Fabrication of In Situ Nanofiber-Reinforced Molecular Composites by

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Applications of Polymer, Composite, and Coating Materials

Fabrication of in-situ nano-fibers reinforced molecular composites by non-equilibrium self-assembly Qingbao Guan, Li Yuan, Aijuan Gu, and Guozheng Liang ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b15037 • Publication Date (Web): 19 Oct 2018 Downloaded from http://pubs.acs.org on October 20, 2018

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ACS Applied Materials & Interfaces

Fabrication of in-situ nano-fibers reinforced molecular composites by non-equilibrium self-assembly Qingbao Guan, Li Yuan*, Aijuan Gu, and Guozheng Liang Jiangsu Key Laboratory of Advanced Functional Polymer Design and Application, Department of Materials Science and Engineering, College of Chemistry, Chemical Engineering and Materials Science, Soochow University, Suzhou, 215123, P. R. China ABSTRACT Although the concept of molecular composites (MCs) is very promising, there are major obstacles arising from the immiscibility of the rigid-rod with the random-coil polymers. Here we developed a novel method for fabricating in-situ reinforced MC system with non-equilibrium self-assembled nano-fibrous structures based on bisphenol A epoxy resin, 4,4′-diaminodiphenylsulfone, bismaleimide, and polyphenylene ether (PPO) oligomer. A variety of spectroscopic and morphological techniques were used to probe the structural evolution from the emergence of nano-fibrils, to growth and aggregation of nano-fibers, and then to formation of in-situ reinforced MC with strong interfacial interaction. The in-situ nano-fibers within polymer matrix could be formed by polymerization force extruding PPO phase through the interspaces within simultaneous interpenetrating network (SIPN) polymers during the cure process of thermosetting resin system. As compared to the control sample, the in-situ nano-fibers reinforced MC exhibited better thermal property and flame retardancy. Especially, the obtained MC showed significant improvement in glass transition temperature and mechanical properties, which were mainly attributed to the restriction of high thermal stability PPO on the segmental motion of

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polymer chains, the toughening and reinforcement behaviors of PPO nano-fibers on the matrix and the chemical interaction at PPO/matrix interface. Keywords: in-situ nano fibers, thermosetting resins, non-equilibrium self-assembly, thermoplastic resin, self-reinforcement, mechanical property, thermal property 1. INTRODUCTION Molecular composites (MCs) consist of rigid-rod molecules dispersed in a matrix of a relatively flexible or semiflexible coil polymer.1 The rigid-rod component, namely molecular fiber acts as reinforcement, which is similar to those in conventional fiber reinforced composites. Because of the large aspect ratio of the molecular fibers and the maximizing interaction between the fibers and the matrix on a molecular level, the advantages of MCs over conventional macroscopic fiber reinforced composites including enhanced thermal, thermoxidative, and hydrolytic stability, and higher specific mechanical properties.2 However, these synergistic effects can only be achieved when the rigid-rod molecular fiber is well dispersed and not phase separated from the matrix component.3 The homogeneous phase can be obtained only at a concentration of the molecular fibers lower than the critical concentration (2-4 wt%).4,

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At

higher concentrations the aggregation of the molecular fibers occurs, as predicted by Flory’s theory.6 The phase separation can be further promoted for the anisotropic rod-like component.7, 8 This heterogeneity deteriorates the mechanical performance of MCs down to the level of macroscopic composites. Possible routes have been developed in an attempt to address these challenges, including advanced synthesis approaches such as in-situ and precursor techniques,9, technologies,11,

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10

rapid preparation

incorporation of favorable interactions such as dipole-dipole, ionic, or

2


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hydrogen bonding between components,13-16 and molecular structure design such as block or graft copolymers.17, 18 However, we notice that the processability of MC is always problematic, as most of the materials are prepared in solution using toxic and/or strong solvents such as sulfuric acid, and can hardly be melt processed. The rigid-rod molecular fibers generally have a much higher melting point than the processing temperature of polymer matrix, and some of them even do not melt inherently.19 From the point of view of the thermodynamics of polymer blends, phase separation seems to be unavoidable during heating because the dispersion of the molecular fiber in the matrix is far from thermodynamically stable.20 Some attempts to prepare thermally processable systems have been reported for blends of aramid/poly(ether sulfone),21 aramid/amorphous nylon,22 and liquid crystalline copolyester/polycarbonate,23 whereas these blends can only be processed by hot compression and cold drawing since the melting point of the rigid-rod component is either much higher than the processing temperature or not observed, and the molecular chain orientation can be easily controlled by the external force . It is well-known that self-assembly is a powerful approach to construct a variety of structures of nano- to micrometer dimensions, such as nanotubes, vesicles, and capsules based on small molecules, polymers, proteins, colloids, and nanoparticles under thermodynamic equilibrium conditions.24-27 Boekhoven et al. reported a molecular fuel-driven self-assembly process that can control the formation of dispersed fibrous structures by tuning the kinetics of fuel conversion.28 However, most of these advances are based on biological or living thermodynamic framework: the target configuration minimizes thermodynamic free energy,29 which is easily violated by mechanical agitation, heating or other driving forces. One way to remedy this problem is to incorporate irreversible covalent bonds and cross-linked networks through cure reaction, permanent self-assembled morphology and structure can be obtained.30 Nevertheless, there is

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limited report on the cure reaction induced self-assembled molecular fiber in MCs. It is worthy to note that continuous energy supply during cure reaction maintains this process far from equilibrium state, leading to a dynamic morphology and structure development of polymers, which can be considered as a non-equilibrium thermodynamic process.31, 32 This brings up the question, can MCs be designed and prepared based on a thermoplastic/thermoset blend system? Our approach toward MCs makes use of thermoplastic/thermoset blends based on bisphenol A epoxy resin (EP), 4,4′-diaminodiphenylsulfone (DDS), bismaleimide (BMI), and polyphenylene ether (PPO) oligomer. The preparation relies on the formation of self-assembled nano-fibrous structure driven by the cure reaction of thermoplastic/thermoset blends. In this system, cure reaction process and morphological development, instead of equilibrium composition, determine flame retarding properties, thermo-mechanical properties of MCs such as Young’s modulus, glass transition temperature (Tg), and self-reinforcement capability. We will present deeper new mechanistic insight into non-equilibrium self-assembly and a more compelling demonstration of the advantages offered by this processing approach compared to that of existing MCs. 2. EXPERIMENTAL SECTION Materials Diglycidyl ether of bisphenol A epoxy resin (EP, epoxy equivalent: 185–190 g/eq) was purchased from Wuxi Resin Plant (China). Low-molecular-weight polyphenylene ether (PPO) with vinyl end-groups (phenolic end-group content ≤ 300 ppm, Tg = 154 C, Mn = 2200 g·mol-1) was obtained from SABIC Innovative Plastics. 4,4′-bismaleimidodiphenylmethane (BMI) was provided by Northwestern Chemical Engineering Institute (China). 4,4′-diaminodiphenylsulfone (DDS) catalyst was obtained from Alfa Aesar. Preparation of MC 4


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The mixture of EP, BMI, and PPO was heated to 130 °C, and when the mixture became clear liquid, DDS was added. The mixture compositions as used for our studies are summarized in Table 1. The mixture was stirred at 135 °C for 40 min. A homogenous prepolymer melt was obtained and then poured into a pre-heated mold, degassed at 135 °C for 1 h and cured following the procedure: 150 °C/2 h + 180 °C/2 h + 200 °C/2 h. The resultant MCs are coded as EP/DDS/BMI/PPO. For comparison, the control sample (coded as EP/DDS/BMI) was prepared by the identical procedure without PPO. Table 1. Backbone composition of control sample and MCs Prepolymer

EP/DDS/BMI EP/DDS/BMI/0.25PPO EP/DDS/BMI/0.50PPO EP/DDS/BMI/0.75PPO

Weight ratio EP DDS BMI PPO

1 1 1 1

0.35 0.35 0.35 0.35

1 1 1 1

0 0.25 0.50 0.75

Epoxy groups (mol·kg-1)

Amino groups (mol·kg-1)

BMI double bonds (mol·kg-1)

PPO double bonds (mol·kg-1)

2.17 1.96 1.79 1.65

1.20 1.09 0.99 0.91

2.36 2.15 1.96 1.80

0.09 0.16 0.22

Characterization The morphology of cryogenically fractured specimens was examined using a scanning electron microscopy (SEM, Hitachi S-4700, Japan). All samples for SEM experiments were sputter coated with a thin gold layer prior to examining. Some fractured surfaces were etched with toluene for 1 h before sputter-coating. The role of the toluene was used to get rid of the thermoplastic or PPO dominated phase. Atomic force microscopy (AFM) images and Young’s modulus of cryogenically fractured specimens were recorded using Bruker Multimode 8 AFM (Germany) with an aluminum reflex coated silicon cantilever probe (Tap 150AI-G, Budget Sensors, Innovative Solution Bulgaria Ltd., Sofia, Bulgaria). The tip has a half cone angle of 10 degrees, a radius of 2 nm, and a

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resonance frequency of 75 kHz. The nominal force constant is 3 N m-1, and the actual spring constant of the probe was calibrated before each test. A tapping mode (PeakForce Tapping) was employed to measure the Young’s modulus. The NanoScope Analysis 1.7 software was used for image analysis. Transmission electron microscope (TEM) images were recorded using a FEI Tecnai G-20 TEM (USA) at 200 kV. Ultrathin sections of approximately 70 nm in thickness were prepared using a Leica UC7-FC7 ultra microtome (Germany). In addition, thin membrane specimens of approximately 500-700 nm in thickness were also directly obtained by curing the EP/DDS/BMI/0.25PPO system for TEM observation. The cure behavior of the prepolymer was determined by differential scanning calorimetry (DSC) using a DSC Q2000 (TA Instruments) with a heating rate of 10 °C·min-1 in a nitrogen atmosphere. Fourier transform infrared (FTIR) spectroscopy was performed by scanning KBr disks of the samples from 400 to 4000 cm-1 (Nicolet, ThermoScience). The conversion (α) of functional groups in EP/DDS/BMI/PPO was analyzed based on their FTIR spectra, and could be calculated according to eqn. 1.



 = 1

Ax A1513   100% A0, x A0,1513 

(1)

where Ax and A1513 were the integral areas of characteristic peaks of functional groups and phenyl ring at 1513 cm-1 at different curing stages, and A0,x and A0,1513 were the integral areas of characteristic peaks and phenyl ring before cure. Dynamic mechanical analysis (DMA) was using a single cantilever beam clamping setup (TA Q800 DMA). DMA tests were carried out from 30 °C to 280 °C using a heating rate of 3 35 mm 10 mm  2 mm . The glass transition

°C·min-1 at 1 Hz. Sample dimension was 6


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temperature (Tg) was determined from the peak temperature in the loss factor (tan delta)-temperature plot. Thermogravimetric analyses (TGA) were performed from 30 °C to 800 °C at a heating rate of 10 °C·min-1 using nitrogen flowing at 50 mL·min-1 (TA Instruments SDTQ500) to study the thermal stability of specimen. Microcombustion calorimetery (MCC) was performed on a FTT0001 microscale combustibility calorimeter (UK). A 5 mg sample was heated to 700 °C at a heating rate of 60 °C·min−1 in a mixed stream of oxygen and nitrogen flowing at 20 and 80 cm3·min−1, respectively. Thermogravimetric

Analysis-Infrared

(TG-IR)

spectra

were

recorded

using

a

thermogravimetric analyzer (TGA F1, Netzsch, Germany) that was interfaced to a FTIR spectrophotometer (TENSOR 27, Bruker, Germany). Ten milligrams of a sample were put in an alumina crucible and heated from 40 to 800 °C with a heating rate of 10 °C /min under a nitrogen atmosphere, and the flowing rate was 45 mL·min-1. Flexural strength was measured using a KQL WDW100 (China) universal testing machine. The three-point bend fixture featured contact points with a 5 mm radius of curvature. Tests were conducted at 25 °C using a cross-head speed of 0.5 mm·min-1, and 10–15 specimens per composition were tested. Fracture toughness (KIC) was measured according to the ASTM standard D5045 using the KQL WDW100 (China) universal testing machine under mode I with a speed of 0.5 mm·min-1. Each specimen was tested using a standard single-edge notched beam. The size of the original crack was 0.45W < a < 0.7W. KIC was calculated according to eqn. 2.

K IC  Y

3PS a 2 BW 2 7

(2)



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where P was the critical load, B is the thickness, W was the width, S was the span, a was the size of the original crack, and Y was defined as: Y  1.93  3.07(

a a a a )  14.53( ) 2  25.11( )3  25.80( ) 4 W W W W

(3)

Tensile tests were conducted on a universal testing machine (MTS CMT-4104) according to Chinese standard GB/T2567-2008 using a cross-head speed of 2 mm·min-1. 10 dumbbell-shaped specimens for each composition were tested. Strain was measured by the digital image correlation (DIC) technique using a commercial system developed by GOM (Braunschweig, Germany). The optical measurement technique was similar to that described by Littell et al.33 Two cameras were connected to a computer equipped with software capable of pattern recognition. After a calibration procedure was completed, a test specimen with a speckle pattern applied to the surface was placed within the calibrated space. The position of any point on the surface of a test specimen could be determined by the software using the stereo images of the specimen. The strain was then calculated from surface displacements measured at specified time intervals during a test. 3.

RESULTS AND DISCUSSION

Morphology of MC The morphologies of control sample (EP/DDS/BMI) and MCs (EP/DDS/BMI/PPO) are investigated using scanning electron microscopy (SEM) and atomic force microscopy (AFM), respectively. The SEM images of fracture surface of EP/DDS/BMI after fracture toughness test show a smooth fracture surface (Figures 1A, 1A′ and 1A′′), whereas the fracture surfaces of EP/DDS/BMI/PPO systems are much rougher (Figures 1B-1D, 1B′-1D′ and 1B′′-1D′′). Interestingly, on close inspection of the rough area of EP/DDS/BMI/PPO, distinct nano-fibrous

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structure with a diameter of approximately 50 nm is observed (Figures 1B′′-1D′′), which indicates that the incorporation of PPO plays an important role in establishing in-situ nano-fibers structure. The in-situ nano-fibers contained MCs based on EP/DDS/BMI/PPO systems are indeed fabricated without hot compression and cold drawing.

Figure 1. SEM images of the fracture surfaces of EP/DDS/BMI (A, A′, A′′), EP/DDS/BMI/0.25PPO

(B,

B′,

B′′),

EP/DDS/BMI/0.50PPO

(C,

C′,

C′′)

and

EP/DDS/BMI/0.75PPO (D, D′, D′′) after fracture toughness test. To understand the length and aspect ratio of the in-situ-formed fibers, TEM measurements were performed on the thin membrane obtained directly by curing the EP/DDS/BMI/0.25PPO system. Figures 2A and 2B show the length of the in-situ-formed fiber within the range of 150– 450 nm, which is much shorter than that of the conventional fibers because the latter is highly aligned under external force or stretch. The evaluated aspect ratio for the fiber is 3–9. Figures 2C-2D show the AFM tapping mode (Peak Force Tapping) Young’s modulus image with corresponding representative cross sections of representative sample EP/DDS/BMI/0.25PPO. 9



The Young’s modulus values reflect the nature of materials. The modulus profile of random selected scan path in modulus image shows distinct modulus changes, of which the dark area (low modulus value) can be attributed to the fibers (PPO-rich domains) and the lighter area (higher modulus value) to the matrix (EP/DDS/BMI), because the thermoplastic PPO possesses less rigid molecular feature compared to EP/DDS/BMI matrix. Figure 2D shows that the variation of modulus is basically regular, then the width of the wave troughs of modulus curve can be regarded as the diameter of PPO fiber, which is approximately 50 nm.

Figure 2. TEM images of EP/DDS/BMI/0.25PPO thin membrane (A and B). AFM images of the fracture surfaces of cryogenically fractured EP/DDS/BMI/0.25PPO sample: Young’s modulus image (C) with corresponding representative cross sections (D). Self-assembled morphologies such as lamellar, cubic bicontinuous, hexagonally packed cylinders, body-centered cubic packed spheres, and disordered micelles are well established in the bulk for toughened thermosets and elastomers formed by step polymerization.30, 34 However, 10


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as far as we are aware, despite the block copolymers or polyblends containing liquid crystalline polymer phase that can be aligned to form micron-sized fibrous structure,35-37 the non-equilibrium self-assembled nano-fibrous structure induced by the cure reaction of thermoplastic/thermoset blend has not been reported before. Mechanism of the formation of in-situ fibers within polymer matrix To assess the in-situ fibrous structural evolution of non-equilibrium self-assembly, the reaction behaviors of EP/DDS/BMI/PPO prepolymer and the mixture of different raw materials were investigated using DSC (Figure 3). Although BMI monomers, DDS and PPO show melting points (Tm) at 160 °C, 180 °C and 250 °C, respectively (Figure 3A), homogenous prepolymer mixtures of EP/DDS/BMI and all EP/DDS/BMI/PPO systems are obtained after stirring at 135 °C for 40 min because of the good compatibility of BMI, DDS and PPO with EP, which can be indicated by no melting peaks in DSC curves of all prepolymers (Figure 3B) and the chemical reactions between raw materials (Figure 3C). The DSC curve of EP/DDS/BMI prepolymer shows two exothermic shoulder peaks at 284 °C and 309 °C (Figure 3B), which arise from BMI homopolymerization38, 39 and EP/DDS reaction40, 41 in a low temperature range, and BMI/DDS reaction42,

43

in a high temperature range as implied by their DSC curves (Figure 3C). The

possible reaction mechanisms are shown in Scheme 1. Because no reaction occurs between BMI and EP as indicated by the DSC curve of BMI/EP and the polymerizations of BMI monomer and EP resin can happen simultaneously as above mentioned, a so-called simultaneous interpenetrating network (SIPN) can be formed based on BMI homopolymerization and EP/DDS reaction networks that interlock and exhibit no or insignificant chemical bonding.44, 45 Musto P. et al. also have well investigated the kinetics and mechanism of the formation of SIPN structure based on EP/DDS/BMI by FTIR and DMA measurements,46 and their experimental results

11



indicate that BMI and EP/DDS pair follow two different and largely independent reactions and the two chemically different networks are likely to build up simultaneously during curing. The researches related to SIPNs of elastomers/polystyrene,45 BMI/polyether polyurethane-crosslinked epoxy,47 BMI/cyanate ester resin,38,48 castor oil or polyurethane/polyester49 can support our

Exo

(A) EP PPO BMI DDS

50

Heat Flow (W/g) Exo

findings as well.

Heat Flow (W/g)

(B) 284

309

246

EP/DDS/BMI

309 237

EP/DDS/BMI/0.25PPO 232

EP/DDS/BMI/0.50PPO EP/DDS/BMI/0.75PPO

100

150

200

250

300

350

Temperature (C)

50

100

150

200

250

300

350

Temperature (C)

Exo

(C) PPO/DDS

Heat Flow (W/g)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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EP/PPO EP/DDS BMI/DDS BMI/PPO BMI/EP

50

100

150

200

250

Temperature (C)

300

350

Figure 3. DSC curves of EP/DDS/BMI and EP/DDS/BMI/PPO prepolymers (A), BMI, EP, DDS and PPO (B) and their mixtures (C). In addition, the cured EP/DDS and BMI/DDS show a higher Tg of 219 °C and a lower Tg of 147 °C, respectively, whereas the cured EP/DDS/BMI displays only one peak at 156°C in tan delta-temperature curve (Figure S1), which indicates that two EP and BMI polymer networks

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exhibit a substantial degree of interpenetration and form SIPN structure.46 Here, it must be mentioned that BMI/DDS system cannot be completely cured for the low cure temperature (~200 °C) as indicated by the high exothermic peak of BMI/DDS at 285 °C (Figure 3C), and the T g value of pure cured BMI resin is not experimentally available in this work because of the high melting point (about 160°C) and the extensive degradation of the uncured BMI at the onset of large-scale molecular mobility. However, the T g value of system as function of BMI can be well interpolated by the Fox equation and the obtained T g in such a way is ~352 °C for a cured BMI resin. 46

O n

O N

N BMI

O

H3C

(A) O

O

CH3

H3C

OO

N

O

N

n

CH3 H2N

NH2

+ O

O

O OH

n

O

O

O

O

EP

OH

S

O

DDS

OH

+ EP

H N

OH N (B)

O

I

S

O

O

S

O

Epoxy resins-amine addition polymers EP/DDS

O N

HO

+

HO

O

HN

N

(C)

N BMI

O

I

O

S

O

O

O

S

O

Scheme 1. Possible reaction mechanisms of BMI homopolymerization (A), EP/DDS cure reaction (B) and reaction between EP/DDS and BMI (C).

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For EP/DDS/BMI/PPO systems, the introduction of PPO can shift the low exothermic peak to a low temperature and weaken the intensity of high exothermic peak of EP/DDS/BMI (Figure 3B), which is attributed to the fact that the C=C group of PPO is prone to react with the H–C= group in the maleimide of BMI via ‘ENE’-reaction at relatively low temperature as indicated by the lower exothermic peak at 215 °C in the DSC curve of the BMI/PPO mixture (Figure 3C),50 and the reaction mechanism of PPO oligomer with BMI is shown in Scheme 2.

O

O N

2

CH2

N O

+

O

H3C

R1

O

O

O

BMI

CH2 CH3 O PPO

O O

N

'ENE' reaction

CH2

O

CH2 O

H2C

O N

R1

O

O

O

O

C H2

N O O

(II)

N O

CH3 H3C

CH3

O

O

CH3 n

R1 = H3C CH3

m

Y

CH3 CH3

Scheme 2. Possible reaction mechanism of BMI/PPO ‘ENE’ reaction. To demonstrate the possible reactions in EP/DDS/BMI/PPO systems and the formation mechanism of in-situ fibers within matrix, the FTIR spectra and the conversion values of functional groups of the representative EP/DDS/BMI/0.25PPO system obtained at various curing stage are provided in Figure 4, and the schematic of the formation mechanism of in-situ fibers

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within matrix are shown in Figure 5. With the prolongation of curing time and the increase of curing temperature, EP/DDS/BMI/0.25PPO shows the gradual decrease in the intensities of characteristic absorptions of C=C group of PPO and H–C= group of BMI at 1706 cm−1 and 827 cm-1, respectively (Figure 4A). Regarding to the EP/DDS reaction, the intensities of absorption peak of epoxy group at 910 cm−1 and two well resolved absorption peaks of primary amine group at 3475 cm−1 and 3375 cm−1 simultaneously decrease as the cure reaction proceeds (Figure 4B). The conversions of H–C= group in BMI and C=C group in PPO are larger than that of amino and epoxy groups (Figure 4C). It is very evident that the polymerization reactions of BMI monomer and EP resin in EP/DDS/BMI/PPO system are also likely to lead to the formation of SIPN structure (Figure 5B), which is similar to the woven mesh network structure, and small interspaces must occur around the interlocked sites of BMI and EP polymer chains owing to no chemical interactions in BMI/EP, EP/PPO and PPO/DDS systems (Figures 3A and 3C). With the reaction going on in the system, more BMI and EP based polymers are produced and they are prone to be severely interlocked with each other, meanwhile the unreacted PPO phase domain is gradually separated (Figure 5C). As BMI and EP continues to react, polymer molecular weights gradually increase, the polymerization shrinkage force can extrude PPO phase as aligned fibrils through the interspaces within SIPN structures (Figures 5D and 5E). Owing to the reaction between PPO and BMI, the strong interface interaction at the interface of PPO/matrix can be realized.

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Figure 4. FTIR spectra of EP/DDS/BMI/0.25PPO, wavenumber ranges of 3500–1450 cm−1 (A) and 950–750 cm−1 (B), and the conversions (α) of functional groups in EP/DDS/BMI/0.25PPO prepolymer (C) at different prepolymerization stages: (a) 135 °C/0 h, (b) 135 °C/0.5 h, (c) 135 °C/0.75 h, and various curing stages: (d) 150 °C/1 h, (e) 150 °C/2 h, (f) 150 °C/2 h + 180 °C/1 h, (g) 150 °C/2 h + 180 °C/2 h, (h) 150 °C/2 h + 180 °C/2 h + 200 °C/1 h, (i) 150 °C/2 h + 180 °C/2 h + 200 °C/2 h. The thickness correction was accomplished by dividing the analytical band by the invariant peak of phenyl rings at 1513 cm−1.

Figure 5. Schematics of the formation mechanism of in-situ fibers within matrix: (A) prepolymer systems, (B) 150°C/1−2 h, (C) 150 °C/2 h + 180 °C/1−2 h, (D) 150 °C/2 h + 180 °C/2 h + 200 °C/1 h and (E) 150 °C/2 h + 180 °C/2 h + 200 °C/2 h. 16


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The morphological development of the representative EP/DDS/BMI/0.25PPO system obtained at various curing stage are monitored using SEM and TEM (Figures 6 and S2). In the early stage of curing (150 °C/1−2 h), EP/DDS/BMI/0.25PPO is a homogenous single-phase mixture (Figure S2). It can be observed that short nano-fibrils start emerging in EP/DDS/BMI/0.25PPO when the cure temperature reaches 180 °C (Figures 6A and 6A′), which is evidently attributed to the fact that the polymerization reaction of EP/DDS/BMI/0.25PPO system induces the formation of PPO fibrils. As the cure reaction proceeds, the molecular weights of polymers gradually grow and the more SIPN structures form, leading to more distinct nano-fibrous structure (Figures 6B, 6B′, 6C and 6C′). Eventually, an in-situ nano-fibers/matrix composites based on a miscible thermoplastic/thermoset mixture (EP/DDS/BMI/0.25PPO) is obtained at the end of cure protocol (Figures 6D and 6D′).

Figure 6. Morphological development of EP/DDS/BMI/0.25PPO induced by cure reaction. SEM and TEM images obtained at various curing stages: 150 °C/2 h + 180 °C/1 h (A, A′), 150 °C/2 h + 180 °C/2 h (B, B′), 150 °C/2 h + 180 °C/2 h + 200 °C/1 h (C, C′), and 150 °C/2 h + 180 °C/2 h + 200 °C/2 h (D, D′).

17



Page 18 of 39

Structure development during the cure reaction in various thermoplastic/thermoset binary blend systems (e.g., epoxy resin/polyether sulphone,51, 52 epoxy resin/polyphenylene ether53) has been widely investigated in the past decades. Although a miscible thermoplastic/thermoset mixture with specific interactions can be obtained at low critical solution temperature, demixing takes place via either cure reaction or spinodal decomposition, this results in a phase separated morphology of dispersed thermoplastic particles that subsequently toughen the thermoset matrix. For instance, Yamanaka et al. reported that the co-continuous phase connectivity of a ternary thermoplastic/thermoset mixture EP/DDS/polyether sulphone can be interrupted by the increase in interfacial tension, resulting in a dispersed granular morphology.51,

52

With an increasing

concentration of the thermoplastic additive, the morphology changes via a co-continuous structure into thermoplastic materials filled with thermoset particles. This phenomenon can be explained by the fact of the distinct solubility parameter (δ) values of EP/DDS and polyether sulphone molecules. Taking EP/DDS/BMI/0.25PPO as a representative example, the interaction parameter χAB can be calculated from the δ of each domain (Table 2) according to eqn. 4:54

 AB =

VAB ( A   B ) 2 RT

(4)

VAB can be calculated as follows:

VAB = (VA  VB ) 2

(5)

where T is the temperature, R is the gas constant, VA and VB are the molar volume of the A- and B- domains, respectively. Therefore, the interaction parameter χAB is an important driving force towards the microphase separation of the A- and B- domains.

18


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Table 2. Calculation of solubility parameters by component group contributions (Hoftyzer-Van Krevelen method 55) for PPO as an example Group

Number

p-phenylene –COO– –O– –CH3 =CH2 δ (MJm-3)1/2

12 6 6 30 6

Fd,i (MJm-3)1/2mol-1 1,270 390 100 420 400 δd = 15.36

Fp,i (MJm-3)1/2mol-1 110 490 400 0 0 δp = 0.74

Eh,i Jmol-1 0 7,000 3,000 0 0 δh = 5.27

Vm 3 cm mol-1 71.4 18.0 3.8 33.5 28.5

The δ can be calculated using eqn. 6:54, 55

 = ( d2   p2   h2 )

(6)

δd, δp, and δh can be calculated as follows:

d = 

Fd ,i

(7)

Vm

p=

h =

F

2 p ,i

(8)

Vm

E

h ,i

(9)

Vm

where Fd,i and Fp,i are the group contributions for dispersive and polar forces, Eh,i is the group contribution for the hydrogen bonding in the polymer, and Vm is the molar volume of the repeating unit at 298 K. The interaction parameters (χAB > 9) of PPO-rich and EP/DDS-rich domain with respect to the EP/DDS/BMI/0.25PPO system are calculated from their δ values based on component group contribution model (Table 3), implying a tendency towards phase separation for the said system.

19



Page 20 of 39

However, the driving force to form nano-fibrous structure in present research suppresses the interfacial tension force of PPO-rich domains to form granules. The competitive progression of morphologies can be rationalized by considering BMI monomer as chain extender of PPO owing to the reaction between BMI and PPO, which allows the nano-fiber length to increase. Meanwhile, the increased SIPN structures resulting from the reactions of PPO/BMI and EP/DDS play the role as a scaffold to maintain nano-fibrous structure. Figure 6D shows a clear interfacial region (circled in green) and self-assembled nano-fibers derived from the matrix of MC. Table 3. The interaction parameter χAB of different domains Domains A and B

χAB

EP/DDS, PPO EP/DDS, PPO/BMI EP/DDS/BMI, PPO

9.37 11.39 13.38

δA (MJm-3)1/2 20.27 20.27 20.96

δB (MJm-3)1/2 16.23 15.80 16.23

VA cm3mol-1 691 691 898

VB cm3mol-1 2,164 2,129 2,164

To further confirm the fiber composition, the fracture surfaces of MCs are etched with toluene for 1 h to get rid of the thermoplastic or PPO dominated phase and Figure 7 shows the SEM images of the etched fracture surfaces of EP/DDS/BMI. After toluene-treatment, plenty of micron-sized voids appear on the surfaces of EP/DDS/BMI/PPO system, and the nano-fibrous morphology completely disappear neither on the surfaces nor inside of holes, suggesting that the fibers are mainly composed of PPO dominated phase, which is soluble in toluene. This phenomenon further confirms that the formation of PPO nano-fibers are induced by cure reaction of EP/DDS/BMI/PPO and embedded as a cluster/bundle in MCs.

20


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Figure 7. SEM images of fracture surfaces of EP/DDS/BMI (A, A′), EP/DDS/BMI/0.25PPO (B, B′), EP/DDS/BMI/0.50PPO (C, C′), and EP/DDS/BMI/0.75PPO (D, D′) after etching in toluene. Thermomechanical properties and thermal stability To explore the thermomechanical properties of resultant MCs with self-assembled nano-fibrous structure, the storage modulus (E′) and tan delta as a function of temperature were studied using DMA. The E′ of EP/DDS/BMI/PPO is lower than that of the reference EP/DDS/BMI at 25 °C (Figure 8A), which is mainly attributed to the less rigid PPO-rich domains in the system and in a good agreement with the AFM analysis. EP/DDS/BMI has a Tg of 156 °C (Figure 8B) corresponding to its SIPN structure based on the BMI homopolymerization and EP/DDS reaction networks46 and shows a rapid decrease in E′ with increasing temperature between 115 °C and 160 °C. Compared to DDS monomers, the vinyl end-capped PPO can easily react with BMI, the incorporation of PPO oligomer (Mn = 2200 g·mol-1) can enlarge the average molecular weight between the crosslinking points and should lower the Tg of EP/DDS/BMI. But actually

the

Tg

values

of

EP/DDS/BMI/0.25PPO,

EP/DDS/BMI/0.50PPO

and

EP/DDS/BMI/0.75PPO are 191°C, 214 °C and 229 °C (Figure 8B), respectively, which are 35−73 °C higher than that of EP/DDS/BMI. The highest Tg of 229 °C for MCs is comparable to that of a typical high-performance thermoset BMI crosslinked with diallylbisphenol (204 °C)

21


43


using the same cure program in present work. The significantly improved Tg of EP/DDS/BMI/PPO is mainly attributed the fact that the kinked phenylene ether structure greatly restrict the segmental motion of polymer chain. In addition, the incorporation of the vinyl end-capped PPO suppresses the action between BMI and DDS that is readily to form crosslinked network with lower Tg.41 The calculated crosslinking density of EP/DDS/BMI is 1013 mol·m-3. For

EP/DDS/BMI/0.25PPO,

EP/DDS/BMI/0.50PPO,

and

EP/DDS/BMI/0.75PPO,

their

crosslinking density values are 1911 mol·m-3, 2957 mol·m-3 and 4438 mol·m-3, respectively (Table S1). The introduction of PPO can significantly increase the crosslinking density values of EP/DDS/BMI. These results are in good agreement with the analysis of Tg. 3000

1.2

(A) EP/DDS/BMI EP/DDS/BMI/0.25PPO EP/DDS/BMI/0.50PPO EP/DDS/BMI/0.75PPO

2500 2000 1500 1000

0.8

156C

0.6

191C

0.4

214C 229C

0.2

500 0

(B)

1.0

Tan Delta

Storage modulus (MPa)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 22 of 39

0.0

50

100

150

200

250

50

Temperature (C)

100

150

200

250

Temperature (C)

Figure 8. Storage modulus (A) and tan delta (B) as a function of temperature for EP/DDS/BMI and EP/DDS/BMI/PPO systems. The introduction of PPO can slightly enhance the initial thermal decomposition temperature (Tdi) at 5wt% weight loss and the temperature (Tdmax1 and Tdmax2) at maximum thermal decomposition rate of EP/DDS/BMI (Figure 9A and Table 4), indicating that EP/DDS/BMI/PPO systems exhibit better thermal stability than EP/DDS/BMI. The heat release rate (HRR) versus time and the characteristic parameters of micro-scale combustion measurement including heat release capacity (HRC), peak heat release rate (PHRR), total heat release (THR) and the 22


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temperature at peak heat release rate (TPHRR) are summarized in Figure 9B and Table 4. Although PPO do not significantly influence the THR and TPHRR values, it can decrease the HRC and PHRR values of EP/DDS/BMI, and as compared to EP/DDS/BMI, EP/DDS/BMI/PPO systems display 23−32% and 23−33% decreases in HRC and PHRR, respectively, indicating that EP/DDS/BMI/PPO systems have better flame retardancy than EP/DDS/BMI. In addition, it should be noted that unlike EP/DDS/BMI, all EP/DDS/BMI/PPO systems exhibit double-peak HRR curves rather than a single peak HRR curve (Figure 9B), and they also show double-peak derivative thermogravimetric (DTG) curve that is different from that of EP/DDS/BMI (Figure 9A), indicating that the two-step flame retarding mechanism of EP/DDS/BMI/PPO differs from that of EP/DDS/BMI. To further investigate the flame retarding mechanism of EP/DDS/BMI/PPO, TG-IR technology was used to detect gaseous products during the decomposition process. Figures 9C-9F show the three-dimensional FTIR spectra of the gaseous products evolved from the combustion processes of EP/DDS/BMI and EP/DDS/BMI/PPO systems. The samples with the same size and weight are taken in TG-IR tests, so the intensity of the absorption peak in the three-dimensional spectra can reflect the amounts of decomposition products. The characteristic absorption peaks of -OH at 3500 cm-1 and =C−H group at 3050 cm-1 are found in the FTIR spectra of all samples. Although the Tdi of EP/DDS/BMI is as high as 389 °C, an absorption peak of CO2 at 2362 cm-1 appears in the FTIR spectrum of EP/DDS/BMI (Figure 9C) when the temperature rises to 200 °C. The introduction of PPO into EP/DDS/BMI can shift the absorption peak of CO2 to high temperature. For EP/DDS/BMI/0.50PPO and EP/DDS/BMI/0.75PPO, the absorption peak of CO2 can be obviously detected at 420 °C (Figures 9D − 9F), and the intensities of the absorption peaks of various hydrocarbons and CO2 are significantly weaker than

23



that of EP/DDS/BMI, suggesting that the introduction of PPO can suppress the decomposition of polymer matrix and EP/DDS/BMI/PPO systems show better flame retardance than

100

(A)

EP/DDS/BMI EP/DDS/BMI/0.25PPO EP/DDS/BMI/0.50PPO 1 EP/DDS/BMI/0.75PPO

Weight (%)

80 60

0

40 20

-1

0

100

200

300

400

500

600

Temperature (C)

700

800

Heat release rate (W/g)

EP/DDS/BMI.

Deriv. weight (wt%/C)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 24 of 39

400

(B)


300 200 100 0 0

100

200

300

400

500

600

700

Time (s)

Figure 9. Heat release rate (A) and thermogravimetric analyses (B) as a function of temperature, and three-dimensional FTIR spectra (C-F) of the gaseous products evolved from the combustion processes of EP/DDS/BMI and EP/DDS/BMI/PPO systems.

24


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Table 4. Thermal properties and micro-scale combustion data of EP/DDS/BMI and self-reinforced MCs (EP/DDS/BMI/PPO).

Sample

Tdi Tdmax1/Tdmax2 (°C) (°C)


389 393 391 395

HRC [J·(g·K)-1 ]

410/0 414/456 415/455 412/456

298 229 221 203

PHRR THR TPHRR -1 (W·g ) (kJ·g-1) (°C) 294.7 226.8 218.4 197.0

16.9 18.5 17.9 18.7

416 412 409 409

The improved thermal stability and flame retardance of EP/DDS/BMI/PPO systems can be mainly explained as follows: First, PPO loses weight over 450 °C (Figure S3) for the scission of ether links, accompanying the rearrangement reactions through methylene bridges,56,57 therefore PPO has better thermal stability than EP/DDS/BMI, which can make PPO act as a thermal barrier to protect the matrix and dissipate heat, and then inhibit the degradation of matrix. Second, the in-situ fibers can effectively absorb volatiles and reduce the diffusion of combustible volatiles for their nano-structures. Third, the excellent interface interaction at the in-situ PPO fibers/matrix can improve the thermal stability and the flame retardance of EP/DDS/BMI/PPO. Although PPO can stabilize the thermal stability of EP/DDS/BMI and suppress the degradation of matrix, the pyrolysis of PPO starts when the temperature rises to 450 °C, then degradation products including xylenol, phenol, methylphenol, dimethylphenol, etc. occur.56,57 As a result, EP/DDS/BMI/PPO systems undergo a two-stage degradation as indicated by the double peaks appeared in DTG and HRR curves (Figures 9A and 9B). Mechanical properties Figure 10 shows the flexural strength and fracture toughness of EP/DDS/BMI and EP/DDS/BMI/PPO

systems.

The

flexural

strengths

of

EP/DDS/BMI/0.25PPO,

EP/DDS/BMI/0.5PPO and EP/DDS/BMI/0.75PPO are 138 MPa, 149 MPa, and 153 MPa, 25



respectively, which increase by 30 %, 40 % and 44 %, respectively, as compared to EP/DDS/BMI (Figure 10A). The fracture toughness (KIC) of EP/DDS/BMI/PPO shows a similar increasing trend to the flexural strength as the PPO content increases. The KIC values of EP/DDS/BMI/0.25PPO, EP/DDS/BMI/0.5PPO and EP/DDS/BMI/0.75PPO are 1.65 MPa·m1/2, 1.74 MPa·m1/2 and 1.79 MPa·m1/2, respectively (Figure 10B), which are 50 %, 58 % and 63 % higher than those of EP/DDS/BMI. It is worth noting that the significant enhancement of both flexural strength and toughness can be obtained with a small loading of PPO relative to the whole resin system as indicated in Table 1. 200

2.4

(A)

0.4 0.0

EP/DDS/BMI/0.75PPO

0.8

EP/DDS/BMI/0.50PPO

1.2

EP/DDS/BMI/0.25PPO

1.6

EP/DDS/BMI

0

EP/DDS/BMI/0.75PPO

40

EP/DDS/BMI/0.50PPO

120

EP/DDS/BMI/0.25PPO

1/2

KIC (MPam )

160

80

(B)

2.0

EP/DDS/BMI

Flexural strength (MPa)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 26 of 39

Figure 10. Flexural strength (A) and fracture toughness (B) of EP/DDS/BMI and EP/DDS/BMI/PPO. The increased flexural strength and toughness of materials can be mainly attributed to the presence of in-situ PPO nano-fibrous structure. Specifically, the PPO nano-fibrous domains can bring the crack pinning and blunting behaviors and may fracture during the load process, which can cause the efficient energy dissipation to toughen the matrix,58, 59 and the PPO nano-fibers have higher aspect ratio and exhibit anisotropic properties, which can reinforce the matrix and improve the strength of matrix. In addition, the in-situ nano-fibers mainly consisting of PPO-rich polymers prefer to disperse well within the matrix and the chemical interaction at PPO/matrix interface can effectively enhance the mechanical property of EP/DDS/BMI/PPO.

26


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The SEM images demonstrate that EP/DDS/BMI/PPO systems shows much rougher fracture surfaces than EP/DDS/BMI, the crack pinning and blunting phenomena around nano-fibrous clusters and the fracture of PPO nano-fiber domains are clearly observed on the fracture surfaces of EP/DDS/BMI/PPO system (Figures 1B-1D), which indicate that EP/DDS/BMI/PPO system have better toughness than EP/DDS/BMI. The bonding at PPO/matrix interface can confirm the strong interface interaction (Figures 1′B-10D′), which prefers to improve the mechanical property of EP/DDS/BMI/PPO system. Figure 11 shows the representative stress-strain diagrams of EP/DDS/BMI and EP/DDS/BMI/0.75PPO measured by the strain gauge and the digital image correlation (DIC) technique. The EP/DDS/BMI specimen show tensile properties of typical thermosets, i.e., the tensile strength of ~30MPa and elongation at break of ~1 %. The tensile strength and elongation at break of EP/DDS/BMI/0.75PPO increase by 30 % and 55 %, respectively, as compared to EP/DDS/BMI. The application of DIC allows us to investigate strain concentrations and characteristic damage morphology in terms of crack initiation and propagation on the specimen surface, and to assess their impact on the homogenized stress-strain response.60-65 The homogenous vertical strain (ɛy) and perpendicular strain (ɛx) distribution in EP/DDS/BMI during tensile test are presented in Figures 11B and 11C, a higher strain rate (0.017 mm/s) is observed at 0.7 % (B point in Figure 11A) close to the middle of the specimen where the crack initiated and the failure event occurred (Figure 11C). Despite the strain concentration of ɛy has been observed in the lower part of EP/DDS/BMI/0.75PPO specimen (Figure 11D), the strain rate clearly shows a top-down distribution, indicating that the load transfer in axial direction is not disrupted by the presence of nano-fibrous structure. Upon close inspection of the ɛx of EP/DDS/BMI/0.75PPO, of which the distribution is similar to that of traditional fiber reinforced polymeric composites.61

27



This again confirms that the formation of an in-situ nano-fibers reinforced MC based on EP/DDS/BMI/PPO. The improvement in tensile property illustrates that the load transfer efficiency in EP/DDS/BMI/PPO is better than that in EP/DDS/BMI.

Figure 11. Representative stress-strain curves (A) and strain distribution in EP/DDS/BMI (B and C) and EP/DDS/BMI/0.75PPO (D and E) specimens, specified time intervals during a test indicated by the symbols and marked as B, C, D and E in the stress-strain curves. Vertical strain (ɛy, parallel to load), perpendicular strain (ɛx, perpendicular to load), and strain rate are shown for each specified time interval.

28


Page 28 of 39

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4. CONCLUSIONS In summary, we have demonstrated that it is possible to prepare in-situ reinforced molecular composite (MC) with non-equilibrium self-assembled nano-fibrous structures induced by cure reaction. The MC is based on bisphenol A epoxy resin (EP), 4,4′-diaminodiphenylsulfone (DDS), bismaleimide (BMI), and polyphenylene ether (PPO) oligomer. The formation of in-situ aligned nano-fibrils arises from the fact that the polymerization force of SIPN thermosetting polymers extrudes PPO phase through the interspaces of SIPN. As compared to the control EP/DDS/BMI sample, EP/DDS/BMI/PPO MC shows improved thermal stability and flame retardance mainly owing to the thermal barrier of PPO with high thermal property, the volatile absorbing function of nano-structure PPO fiber and the chemical interaction at PPO/matrix. Especially, the Tg of EP/DDS/BMI can significantly increase by 35−73 °C by incorporating PPO, because the kinked phenylene ether structure greatly restricts the segmental motion of polymer chain. The mechanical property of MC significantly increases mainly owing to the toughening and reinforcing functions of the in-situ aligned PPO nano-fibers, and the obtained EP/DDS/BMI/0.75PPO shows a 44 % increase in flexural strength, a 63 % increase in fracture toughness, a 30 % increase in tensile strength and a 55% increase in elongation at break, respectively. Herein the structure development in EP/DDS/BMI/PPO is completely different from that of traditional thermoplastic/thermoset MCs where the spherical thermoplastic particle phase is readily to be separated during the polymerization process of thermoset and dispersed in the thermoset matrix. To the best of our knowledge, the approach described here is the first non-equilibrium self-assembled nano-fibrous structure induced by cure reaction, thus leading to a much richer structural diversity compared to equilibrium-processed materials. The present

29



Page 30 of 39

findings reveal the new understanding of the morphological development of self-assembled nano-fibrous structure from thermoplastic/thermoset systems. ASSOCIATED CONTENT Supporting Information Additional DMA curves of EP/DDS/BMI, EP/DDS and BMI/DDS (Figure S1), SEM images of EP/DDS/BMI/0.25PPO (Figure S2), TGA curve of PPO (Figure S3) and the crosslinking density of EP/DDS, EP/DDS/BMI and EP/DDS/BMI/PPO systems (Table S1). AUTHOR INFORMATION Corresponding Author *[email protected] (L.Y.) Notes The authors declare no competing financial interests. ACKNOWLEDGMENT The authors thank the National Natural Science Foundation of China (No. 51273135 and 51703148), and the support of the Priority Academic Program Development of Jiangsu Higher Education Institutions (PAPD), State and Local Joint Engineering Laboratory for Novel Functional

Polymeric

Materials,

and

the

China

Postdoctoral

Science

Foundation

(2017M611901). We acknowledge helpful discussion with Prof. Stephen J. Picken. REFERENCES (1) Lefaux, C. J.; Kim, B. S.; Venkat, N.; Mather, P. T. Molecular Composite Coatings on

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