electromechanical Property. The electric field induced actuation area strains of PUEs were determined by the circular shape of a homemade device with wide-angle lens emplaced camera. The graphite, silicone oil, and curing agent30 composited conductive slurry composite was used as compliant electrodes to be coated on both sides of PUEs film. The actuation area strain was calculated by analysis of captured video images. The thickness of PUEs films was about 0.1−0.2 mm. 2.10. Dynamic Mechanical Thermal Analysis (DMTA). The dynamic mechanical properties of PUEs were measured by DMA Q800 instrument (TA Corp. USA) at 0.1% of strain and 1 Hz of frequency. The sample was heated from −80 to 150 °C with 3 °C min−1 heating rate. The samples were prepared as rectangular shape with 1 mm thickness, 12 mm length, and 4 mm width, respectively.

3. RESULTS AND DISCUSSION 3.1. Structure Characterization of PUEs. To understand the chain extender length effect on crystal structure and microphase separation of PUEs, we measured the WAXS of PUEs as shown in Figure 2. In Figure 2, the EG-PU of the

Figure 3. DSC thermograms of PUEs with various lengths of chain extender obtained during the second heating process with 10 °C min−1 heating rate.

crystal of 4, 6-polyurethane is triclinic crystal unit cell with parameters a = 4.98 Å, b = 4.71 Å, c = 19.4 Å, α = 116°, β = 105°, and γ = 109°, and the crystal of 6, 6-polyurethane is triclinic crystal cell with parameters a = 5.05 Å, b = 4.54 Å, c = 21.9 Å, α = 112°, β = 108°, and γ = 108°. Based on crystalline parameters of m,6-aliphatic polyurethanes, we can assign the crystal structure of EG-PUE to crystal type II of 2, 6polyurethane, and the diffraction peak at 6.7° (corresponding to 13.3 Å of interplane distance) can be assigned to the (001) plane. The chain conformations in the crystalline phase of m,6aliphatic polyurethanes adopted planar zigzag conformation with perfect hydrogen bonding for the m = 4 and 6 and type I of m = 2, whereas they adopted twisted conformation (tilting the urethane group planes approximately 30° with respect to the fiber axis) with deficient hydrogen bonding for the type II of m = 2.29 The chain conformations of PUEs were determined by FTIR as shown in Figure 4A. As shown in Figure 4A, the band at 1464 cm−1 is contributed from HDI residuals. Therefore, we can determine that the band at 1476 cm−1 should be contributed from aliphatic diol residuals. The bands at 1476 and 1464 cm−1 can be assigned to CH2 wagging vibration mode. Accordingly, the relative absorbance of band at 1476 cm−1 should be dependent on the methylene group numbers of aliphatic diol. In other hand, the band at 1475 cm−1 is assigned to trans conformer31−33 in the EG-based polyesters such as poly(ethylene terephthalate) and poly(ethylene 2,6-naphthalate). Similarly, we can assign the band at 1476 cm−1 to trans conformer. The stretching vibration bands of C−O from diol residuals of PUEs are observed at around 1048 and 1065 cm−1 which are assigned to gauche and trans conformers, respectively.32,34 Interestingly, the relative absorbance of the band at 1048 cm−1 is higher than that of the band at 1065 cm−1 in EG-PUE, whereas the relative absorbance of the band at 1048 cm−1 is lower than that of the band at 1065 cm−1 in BDO-PUE. It give evidence for the C−O groups in EG-PUE and BDO-PUE having different conformations. Since the molar ratio of chain extender and HTBN is 7.89:1, the absorbance of C−O groups is mainly contributed from the chain extender of C−O groups. Accordingly, we can determine that the glycol residual (−O− CH2−CH2−O−) in EG-PUE should preferentially adopt gauche−trans−gauche conformation to form hydrogen bonding, whereas the BDO residual in BDO-PUE should

Figure 2. WAXS profiles of PUEs with various lengths of chain extender.

WAXS profile shows weak crystalline diffraction peaks at 6.7 and 22.8° which correspond to 13.3 and 3.9 Å interplanar spacing, respectively. The WAXS profile of BDO-PU shows crystalline diffraction peaks at 20.5, 22.2, and 24.1°, and the WAXS profile of HDO-PU shows crystalline diffraction peaks at 11.8, 19.8, 21.6, and 23.7°. The results indicate that the crystalline phase and crystallinity of PUEs are dependent on the length of chain extender. The melting temperatures (Tm) of EG-PUE, BDO-PUE, and HDO-PUE are observed at 138, 164, and 153 °C in the DSC thermosgrams, respectively, as shown in Figure 3. The melting enthalpies of EG-PUE, BDOPUE, and HDO-PUE are calculated as 12.5, 21.6, and 26.4 J/g, respectively. In addition, the EG-PUE shows an obvious cold crystallization peak at 65.2 °C and a glass transition temperature at 28.9 °C. Those results demonstrate that the EG-PUE has lower crystallinity and weaker crystallizability than BDO-PUE and HDO-PUE. On the other hand, the crystalline structure of PUEs can be determined based on the crystalline structure of m,6-polyurethanes (m = 2, 4, and 6) reported by Saito et al.29 The crystal type I of 2,6-polyurethane is triclinic crystal cell with parameters a = 4.93 Å, b = 4.58 Å, c = 16.8 Å, α = 113°, β = 103°, and γ = 109°, whereas crystal type II of 2,6polyurethane is triclinic crystal cell with parameters a = 4.59 Å, b = 5.14 Å, c = 13.9 Å, α = 90°, β = 90°, and γ = 119°. The D

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Figure 5. SAXS profiles of PUEs with various lengths of chain extender.

corrected SAXS data can provide information on the average long period (L) between HD and soft domain (SD) along the direction of lamellar stack. Accordingly, the long period of lamellar stacks can be calculated by L = 2π/qmax (where qmax is the q value at maximum intensity of scattering peak). In Figure 5, the Lorentz-corrected SAXS profiles of PUEs show two scattering peaks (q1 and q2) around the regions of 0.015−0.040 and 0.050−0.090 Å−1, respectively. It means that the two period structures may coexist in the PUEs. In our previous work,26 we demonstrated that the q1 and q2 in the HDI, HTBN, and BDO composited PUEs are assigned to crystallized lamellar structure and microphase-separated structure, respectively, because the peaks of q1 and q2 do not have correlation between them and the q1 is dependent on the crystallinity. It means the PUEs have multiphase structure including crystalline phase and microphase-separated HD and SD phase. For BDO-PUE, the long period of crystallized lamellar structure and microphase-separated structure related scattering peak appear at around 0.032 and 0.065 Å−1, respectively, corresponding to 19.6 and 9.7 nm. In addition, the relative intensity of peak q1 is more intense than that of peak q2 in the EG-PUE compared to BDO-PUE. It indicates that the EDO-PUE has relatively lower degree of microphase separation than BDO-PUE and HDO-PUE because the relative lower scattering intensity of q2 in EG-PUE is caused by lower electron density difference between HD and SD. The morphology of PUEs was observed by AFM with AM− FM viscoelastic mapping mode as shown in Figure 6. The viscoelastic AFM technique can probe multiphase morphology of PUEs based on high resolution of mechanical property difference. In the Figure 6C, the BDO-PUE and HDO-PUE show the higher and lower modulus domains periodically existing morphology with a few tens of nanometers, whereas the EG-PUE shows higher modulus domains aggregated morphology with a few tens of nanometers. Based on SAXS results, we can determine that the periodically existing morphology and aggregated morphology are contributed from crystallization of HS. In comparison with BDO-PUE and HDO-PUE, the formation of aggregated morphology in EG-PUE may be caused by the gauche conformation of C−O for hydrogen bonding. In addition, a few tens of nanometers to a few hundreds sizes of relatively higher modulus region (called HD-rich region) and lower modulus regions (called SD-rich region) are observed in AFM modulus images as shown in Figure 6B and C. Therefore, we suggested that the

Figure 4. FTIR spectra of PUEs with various lengths of chain extender displayed at (A) 1600−800 cm−1, (B) 1760−1660 cm−1.

preferentially adopt the all-trans conformation from hydrogen bonding. Similarly, the C−O conformation of EG in the diphenylmethane diisocyanate (MDI), poly(tetramethylene adipate) (PTMA), and EG composted PUEs was assigned to gauche conformers.35 It was explained that the short chain of EG is too short to permit packing of the MDI units in the same way as for the longer chain extenders. In our case, the gauche conformation of C−O linkage in EG-PUE may be caused by titling the urethane group planes approximately 30° with respect to the fiber axis to form hydrogen bonding. However, relative absorbance of bands at 1048 and 1065 cm−1 in HDOPUE appear similar. It means that the C−O in HDO-PUE adopts both gauche and trans conformers to form hydrogen bonding. The infrared peaks of CO appear in the region of 1760− 1660 cm−1 as shown in Figure 4B. The absorbance bands of CO can divide into three peaks at 1717, 1701, and 1687 cm−1, respectively, which can be assigned to free CO, hydrogen-bonded (H-bonded) CO in disordered phase, and H-bonded CO in ordered phase, respectively.36 The absorbance fraction of H-bonded CO (f H) in PUEs can A be calculated by fH = A +HA , where Af and AH represent the f

H

absorbances of free CO and H-bonded CO (including both H-bonded CO from ordered and disordered phases), respectively. The f H values of EG-PUE, BDO-PUE, and HDOPUE were calculated as 0.80, 0.86, and 0.87, respectively. In addition, the wavenumbers of H-bonds CO in EG-PUE and BDO-PUE (or HDO-PUE) are 1688 and 1684 cm−1, respectively. The results indicate that the hydrogen-bonding strength of interhard segments in EG-PUE is slightly weaker than that in BDO-PUE and HDO-PUE. It may be owing to the less compact packing of hydrogen-bonded HSs caused by gauche conformer. The microphase-separated morphology of PUEs was determined by SAXS as shown in Figure 5. The LorentzE

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Figure 6. AFM images with AM−FM viscoelastic mapping mode of PUEs with various lengths of chain extender: (1) EG-PUE, (2) BDO-PUE, (3) HDO-PUE; (A) height image, (B) modulus image, (C) amplified modulus image.

HD-rich region should be composed of crystallized lamellar structure with crystalline and amorphous phases and microphase-separated structure with HD and SD phases. Similarly, the SD-rich regions should contain a few HD phases depending on the degree of microphase separation. As far as we know, such complex morphology of PUEs has not been reported yet. The viscoelastic AFM method is a new powerful tool to understand the multiphase structure of PUEs. The mechanism of multiphase structure formation in crystalline PUEs needs more detailed studies. Accordingly, the SSs should exhibit different responses by mechanical and electric fields due to the SSs constrained by various spaces of HD-rich and SD-rich regions. 3.2. Dynamic Mechanical Responses of PUEs. To understand the SS and HS motions of PUEs, the storage modulus and tan δ curves of PUEs were obtained by DMTA as shown in Figure 7. The storage modulus of PUEs is sharply decreased at around −50 °C and then continuously decreased until about 150 °C with increase of temperature. Accordingly, strong relaxation peaks appear in the tan δ curves of PUEs at around −36 °C which can be assigned to Tg of SS in SD-rich regions. The Tg values of SS in EG-PUE, BDO-PUE, and HDO-PUE are determined by DMTA as −36.5, −36.5, and −38.8 °C, respectively. The DSC results (Figure 3) show that the Tg values of SS in EG-PUE, BDO-PUE, and HDO-PUE are about −38.8, −42.4, and −38.1 °C, respectively.

Figure 7. Temperature-dependent storage modulus (A) and (B) tan δ curves of PUEs with various lengths of chain extender.

F

DOI: 10.1021/acs.macromol.8b01171 Macromolecules XXXX, XXX, XXX−XXX

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Figure 8. Temperature−frequency dependent dielectric constant (A) and dielectric loss (B) of PUEs with various lengths of chain extender: at (1) 0.1 kHz, (2) 1 kHz, and (3) 10 kHz.

29.6 and 20.9 °C in the tan δ curves of BDO-PUE and HDOPUE are contributed from SS motions in HD-rich regions, respectively. 3.3. Dynamic Electric Responses of PUEs. The electric field induced dipole polarization responses of PUEs were determined by temperature−frequency dielectric property measurements in the temperature range from −40 to 140 °C at 0.1, 1, and 1 kHz frequency, respectively, as shown in Figure

In addition, two weak relaxation peaks appear in the tan δ curve of EG-PUE at 34.4 and 71.0 °C, whereas one weak relaxation peak appears at 29.6 and 20.9 °C in the tan δ curves of BDO-PUE and HDO-PUE, respectively, as shown in Figure 7B. Based on SAXS and AFM results, we can determine that the relaxation peaks at 34.4 and 71.0 °C in the tan δ curve of EG-PUE may be contributed from the SS and HS motions in HD-rich regions, respectively. Similarly, the relaxation peaks at G

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increased with increasing the methylene group numbers of chain extender. The break tensile strength, Young’s modulus, and break elongation of PUEs are summarized in Table 3. The

8. At 0.1 Hz frequency, the dielectric constants of PUEs are increased with increase of temperature up to about 20 °C and then are reincreased at above 100 °C as shown in Figure 8A1. With increase of the frequency to 1 or 10 kHz, the trends of temperature-dependent dielectric constant change of PUEs are similar to that in 0.1 kHz frequency as shown in Figure 8A2 and A3. Also, the dielectric constant of PUEs is increased more rapidly with increase of temperature in the temperature range from −40 to 20 °C at higher frequency. The dielectric loss of PUEs at 0.1 kHz shows minimum values at around 20 °C as shown in Figure 8B1. The minimum values of dielectric loss in PUEs are shifted to higher temperature as increase of frequency. In addition, dielectric loss relaxation peaks appear at −20 °C and 10 kHz frequency. The DMTA and AFM results indicated that the Tg of SS in SD- and HD-rich regions appeared at around −38 and 20 °C, respectively. However, HS motions in HD-rich regions require higher temperature for orientation of dipoles to polarization. Therefore, the dielectric constants of PUEs at the temperature range from −40 to 20 °C are dominantly contributed from dipole polarization of polar groups from SSs in SD-rich and HD regions. On the other hand, the EG-PUE showed lower dielectric constant than BDO-PUE and HDO-PUE at all temperature and frequency ranges. It may be owing to inability of dipole polarization of constrained SS and HS by larger HDs in the frequency ranges. On the other hand, the dielectric constant of EG-PUE, BDOPUE, and HDO-PUE is 6.08, 7.92, and 6.78 at 1 kHz and 20 °C, respectively. The higher dielectric constant of BDO-PUE may be caused by period nanostructured structure.23 3.4. Strain−Stress Curves of PUEs. The strain−stress curves of PUEs are shown in Figure 9A. All PUEs show at least 450% of strain, and the break tensile strength of PUEs is

Table 3. Dielectric Constant and Dissipation Factor, Breakdown Strength, Mechanical Properties, and Electromechanical Properties of PUEs property

EG-PUE

BDO-PUE

HDO-PUE

dielectric constant @1 kHz and 20 °C dissipation factor @1 kHz Young’s modulus (MPa) break tensile strength (MPa) break elongation (%) actuation area strain, SA (%) @ 40 kV/mm Maxwell stress (σM) (MPa)a actuation area strain, sM,A, by Maxwell stress (%)b contribution of electrostriction effect on actuation area strain, f E (%)c

6.08

7.92

6.78

0.040 7.8 ± 0.4 9.5 ± 0.2 972 ± 36 3.2

0.026 16.3 ± 0.8 11.7 ± 0.5 447 ± 18 1.8

0.027 16.6 ± 1.5 15.6 ± 1.5 470 ± 22 2.4

0.086 1.10

0.112 0.64

0.096 0.58

66

64

76

a

Calculated by σM = ε0εrE2 at 40 kV/mm, 1 kHz, and 20 °C.

b

Calculated by sM,A =

(

1 1 + sZ,M

)

− 1 × 100%; sZ,M = −

kV/mm, 1 kHz, and 20 °C. Calculated by fE = 1 − c

40 kV/mm, 1 kHz, and 20 °C.

sM,A sA

εε0E 2 Y

at 40

× 100% at

EG-PUE shows a Young’s modulus of 7.8 MP, which is significantly lower than those of BDO-PUE (16.3 MPa) and HDO-PUE (16.6 MPa). As is well-known, the mechanical properties of PUE are strongly influenced by the stability of their microphase separation in formed HDs and HS contents. To understand the real stress changes during deformation, the true stress−true stain curves of PUEs are shown in Figure 9B. In the initial region of the true stress and true strain curves (true strain E > 25 kV/mm)

electrostriction effect coefficient, Q (m4/C2), at 1 kHz and 20 °C

−3.06 × 10−17

−1.29 × 10−17

−1.22 × 104

−1.55 × 10−17

−9.00 × 10−18

−2.92 × 103

−2.21 × 10−17

−8.82 × 10−18

−8.48 × 103

I

DOI: 10.1021/acs.macromol.8b01171 Macromolecules XXXX, XXX, XXX−XXX

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Macromolecules

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whereas HD-rich regions of BDO-PUE and HDO-PUE showed alternative period morphology. The mechanical and electric fields induced responses of SS motions in PUEs were strongly influenced by the various SSs constrained structures. The electric field induced strain of PUEs was actuated by both Maxwell stress and electrostriction effect, of which contribution of electrostriction effect was above 66% in total actuation strain at 40 kV/mm, 1 kHz, and 20 °C. In addition, we found that the effect of electrostriction on the actuation strain played an important role in enhancement of the actuation strain of PUEs at lower electric field. Among PUEs, the EG-PUE showed −3.06 × 10−17 m2/V2 of electromechanical coefficient and 1.22 × 104 m4/C2 of electrostriction effect coefficient at 1 kHz and 20 °C. Our results can provide new insight on understanding the relationship between multiphase structure and electromechanical properties.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Yongri Liang: 0000-0003-3576-613X Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This research work was supported by Natural Science Foundation of Beijing Municipality (2172023) and the Promotion of Innovation in Beijing Universities Fund (TJSHG201510017024). The 2D WAXS experiments were supported by Beijing Synchrotron Radiation Facility (BSRF) in China. The SAXS experiments were supported by Shanghai Synchrotron Radiation Facility (SSRF) in China. We thank Prof. Zhi-min Dang (Tsinghua University, China) for helping us with measurements of temperature−frequency dielectric properties.

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K

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