Shape memory and self-healing materials from supramolecular block polymers

Accepted Manuscript Shape memory and self-healing materials from supramolecular block polymers Jiuyang Zhang, Mengmeng Huo, Min Li, Tuoqi Li, Naixu Li, Jiancheng Zhou, Jing Jiang PII:

S0032-3861(17)31111-4

DOI:

10.1016/j.polymer.2017.11.043

Reference:

JPOL 20156

To appear in:

Polymer

Received Date: 12 August 2017 Revised Date:

25 October 2017

Accepted Date: 19 November 2017

Please cite this article as: Zhang J, Huo M, Li M, Li T, Li N, Zhou J, Jiang J, Shape memory and self-healing materials from supramolecular block polymers, Polymer (2017), doi: 10.1016/ j.polymer.2017.11.043. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

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Shape Memory and Self-healing Materials from Supramolecular Block Polymers Jiuyang Zhang,a,* Mengmeng Huo,a Min Li,a Tuoqi Li,b Naixu Li,a Jiancheng Zhou,a,* Jing Jiang c a

School of Chemistry and Chemical Engineering, Southeast University, Nanjing, 211189,

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China b

The Dow Chemical Company, 2301 N. Brazosport Blvd, B-1608, Freeport, TX USA 77541

c

School of Chemistry and Chemical Engineering, Nanjing University, Nanjing 210093, China

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E-mail: [email protected]; [email protected]

Abstract: Shape memory thermoplastics are an important class of engineering polymers with characteristics.

A

series

of

diblock

polymers,

polymethylmethacrylate-b-

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smart

poly(butylacrylate-co-2-acrylamido-2-methyl-1-propanesulfonic acid) (PMMA-b-P(BA-coAMPS)) with varying molecular weights and compositions were conveniently synthesized. The monomer selection as well as the design of molecular architecture facilitated microphase

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separation in the prepared polymers, leading to two networks formed by the vitrified PMMA domain and the supramolecular interactions from the AMPS component, respectively. As a result, these block polymers exhibited excellent shape memory property (shape recovery ratio:

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95%). Besides, the materials were also mechanically tough with an outstanding breaking strain of around 500% and a high tensile strength over 10 MPa. Furthermore, these shape

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memory polymers (SMPs) were built merely relying on physical interactions, retaining the thermal processability of engineering thermoplastics: it was facile to melt compound the materials into various desired shapes (dog-bone, disk, and film, etc.). More importantly, the prepared SMPs also showed exceptional self-healing capability under ambient conditions, benefited from the supramolecular interactions (ionic interactions) within the polymer network. This work elucidates a new path towards fabricating multifunctional low-cost shapememory materials for a host of industrial applications.

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1. Introduction The continuously growing demand for ‘smart materials’ in technological devices has greatly driven the development of stimuli-responsive materials.[1-6] Shape memory polymers (SMPs) are an important class of stimuli-responsive materials capable of retaining the permanent

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shape via transforming from any temporary shape under proper conditions.[7-12] SMPs have aroused a wide spread attention in the past decade, not only in fundamental researches but also in a plethora of ‘smart’ applications, such as ‘smart’ textiles,[7, 8]

biomedical

. Although SMPs have been

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materials[14] and self-repairing industrial components[6,

8, 13]

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extensively explored, there are still several stringent bottlenecks in restricting their broad industrial applications.[7, 8]

The first hurdle lies in the processability of SMPs.[7, 8] The crosslinking techniques utilized to build SMPs fall into two categories, i.e., chemical and physical crosslinking.[7] The majority of current literature reports focus on the former,[7, 8, 14, 15] due to the excellent mechanical

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performance and chemical stability imparted by the chemical crosslinkers. However, this permanent crosslinking inevitably sacrifices the solvent and thermal processability of the

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resultant materials, rendering them unsuitable for industrial scale production.[11, 15-17] On the other hand, physically crosslinked SMPs are generally attained via phase separation (such as 19]

) and crystallization of one or more phases (for

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polymethyl methacrylate vitrification[18,

example: polyethyl oxide crystallization[18]), allowing the products to be thermally processable.[7, 11] Yet one disadvantage of these physical networks is the intrinsic deficiency in mechanical performance, which detracts from many applications where mechanical robustness is essential.[7, 11, 20] Furthermore, most physically crosslinked SMPs are limited to such building components as semicrystalline or supramolecular polymers.[10,

20-23]

Consequently, so far, only a few commercially viable polymers have been used as physically crosslinked SMPs with desired features, for instance, Nafion,[12] polyurethanes[7,

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8, 22]

and

ACCEPTED MANUSCRIPT polyether ether ketones[24]. To fabricate new SMPs endowed with processability as well as outstanding mechanical performance warrants further studies. Another challenge confronting SMPs is to achieve multi-functionality in addition to the shape memory characteristic.[7,

8, 11]

Some pioneering work has successfully conjugated other

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functionalities with SMPs, including self-healing capability,[25-29] biocompability[30-32] and switchable optical properties[30, 33]. Among these smart features, self-healing gains particular interests,[7, 8, 27] and has been realized via the utilization of reversible covalent bonds[27, 34, 35] 26, 36]

Unfortunately, the

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or constructing vascular systems within polymeric matrices.[25,

delicate molecular design and preparation of special chemical reagents are usually time and

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cost consuming. Therefore, it is preferred to develop a facile and efficient approach at manageable cost to granting SMPs the self-healing property.

The third deficiency of conventional SMPs is the limited exploration of polymer architectures. Specifically, the selection of building polymers is currently confined to homo- or random-

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polymers, leaving a large uncharted territory of other various architectures, like block, star and graft polymers, etc.[7, 11, 17, 30] These complex architectures provide a remarkably wide space of structural parameters[37] for designing and tailoring the microstructure and

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consequent properties of SMPs. Until recently, there are a few reports on utilizing poly(ʟlactide) (PLLA) or poly(ε-caprolactone) (PCL) based multiblock copolymers[38-40] or PCL 42]

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based star polymers[41,

to construct shape memory materials. In several studies, the

researchers were able to achieve excellent mechanical performance (e.g. tensile strength > 10MPa, breaking strain > 800%) via judiciously controlling the semicrystalline block length and the equivalent repeating number of connected triblocks in multiblock copolymers.[39, 43] However, the behaviour of these SMPs is dominated by the crystallization and melting of PLLA or PCL segments, making it difficult to optimize the shape-memory and mechanical properties simultaneously. A feasible solution to this dilemma may require appropriately using the diverse architectural parameters of block polymers along with probing other shape -3-

ACCEPTED MANUSCRIPT memory mechanisms. Although various diblock copolymers have been applied in diverse aspects, including self-healing,[44,

45]

lithography[37,

46]

and biomedical applications,[47] the

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utilization of block polymers in SMPs is still far less investigated.

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Scheme 1. Illustration of the self-healable shape memory thermoplastic material fabricated from supramolecular block polymers. The shape memory effect (left part) is realized by vitrifying and liquefying the microphase separated PMMA domain. The self-healing property is achieved relying on the high chain mobility of the rubbery phase as well as the “rebonding” of supramolecular interactions (ionic interactions).

In this paper, we introduced a new type of self-healable shape-memory materials to tackle the aforementioned aspects limiting the application of conventional SMPs. A series of diblock polymethylmethacrylate-b-poly(butylacrylate-co-2-acrylamido-2-methyl-1-

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polymers,

propanesulfonic acid) (PMMA-b-P(BA-co-AMPS)) with varying molecular weights and

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compositions were conveniently synthesized via the reversible addition−fragmentation chaintransfer (RAFT) polymerization. The starting materials are low-cost monomers (MMA, BA and AMPS) for engineering polymers. The selection of monomers and the design of molecular architecture lead to a strong Flory-Huggins interaction parameter (χ) between PMMA and P(BA-co-AMPS) segments, facilitating microphase separation for the materials, and therefore, segregated PMMA domains upon vitrification work as physical crosslinkers. Meanwhile, AMPS units are broadly applied in polymeric materials for their supramolecular interactions (dipole-dipole/ionic interactions)[48-50] (Scheme 1), which could offer the second -4-

ACCEPTED MANUSCRIPT crosslinkers. As such, the prepared PMMA-b-P(BA-co-AMPS) diblocks behave like typical SMPs with excellent shape memory characteristics: shape recovery ratio 95%. Notably, these materials are also endowed with superior mechanical properties, exhibiting breaking strain and tensile strength at about 500% and 10 MPa, respectively. Furthermore, these

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supramolecular block polymers are thermally processable due to the physical crosslinking and are facile to be melt compounded into any desired shape (dog-bone, disk, and film). More importantly, the prepared SMPs also possess the self-healing capability under ambient

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conditions, benefited from the ionic interactions within the polymer network. To the authors’ knowledge, it is the first report on mechanically tough, self-healable shape-memory

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thermoplastics built with supramolecular block polymers. We believe this work provides a great platform for designing and fabricating multifunctional and mechanically robust SMPs that are ready to be applied in a myriad of practical services.

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2. Results and Discussion

Diblock polymers PMMA-b-P(BA-co-AMPS) were prepared via RAFT polymerization using commercially viable monomers. PMMA was synthesized as the first block and macroinitiator

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for the later polymerization (Figure 1A). The molecular weight of PMMA was obtained by

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end group analysis of 1H NMR spectra (Figure S1 - S2), according to the ratio of integrated areas under peaks of phenyl protons (7.0 - 7.5 ppm) in the RAFT agent and peaks of methyl protons in methacrylate (3.5 ppm). The narrow polydispersity (Ð) of the resultant PMMA macroinitiators obtained from gel permission chromatograph (GPC) indicated the excellent control over the polymerization (Table 1 and Figure S3). The second block was formed by copolymerization of BA and AMPS monomers. Here, AMPS monomer was chosen not only due to its unique intermolecular interactions, but also its excellent solubility in common organic solvents (its sodium salts, AMPS-Na, was less soluble in organic solvent, like DMF and acetone). The molecular weight of the second block and the molar ratio of MMA, BA and -5-

ACCEPTED MANUSCRIPT AMPS in polymers were also characterized via 1H NMR spectra as shown in Figure 1B and Figure S4 - S6 (Detailed analysis of 1H NMR peaks was shown in Figure S4). Unfortunately, the polydispersity of the diblock polymers were not able to be analyzed by GPC due to the strong electrostatic interaction between the polymer specimens and columns.[51] To

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demonstrate the controlled polymerization of the second block, a kinetic study over the chainextension was performed, and results were plotted in Figure 1C. The linear fitting for both BA and AMPS revealed a controlled polymerization of both monomers. A series of PMMA-

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b-P(BA-co-AMPS) diblock polymers with varying molecular weights and compositions were synthesized and the molecular characteristics have been summarized in Table 1. For

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nomenclature, Mx-b-ByPz is to denote the diblocks, with x, y and z referring to the number of repeat units of MMA, BA and AMPS, respectively. Meanwhile, a random polymer Mx-co-Byco-Pz was also synthesized as the control and its molecular information was shown in Table 1

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and Figure S7.

Figure 1. (a) Synthetic scheme for Mx-b-ByPz block polymers; (B) 1H NMR spectrum for the diblock polymer: M320-b-B1080P170; (C) Kinetic studies: semilogarithmic plots of chain extension to prepare diblock copolymers (M320-b-B1080P170 is the example shown above) using PMMA (320 kDa) as the macroinitiator. Solid lines are the linear fitting for the experimental

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ACCEPTED MANUSCRIPT data points. For nomenclature, Mx-b-ByPz is to denote the diblocks, with x, y and z referring to the number of repeat units of MMA, BA and AMPS, respectively.

The thermal properties of prepared polymers were characterized by differential scanning calorimetry (DSC). As evident in Figure 2A and S9, there are two different glass transitions

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(Tg) for all diblock polymers, indicating microphase separation between the PMMA and P(BA-co-AMPS) domains in the materials, and the associated Tg values have been listed in Table 1. Compared with the Tg of PBA homopolymer (-60°C),[52, 53] incorporation of the

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AMPS units raises the Tg of P(BA-co-AMPS) domain to the range from -25 to 50 °C, depending on the content of AMPS. This increase is likely due to the presence of

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supramolecular interactions (ionic interactions), associated with the AMPS segments.[48-50] The morphology of the diblock polymers resulted from microphase separation is determined by the overall degree of polymerization (N), the composition f (fMMA = NMMA/N) and the Flory-Huggins interaction parameter χMMA/BxPy. According to a rough estimation based on the

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solubility parameters of three monomers (see detailed calculation in the supporting information), all diblocks have a high χN value (>40) at 25°C. Meanwhile, the fMMA value is ranging from 0.20 to 0.29 for these polymers. Given the phase diagram of typical linear AB

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diblock polymers,[54] the diblocks in this study are likely to form morphologies with spherical or cylindrical PMMA domains packed in the P(BA-co-AMPS) matrix. To further investigate

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the microstructure generated by microphase separation, small angle x-ray scattering (SAXS) was utilized (Figure 2B). However, the higher order peaks are barely discernible on the 1D SAXS profiles, making it difficult to draw any explicit conclusion on the morphology of the diblocks. It reveals the absence of long-range ordering in these materials. On the other hand, the domain spacing (D) is strongly dependent on the AMPS content. For instance, M320-bB1080P170 has a greater D at 53 nm than M320-b-B1050P50 that contains a lower fraction of AMPS and the D is at 44 nm. The latter diblock has a shorter P(BA-co-AMPS) block length and also a lower χMMA/BxPy. The results were confirmed by atomic force microscopy images -7-

ACCEPTED MANUSCRIPT (Figure S8). Both M320-b-B1050P50 and M320-b-B1080P170 were microphase separated into disordered structures. In contrast, the control random polymer M310-co-B1030-co-P155 shows a homogenous phase and a single glass transition (Figure 2), because random polymers lack the

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capability to phase separate.

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Figure 2. (A) Differential scanning calorimetry (DSC) curves of Mx-b-ByPz block polymers and M310-co-B1030-co-P155 random polymers. Arrows indicate the glass transition temperatures. (B) 1-D SAXS profiles of melt-compounded random polymer M310-co-B1030-co-P155, block polymers M320-b-B1080P170 and M320-b-B1050P50. Arrows denote the position of primary peaks. Table 1. Molecular characteristics and thermal properties of PMMA homopolymers, Mx-bByPz block polymers and random copolymers. For nomenclature, Mx-b-ByPz is to denote the diblocks, with x, y and z referring to the number of repeat units of MMA, BA and AMPS, respectively.

a

(kg/mol 32 52 205 213 178 228 196

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PMMA-32k PMMA-52k M320-b-B1080P170 M320-b-B1030P233 M320-b-B1050P50 M520-b-B1110P160 M310-co-B1030-co-P155

Mn

Ðb

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Samples

1.2 1.4 -f -f -f -f -f

FMMA c 1 1 0.20 0.20 0.22 0.30 0.21

FAMPS c 0.11 0.15 0.04 0.09 0.10

Tg (oC) d Tg, ByPz e -12.2 53 -23.6 -19.1

Tg, M e 101 97 103 120 h 112 107 Tg : 2.9 g

Mn determined by 1H NMR end group analysis. b Ð was determined by GPC at 40 °C using polystyrene standards and tetrahydrofuran (THF) as the mobile phase. c FMMA and FAMPS are the molar fractions of MMA and AMPS units in associated polymer chains, respectively, calculated via of the integrated areas under MMA and AMPS peaks in 1H NMR spectra. d Glass transition temperature (Tg) was obtained from the second heating trace of DSC measurements. e Tg, ByPz and Tg, M are the glass transition temperatures for P(BA-co-AMPS) phase and PMMA phase, respectively. f Block polymers are not able to be characterized by GPC due to the ionic interaction with columns. g Glass transition temperature for random polymer M310-co-B1030-co-P155. h A high Tg value for PMMA domain was observed, possibly due to the restricted chain mobility resulted from the very high content of AMPS.

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ACCEPTED MANUSCRIPT The microphase separated PMMA domains upon vitrification act as the physical crosslinkers for the polymeric continuum. Simultaneously, the supramolecular interaction contributed by the AMPS segments further reinforces the matrix. The mechanical properties of all prepared materials were characterized using uniaxial tensile tests under ambient conditions. Dog-bone

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specimens were prepared via melt-compounding at 130 °C, below the degradation temperature of 200 °C (obtained from thermogravimetric analysis, Figure S10). Representative engineering stress versus strain curves are plotted in Figure 3A. Average

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values of the mechanical properties (determined from at least three independent measurements over identical specimens) including elastic modulus (E), strain at break (εb),

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tensile strength (σb), yield strength (σy) and tensile toughness are summarized in Table S1. Most of block polymers show excellent tensile properties (εb at 250% – 800%, σb at 0.5 – 12 MPa and tensile toughness 19 – 29 MJ/m3), significantly outperforming the control random polymer that has a low toughness of 6.9 MJ/m3. It is due to the fact that M310-co-B1030-co-P155,

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unlike the diblock counterparts, is incapable of phasing separating, leading to MMA and AMPS single units or short segments with much smaller sizes than those microphase separated domains in diblocks (Figure 2B). Consequently, in the random polymer, these

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MMA and AMPS components are not effective physical crosslinkers to reinforce the polymeric continuum, rendering poor mechanical performance. The mechanical behavior of

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diblock polymers is also strongly composition dependent, especially on the AMPS content. Drastically enhanced pulling stress level as well as σy and σb can be observed with increasing the fraction of AMPS from M320-b-B1050P50 to M320-b-B1080P170 and further to M320-bB1030P233. Additionally, the content of glassy PMMA domains in the materials is also crucial to the tensile properties. Figure 3A indicates that a longer PMMA block (M520-b-B1110P160) leads to a higher σb (11 MPa) but a relatively lower εb (210%) compared with the analogue M320-b-B1080P170. The diverse structural parameters of this molecular architecture offer a wide

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conditions.

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Figure 3. (A) Representative engineering stress versus strain curves for different Mx-b-ByPz block polymers and M310-co-B1030-co-P155 random polymers; (B) Shape-memory programming for M320-b-B1080P170; εm and εu represent the strain upon the loading at 25 °C and retrieving of external stress at 80 °C, respectively. (C) Photos of the shape recovery process of M320-b-B1080P170 at 80 °C. Photos were taken after different time intervals upon reaching 80°C. Scale bar: 1cm.

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The ionic interactions of AMPS units in polymers not only reinforced the materials, but also formed a stable network for shape-memory. The two networks generated by the PMMA vitrification and ionic interactions of AMPS units impart great shape-memory property to these polymers. Dynamic mechanical analysis (DMA) was utilized to characterize the shape memory behavior of the materials, as shown in Figure 3B for the diblock M320-b-B1080P170. After heating the specimen to 80 °C, a static loading of 0.03 MPa was applied at about 29 min to uniaxially stretch the specimen, causing an immediate strain εm of ~ 30%. This stretched state was maintained while cooling the specimen to 25 °C. And then the temperature was brought back to 80 °C and the stress was removed for recovery. The strain of the specimen - 10 -

ACCEPTED MANUSCRIPT dropped to εr (~ 1.5%) after 80 minutes heating. A high shape fixity ratio at 98% and a shape recovery ratio of 95% were shown in Figure 3B (See supporting information, Part 2, for the definition and calculation). Such shape fixing and recovery ratios originate from the ionic interactions[48, 50] associated with the AMPS segments. These supramolecular interactions are

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not affected by elevated temperature up to 140 °C or even higher (see more details below), forming a stable network to ensure the fix of permanent shapes. Meanwhile, the liquefaction above 80 °C and vitrification below 80 °C of the PMMA domains make the shape

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programming possible. It is also noted that the order-disorder transition temperatures (TODT) of these diblocks are fairly high (see the estimation in supporting information), which means

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the shape transformation is unlikely to occur in a disordered state where the inter-mixing of PMMA and P(BA-co-AMPS) domains may entirely disrupt the two network structures. Therefore, these diblock polymers have another merit of a wide temperature window for shape memory. Figure 3C gives the visualization of how specimens transforming in the

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recovery process (M320-b-B1080P170 as the example). The original specimen from meltcompounding was shape-fixed as a flat strip at 80 °C (close to Tg, PMMA) and then allowed to cool to adopt a temporary shape as a helix (same as the 0s photo in Figure 3C). After

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reheating to 80 °C, the helix was capable to recover to the original shape within 30 seconds, indicating the outstanding shape memory capability. The same phenomenon was observed in

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Figure S14. Samples kept their shape after fixing and could recover after a second heating. The discussion was also concluded in the revised manuscript and SI.

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Figure 4. (A) Representative photos of the M320-b-B1080P170 dog-bone specimen for the selfhealing test. Initially, a dog-bone specimen was completely cut into two pieces, and one piece was stained by methylene blue to show the blue color. The two pieces were allowed to stay in contact and self-heal under ambient conditions for 3 hours before stretching. Scale bar: 1 cm. Representative engineering stress versus strain curves for self-healing tensile specimens of (B) M320-b-B1080P170 and (C) M320-b-B1030P233 after different healing time. The presence of rubbery BA components in the diblock polymers not only grants superior

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ductility, but also makes these materials self-healable, realizing multi-functionality of SMPs. Figure 4 is a clear demonstration of the self-healing behavior of the Mx-b-ByPz diblocks.

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Dog-bone specimens were initially cut into two separate pieces and then both pieces (one piece stained with methylene blue for better visualization) were placed in contact with each other for self-healing under ambient conditions. After 3 hours, the two pieces were successfully healed into one complete dog-bone sample and it was stretchable as indicated in Figure 4A. Figure 4B and 4C present the tensile curve (σ vs ε) comparison between the intact and self-healed specimens for two representative diblock polymers. Specimens were cut to the depth of 70% – 90% of the initial thickness and left at room temperature in air for healing. It is evident that the self-healed M320-b-B1080P170 specimens were able to recover

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ACCEPTED MANUSCRIPT ~70% of all tensile properties (ε, σ and toughness) within only one hour and to reach near 90% within three hours of self-healing (see Table S1). Such decent self-healing property, on one hand, is benefited from the low Tg of P(BA-co-AMPS) domains, which gives these chains/segments high mobility at room temperature to re-entangle across the broken interface,

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leading to quick self-healing. Unfortunately, different from full recovery elastomers, repeating of the self-healing measurement was not able to conduct due to the permanent plastic deformation of the stretched samples. Without AMPS units, the M320-b-B1120 is extremely

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mechanically weak. The damaged M320-b-B1120 could be healed via chain entanglement but the interface was very weak (Figure S13). Meanwhile, compared with samples (M320-b-B1120)

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without AMPS, M320-b-B1050P50 with only 50 units of AMPS showed much improved selfhealing behavior (Figure S13): around 91% in toughness was recovered, far more than it of M320-b-B1120 (12%). The AMPS segments effectively help the mobile chains to rebuild strong ionic interactions with other segments in the vicinity, resulting in exceptional recovery and

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mechanical strength. This conclusion could be also supported by the behaviour of M320-bB1030P170 (Tg = −12 °C) and M320-b-B1030P233 (Tg = 50 °C) as shown in Figure 4C. Samples with AMPS units are capable to be self-healed to recover their original mechanical strength

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and toughness. M320-b-B1030P233 is also capable to self-heal under ambient conditions. However, compared with M320-b-B1080P170, the higher AMPS content results in a slower

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recovery rate (longer than 6h) and lower recovery ratio in toughness (60%) (Healing efficiency was summarized in Table S1). This is attributed to the confined chain-mobility (high Tg at 50 °C) of P(BA-co-AMPS) for M320-b-B1030P233. It is interesting that the healed M320-b-B1030P233 sample reached the maximum strain after 3 hours instead of 6 hours. It may be ascribed to the little moisture absorbed by ionic AMPS during the long time healing. Further researches are required in future to explore the detailed reasons of this phenomenon.

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Figure 5. (A) Dynamic mechanical analysis (temperature sweep) for block polymers M320-bB1080P170 and M320-b-B1050P50, and random polymer control M310-co-B1030-co-P155. Open and closed symbols are for storage modulus (G') and loss modulus (G''), respectively. The dash line denotes a common transition for samples occurring at 140 °C. Arrows indicate the drop of G' and G''. (B) Complex viscosities of these materials at 170 °C; (C) Photos of various melt-compounded geometries (M320-b-B1080P170): dog-bone specimen, film (picked by a white stick) and disk. Scale bar: 1 cm.

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Unlike chemically-crosslinked SMPs being unprocessable, these Mx-b-ByPz diblock polymers still retain the thermal processability as conventional engineering thermoplastics, and the presence of ionic interactions is not to the detriment of thermal processing, as also observed in other block polymer systems.[55-57] DMA was applied to fully investigate the transitions occurring with increasing temperature from -30°C to 180°C. Figure 5A reveals a general trend that the storage moduli (G') of block or random polymers decreased as the temperature rises. According to the tan δ (= G''/G') vs temperature curves of these materials (Figure S12), the random polymer M310-co-B1030-co-P155, and block polymers M320-b-B1080P170 and M320-b- 14 -

ACCEPTED MANUSCRIPT B1050P50 showed glass transitions at 1.7 °C, -21 °C and -29 °C, respectively, consistent with the Tg, BxPy values obtained from DSC measurements. However, for block polymers, the other glass transition for the PMMA domain was not observed, which may be due to the existence of ionic interactions in the matrix. Specifically, plateaus of G' exhibited at the intermediate

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temperature range (60°C – 140°C), ascribing to the network structure from the supramolecular interactions, which was similar to previous reports on sulfonic groupcontaining block polymer networks.[55-57] This strong network mitigated the decreasing effect

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on G' associated with the PMMA glass transition. The plateau value for block polymers was higher than that of the random polymer because of the microphase separation. Besides,

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compared with M320-b-B1050P50, a higher content of AMPS as in M320-b-B1080P170 also led to a greater G' plateau value and caused such plateau to extend to higher temperatures (> 190°C). These results suggest that the microphase segregated P(BA-co-AMPS) domains and more AMPS segments contribute to stronger supramolecular interactions. Interestingly, both the

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random polymer M310-co-B1030-co-P155 and block polymer M320-b-B1050P50 showed a clear decrease in G' as well as in G'' at around 140 °C. It possibly corresponds to the dissociation of AMPS-involved supramolecular interactions. The dissociation temperature of ionic

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interaction in sulfonic acid group-containing polymers has been reported in the range of 140 – 180 °C.[55-58] Figure 5B and S12 present the frequency sweeps at high temperatures (170 °C),

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evidently showing that under relatively high shear rates (10 – 100 rad/s, typical industrial processing conditions[59]), these polymer possess a relatively low complex viscosity at ~ 2000 Pa•s. It indicates excellent thermal processability of the materials. Additionally, these diblock polymers can be easily melt-compounded into various shapes, such as dog-bone, disk and film (Figure 5C), further confirming the great thermal processability.

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3. Conclusions In summary, we successfully prepared a new class of self-healable shape-memory polymers (SMPs) based on the diblock architecture endowed with supramolecular interactions. The diblock polymers, PMMA-b-P(BA-co-AMPS), behaved like typical shape memory materials

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with a high recovery ratio of about 95%. Meanwhile, these SMPs were mechanically tough, exhibiting an exceptional breaking strain of 500% and a high tensile strength of 10 MPa. The materials were effectively reinforced with the two network structures formed by the

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microphase separated PMMA domains upon vitrification as well as the ionic interaction from AMPS segments. Additionally, these block polymers displayed the functionality of rapid self-

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healing under ambient conditions, again benefitted from the microphase separation and the existence of supramolecular interactions. Another merit of these SMPs lied in the retention of thermal processability, making them very promising for industrial applications. We believe these cost-effective, mechanically robust, self-healable and thermally processable SMPs are

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excellent candidates of ‘smart materials’ to meet various service requirements of technological devices or other high-end products. This work also enlightens the design and fabrication of other multifunctional engineering thermoplastics via offering the approach of

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interactions.

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judiciously selecting the molecular architecture and making full use of supramolecular

4. Experimental Section Preparation of Polymethylmethacrylate (PMMA). [PMMA320, PMMA520]. PMMA was prepared via the Reversible Addition-Fragmentation Chain Transfer Polymerization (RAFT). PMMA320 and PMMA520 were prepared in the same fashion, but with different feeding ratios of MMA monomer and the RAFT agent. For example, PMMA320 was prepared by adding MMA (20 g, 21.1 mL, 200 mmol), AIBN (0.048g, 0.29mmol), RAFT agent (0.37 g, 1.38 mmol) into a 50 ml Schlenk flask with 10 mL DMF. The solution was degassed by three - 16 -

ACCEPTED MANUSCRIPT freeze-pump-thaw cycles and backfilled with nitrogen. The solution was placed in a water bath at 70 °C and stirred for 4 h. The product was precipitated three times in methanol. The solids were dried under vacuum for 12 h at 40 °C to obtain pink solids. Yield: 12.3g, 61.5%. For PMMA520, MMA (20 g, 21.1 mL, 200 mmol), AIBN (0.024g, 0.15mmol) and RAFT

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(0.19 g, 0.69 mmol) were used for polymerization.1H NMR (Acetone-D, δ, ppm) (Figure S1 and S2): 7.65 (benzene ortho, 2H), 7.95 (benzene para, 2H), 7.50 (benzene meta, 2H), 1.69 (CCH3, 3H), 2.56 (-CCH2C, 2H), 3.68 (-COOCH3, 3H). GPC curves were shown in Figure S3.

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Preparation of PMMA-b-P(BA-co-AMPS) [M520-b-B1100P160, M320-b-B1030P233, M320-b-

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B1080P170, M320-b-B1050P50]. M520-b-B1100P160, M320-b-B1030P233, M320-b-(B1080P170) and M320b-(B1050P50) were prepared in the same fashion, but with different feeding ratios of BA and AMPS. For example, M320-b-B1080P170 was prepared by adding PMMA320 (1 g, 0.031 mmol), BA (7.2 g, 8.1 mL, 56.3 mmol), AMPS (1.3 g, 6.2 mmol) and AIBN (2.2 mg, 0.013 mmol) in a Schlenk flask with 9 mL DMF. The solution was degassed by three freeze-pump-thaw

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cycles and backfilled with nitrogen. Then the solution was placed in a water bath at 70 °C and stirred for 4.5 h. After reaction, the product was precipitated three times in excess amount of

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water to remove BA. The polymer was then dried under vacuum for 12 h at 40 °C to obtain pink solids. Yield: 4.6g, 48.4%. 1H NMR (DMSO-d6, δ, ppm) (Figure 1B): 3.68 (-

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COOCH3,3H), 4.08 (-COOCH2CH2CH2CH3, 2H), 1.62 (-COOCH2CH2CH2CH3, 2H), 1.45 (COOCH2CH2CH2CH3, 2H), 0.9 (-COOCH2CH2CH2CH3, 3H), 1.44 (-NH2C(CH3)2CH2-, 6H). For M320-b-(B1030P233) and M320-b-(B1050P50), BA (7.2 g, 56.3 mmol), AMPS (2.6 g, 12.4 mmol) and BA (7.2 g, 56.3 mmol), AMPS (0.78 g, 3.7 mmol) were used for polymerization, respectively. For M520-b-(B1100P160), PMMA520 (1.5g, 0.03mmol), BA (7.2 g, 8.1 mL, 56.3 mmol), AMPS (1.3 g, 6.2 mmol) and AIBN (2.2 mg, 0.013 mmol) were used for polymerization. 1H NMR spectra were shown in Figure S4 - S6.

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ACCEPTED MANUSCRIPT Kinetic Studies of Polymerization for M320-b-(B1080P170). The kinetic studies followed the procedures as below. PMMA320 (1 g, 0.031 mmol), BA (7.2 g, 8.1 mL, 56.3 mmol), AMPS (1.3 g, 6.2 mmol) and AIBN (2.2 mg, 0.013 mmol) were added in a Schlenk flask with 9 mL DMF. The solution was degassed by three freeze-pump-thaw cycles and backfilled with

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nitrogen. Then the solution was placed in a water bath at 70 °C to start the reaction. Aliquots were taken after different time intervals (0.5 h, 1 h, 1.5 h, 2.5 h, 3.5 h, 4.5 h and 6 h) under the protection of nitrogen. 1H NMR spectra were used to determine the conversion of monomers.

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Preparation of PMMA-co-PBA-co-PAMPS (M310-co-B1030-co-P155). MMA (1.1 g, 11

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mmol), BA (7.2 g, 56.3 mmol), AMPS (1.3 g, 6.2 mmol), RAFT agent (18.5 mg, 0.069 mmol) and AIBN (2.4 mg, 0.015mmol) were added in a Schlenk flask with 9 mL DMF to synthesize the random polymer M310-co-B1030-co-P155 as a control. The solution was degassed by three freeze-pump-thaw cycles and backfilled with nitrogen. Then the solution was placed in a water bath at 70 °C and stirred for 4.5 h. After reaction, the product was precipitated three

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times in excess amount of water to remove BA. The polymer was dried under vacuum for 3 h at 40 °C to obtain pink solids. Yield: 5.1 g, 53.1%. 1H NMR (Chloroform-D, δ, ppm) (Figure

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S7): 3.68 (-COOCH3,3H), 4.08 (-COOCH2CH2CH2CH3, 2H), 1.62 (-COOCH2CH2CH2CH3, 2H), 1.45 (-COOCH2CH2CH2CH3, 2H), 0.9 (-COOCH2CH2CH2CH3, 3H), 1.44 (-

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NH2C(CH3)2CH2-, 6H).

Mechanical Test of Self-Healing. Dog-bone-shaped samples (length: ~38 mm, width: ~15 mm, thickness: ~0.5 mm) were prepared by melt-compounding of M320-b-B1080P170 and M320b-B1030P233 at 130 °C following the same procedures as described above. The specimen was cut to 70% ~ 90% depth of sample thickness. The damaged samples were self-healed under ambient conditions for a certain time interval. The self-healing time intervals being tested were 1 hour, 3 hour and 6 hours. The healed dog-bone specimens were used for uniaxial tensile tests at room temperature following the procedures as described above. - 18 -

ACCEPTED MANUSCRIPT Photo Recording of the Self-Healing Process 0.1 mg methylene blue stain was added into 0.5 g M320-b-B1080P170 to stain the polymer. Dog-bone-shaped specimens were prepared in the same fashion as above. The specimens were cut into two separate pieces. The stained piece and untreated piece were placed in contact with each other to self-heal for 3 hours. The healed

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dog-bone specimen was stretched and a digital camera was used to record photos of the same specimen under different stretching lengths.

Shape Memory Measurement via DMA. The melt-compounded specimen (M320-b-

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B1080P170) (length: ~13 mm, width: ~5 mm, thickness: ~0.5 mm) was used for shape memory

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measurements. It was first heated to 80 °C at a rate of 3 °C/min from 25 °C and maintained at 80 °C for 5 minutes. Then, the specimen was cooled to 25 °C at a rate of 3 °C/min and a static pulling load of 0.03 MPa was applied to the specimen during the cooling process. After maintaining at 25°C for 5 minutes, the stress was removed. Finally, the sample was reheated to 80 °C at a rate of 3 °C/min and held for 70 minutes. The strain was monitored during the

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entire measurement to obtain the recovery ratio of the specimen. Photo Recording of the Shape Memory Process The melt-compounded specimen (length:

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~38 mm, width: ~5 mm, thickness: ~0.5 mm) (M320-b-B1080P170) was placed in an oven at 80 °C for three minutes and then was twisted into a helix with a length of 20 mm. The helix

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was taken out from the oven and allowed to cool to room temperature on a lab bench and then placed back in the oven at 80 °C to observe the shape recovery. A digital camera was used to record the whole process.

Acknowledgements Funding for this work was provided by the National Natural Science Foundation of China (Grant No. 21504013 and 21774020) and the start-up in Southeast University (No.

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ACCEPTED MANUSCRIPT 1107047110). SAXS experiments were performed in Prof. Liangbin Li’s lab at the University of Science and Technology of China. We would like to extend our sincere thanks to Prof. Liangbin Li and Xiaowei Chen. We also appreciated the help with rheological experiments

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from Prof. Xiaoliang Wang’s lab in Nanjing University.

Reference:

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[56] Weiss, R. A.; Sen, A.; Pottick, L. A.; Willis, C. L., Block Copolymer Ionomers .2. Viscoelastic and Mechanical-Properties of Sulfonated Poly(Styrene-Ethylene ButyleneStyrene). Polymer 32 (1991) 2785-2792. [57] Register, R. A., Ionomers: Synthesis, structure, properties and applications. Springer Netherlands: New York, 1997. [58] Weiss, R. A.; Agarwal, P. K.; Lundberg, R. D., Control of Ionic Interactions in Sulfonated Polystyrene Ionomers by the Use of Alkyl-Substituted Ammonium Counterions. J. Appl. Polym. Sci. 29 (1984) 2719-2734. [59] Zhang, J. Y.; Li, T. Q.; Mannion, A. M.; Schneiderman, D. K.; Hillmyer, M. A.; Bates, F. S., Tough and Sustainable Graft Block Copolymer Thermoplastics. Acs Macro Lett. 5 (2016) 407-412.

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Supporting Information

Shape Memory and Self-healing Materials from Supramolecular Block Polymers Jiuyang Zhang,a,* Mengmeng Huo,a Min Li,a Tuoqi Li,b Naixu Li,a Jiancheng Zhou, a,* Jing a

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Jiang c School of Chemistry and Chemical Engineering, Southeast University, Nanjing, 211189,

China

The Dow Chemical Company, 2301 N. Brazosport Blvd, B-1608, Freeport, TX USA 77541

c

School of Chemistry and Chemical Engineering, Nanjing University, Nanjing 210093, China

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b

1. Materials

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2-Acrylamido-2-methyl-1-propanesulfonic acid(AMPS) (98%) was provided by Shanghai Aladdin Biochemical Technology Co., Ltd. Dimethylformamide (DMF) (99.5%) was purchased

from

Sinopharm

Chemical

Reagent

Co.,

Ltd.

4-Cyano-4-

(phenylcarbonothioylthio)pentanoic acid (RAFT agent) (Sigma Aldrich), butyl acrylate (BA) (99%, Aladdin) and methyl methacrylate (MMA) (99.0%, Aladdin) were purified through

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alumina oxide columns before using. 2,2-Azobis(2-methylpropionitrile) (AIBN) (98%, aladdin) was purified by recrystallization in ethanol before using. Methylene blue stain agent was purchased from Shanghai Aladdin Biochemical Technology Co., Ltd. All other reagents

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were used as received without further purification.

2. Instrumentation

H NMR 1H NMR spectra were collected on a Bruker Av400III HD instrument at 400 MHz,

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with chloroform-d, acetone-d6 or DMSO-d6 as the solvent. Gel Permeation Chromatography (GPC) Tests were performed on a PL-GPC 220 (Polymer Laboratories) at 40°C with tetrahydrofuran (THF) as the mobile phase. Polystyrene (PS) standards were used for the calibration and the detection was operated at a low rate flow of 1 mL/min. Solid Mechanical Property Experiments The dog-bone-shaped specimens with the geometry following ASTM standard D1708 (length: ~38 mm, width: ~15 mm, thickness: ~0.5 mm) were prepared by melt-compounding, which was carried out at 130 °C. Room temperature uniaxial tensile tests were performed on a SANS CMT4503 tensile tester with a crosshead - 24 -

ACCEPTED MANUSCRIPT speed of 5.0 mm/min. The reported results for tensile properties represented an average over at least three independent measurements on three identical specimens for each polymer sample.

Small Angle X-ray Scattering (SAXS) SAXS experiments were carried out using the

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beamline (X-ray Diffraction and Scattering) at the National Synchrotron Radiation Laboratory (Hefei, China) with a X-ray wavelength of 0.154 nm. A Mar345 image plate (2300 × 2300 pixels with a pixel size of 150 mm) was employed to collect 2D SAXS patterns. Samples were prepared as 0.5 mm thin film and all measurements were taken with 10 minutes

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exposure time. SAXS data was analyzed using the Fit2D software from European Synchrotron Radiation Facility, and the scattering intensity was reported in arbitrary units. In all cases the scattering patterns were cylindrical symmetric and were therefore reduced to the

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1-D form of intensity versus scattering wave vector magnitude, q = 4πλ-1sin(θ/2), where θ is the scattering angle.

Rheological Experiments A HAAKE RheoStress 600 strain-controlled rheometer with the 25 mm diameter parallel plate geometry and a forced convection oven with a nitrogen

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atmosphere was used for the measurements. All specimens were loaded to give a gap of approximately 1 mm. Temperature sweeps were first applied with temperature ranging from 40 °C to 190 °C with fixed strain (0.5%) and frequency (1 rad/s), and a ramping rate of 3 °C/min. Frequency sweeps were conducted in the range of 0.01 ≤ ω ≤ 100 rad/s at 170 °C.

regime.

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Strains were controlled to maximize the torque while maintaining the linear viscoelastic

Differential Scanning Calorimetry (DSC) Thermal transitions were determined by a SDT-

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Q600 thermal analyzer (TA). The specimens (5 - 10 mg) were loaded into hermetically-sealed aluminum pans. They were first heated to 150 °C rapidly and held at 150°C to erase any thermal history, followed by quenching to -80 °C, and then reheated to 150 °C at a rate of 20 °C/min. The glass transition temperatures were determined during the second heating runs. Thermogravimetric Analysis (TGA) A SDT-Q600 instrument was used for the thermogravimetric analysis. The specimens (3 - 5 mg) were loaded into open aluminum pans, which were heated to 800°C at a rate of 20 °C/min from room temperature in the air atmosphere. Dynamic mechanical analysis (DMA) Quantitative investigation of the shape memory - 25 -

ACCEPTED MANUSCRIPT function of diblock polymers was carried out on a TA Instruments Q800 in a stress controlled thin film tension mode with a specimen geometry: length: ~13 mm, width: ~5 mm, thickness: ~0.5 mm. The static preload was 0.03 MPa. The shape fixing ratio (Rf) was calculated by the following equation:

f

= ε u ×100%

ε

(1)

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R

m

Here, Ɛm is the strain after stretching and fixing and Ɛu is the strain after removal of stress. The strain recovery ratio (Rr) is calculated according to equation (2): r

= (1 − ε r ) ×100% (2)

ε

m

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R

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Here, Ɛr is the strain after recovery.

3. Synthetic Results

Table 1 summarized all the molecular information of the reported polymers. The results were collected from 1H NMR, GPC spectra and DSC curves in the manuscript and supporting

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information.

Figure S1. 1H NMR spectrum for the PMMA homopolymer/macroinitiator (Mn = 320 kDa). The inset is the zoomed-in spectrum displaying phenyl protons of the RAFT agent. The molecular weight of this PMMA polymer was calculated by comparing the integrated peak areas for phenyl protons and methyl protons (3.6 ppm). The molecular weight of the second - 26 -

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block ByPz in Mx-b-ByPz diblocks were then obtained according to the molar fraction of MMA, BA and AMPS based on the corresponding 1H NMR spectra.

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PMMA (320 kDa) PMMA (520 kDa)

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Response

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Figure S2. 1H NMR spectrum for the PMMA homopolymer/macroinitiator (Mn = 500 kDa). The inset is the zoomed-in spectrum displaying phenyl protons of the RAFT agent. The molecular weight of this PMMA polymer was calculated by comparing the integrated peak areas for phenyl protons and methyl protons (3.6 ppm).

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Time (min) Figure S3. SEC curves for the two PMMA macroinitiators with Mn (from 1H NMR) of 320 kDa and 520 kDa, respectively. - 27 -

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Figure S4. 1H NMR spectrum for the diblock polymer: M320-b-B1030P233. As shown in Figure S1, Mn of PMMA was calculated. Then, the number of repeat units of BA can be determined by peaks at 4.0 ppm by comparing with integral of PMMA peaks. The number of AMPS units are determined by peaks at 1.5 ~ 1.4 ppm, which overlapped methylene protons of BA and could be subtracted according to the already known BA units. Meanwhile, the number of AMPS units was also determined from 1H NMR conversion after polymerization, consistent with the above analysis.

Figure S5. 1H NMR spectrum for the diblock polymer: M320-b-B1050P50.

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Figure S6. 1H NMR spectrum for the diblock polymer: M500-b-B1110P160.

Figure S7. 1H NMR spectrum for the control random copolymer: M310-co-B1030-co-P155. The splitting of peaks at 3.5 ppm was due to random nature of the polymers, which was different from block copolymers.

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4. Estimation of Flory-Huggins Parameters and Order-Disorder Transition Temperature (TODT) Based on the functional group contribution method,[1,2] the solubility parameters of all components are estimated as: δ(MMA) = 9.18 √(cal/cm3)

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δ(BA) = 9.13 √(cal/cm3) δ(AMPS) = 12.41 √(cal/cm3)

For the second block P(BA-co-AMPS), it is a random polymer, according to the mixing

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rule,[3] the equivalent solubility parameters are: δ(B1080P170) = 9.58 √(cal/cm3) δ(B1050P50) = 9.28 √(cal/cm3)

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δ(B1030P233) = 9.74 √(cal/cm3) δ(B1100P160) = 9.55 √(cal/cm3)

Using the following relationship3 for estimating χ:

where

−

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=

is reference volume and we can use the MMA monomer volume, so 0.94 /

= 106.51

/

and

= 1.987 # /$

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100.12 /

The result at 25°C (298.15K) is:

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χ(M320/B1080P170) = 0.029, χN = 45.5, for this diblock, TODT ≈ 1009 °C χ(M320/B1050P50) = 0.002, χN = 2.84, for this diblock, TODT ≈ -200 °C χ(M320/B1030P233) = 0.056, χN = 88.7, for this diblock, TODT ≈ 2260 °C χ(M520/B1100P160) = 0.025, χN = 44.5, for this diblock, TODT ≈ 971 °C

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=

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Figure S8. Atomic force microscopy (AFM) images for M320-b-B1050P50 and M320-b-B1080P170. Both polymers are microphase separated into disordered structures. Scale bar: 500 nm.

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Polymers were spinning cast on silicon wafer and the resulted thin films were annealed under 130 oC for 30 minutes before characterization.

5. Thermal Properties

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Polymer powders are characterized by TGA. The degradation temperature is 200 °C, above the thermal processing temperature, indicating excellent thermal stability.

Figure S9. DSC curves of M320-b-B1030P233 and M520-b-B1110P160. Arrows denote the glass transition temperatures of the PMMA and P(BA-co-AMPS) domains.

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M320-b-B1080P170

80 60

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Weight (%)

100

40

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20 0 400

600

800

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200

Temperature (°C)

Figure S10. Thermogravimetric analysis (TGA) of M320-b-B1080P170.

6. Rheological Properties

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Melt-compounded samples were characterized by rheometer. Tan δ values from dynamic mechanical analysis (temperature sweep) were plot according to Figure 5A. Figure S10 showed transition temperatures corresponding to Tg’s and the dissociation temperature for the supramolecular interactions. At high temperature (> 160 oC), M310-co-B1030-co-P155 was melt into liquids, resulting in varied specimen thickness. So, the data points became very rough.

2.0 T = -29 oC

M320-b-B1080P170 M310-co-B1030-co-P155

T = 1.7 oC

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Tan δ

1.5

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M320-b-B1050P50

T = 140 oC

1.0

0.5

T = -21 oC

-30

0

30

60

90

120

150

180

o

Temperature ( C) Figure S11. Tan δ values from dynamic mechanical analysis (temperature sweep) for random polymer M310-co-B1030-co-P155, block polymers M320-b-B1080P170 and M320-b-B1050P50. Arrows denote the transition temperatures.

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104

103

10

,

M320-b-B1050P50

,

M320-b-B1080P170

,

M310-co-B1030-co-P155

1

0.1

1

10

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102

100

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G' and G'' (Pa)

Dynamic frequency sweep for samples under 170 °C. Under this temperature, M320-bB1080P170 and M320-b-B1050P50 behaved like solids, suggesting an phase separated state. Random copolymers, M310-co-B1030-co-P155, showed typical liquid-like behavior.

Temperature (oC)

Figure S12. Rheological behavior of the random copolymer M310-co-B1030-co-P155, and diblock polymers M320-b-B1080P170 and M320-b-B1050P50 at 170 °C. Closed and open symbols represent the storage modulus (G') and loss modulus (G''), respectively.

7. Tensile Properties:

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Mechanical results were collected from tensile tests and summarized in Table S1.

Figure S13. Representative engineering stress versus strain curves for self-healing tensile specimens of M320-b-B1120 and M320-b-B1050P50 after healing for 3 hours. Compared with samples (M320-b-B1120) without AMPS, M320-b-B1050P50 with only 50 units of AMPS showed

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it of M320-b-B1120 (12%).

Figure S14. Photos of M320-b-B1080P170 of (A) in original state; (B) after stretching under 80 o

C and then cooling to room temperature (force remained until cooling to room temperature);

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(C) Maintaining no shape change for two days under room temperature; (D) Heating again under 80 oC for 30 minutes. Samples kept their shape after fixing and could recover after a second heating.

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Table S1. Tensile properties of Mx-b-ByPz block polymers and random copolymers. σy (MPa)

Ea (MPa)

εb (%)

Toughness b (MJ/m3)

Self-healing Efficiency c

10.3 ± 1.3

42.1 ± 5.8

366 ± 36

29

-

M320-b-B1030P233(SH-1h)

4.8 ± 0.2

36.7 ± 0.3

360 ± 20

10

34%

M320-b-B1030P233(SH-3h)

8.4 ± 0.2

40.1 ± 6.5

454 ± 30

24

82%

M320-b-B1030P233(SH-6h)

7.0 ± 1.8

48.7 ± 17

250 ± 41

17

59%

M320-b-B1080P170

7.4 ± 1.1

5.4 ± 1.6

530 ± 265

21

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M320-b-B1080P170(SH-1h)

5.6 ± 0.8

6.0 ± 1

369 ± 29

15

71%

M320-b-B1080P170(SH-3h)

6.6 ± 0.6

6.5 ± 0.5

419 ± 8

18

86%

Samples

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M320-b-B1030P233

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1.0 ± 0.1

0.7 ± 0.1

715 ± 46

4.1

-

M520-b-B1110P160

10.4 ± 0.1

230 ± 27

19

-

M310-co-B1030-co-P155

1.3 ± 0.3

49.2 ± 2.1 1.2 ± 0.2

715 ± 81

6.9

-

a

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Room temperature Young’s modulus (E) determined from the slope of the fitting line to the linear elastic regime of stress-strain curves. b Tensile toughness determined by integrating the area under stress-strain curves up to the specimen breaking point. c Self-healing efficiency was calculated by ratio of the toughness of healed samples to their original ones.

8. Reference

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3.

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2.

J. Brandrup, E. H. Immergut, Eds.; Polymer Handbook, 4th ed.; Wiley: New York, NY, 1998; pp 1920. Gong, Chunli, et al. "Effect of sulfonic group on solubility parameters and solubility behavior of poly (2, 6‐dimethyl‐1, 4‐phenylene oxide)." Polymers for advanced technologies 18.1 (2007): 44-49. Hiemenz, Paul C., and Timothy P. Lodge. Polymer chemistry. CRC press, 2007; pp 276277.

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A new class of self-healable shape-memory polymers (SMPs) was successfully prepared based on the diblock architecture and supramolecular interactions. These SMPs had a shape recovery ratio at 95% and were mechanically tough, exhibiting an exceptional breaking strain of 500% and a high tensile strength of 10 MPa. Additionally, these SMPs are self-healable under ambient conditions.