Controlled 3D Shape Transformation Activated by Room Temperature

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Applications of Polymer, Composite, and Coating Materials

Controlled 3D Shape Transformation Activated by Room Temperature Stretching and Release of Flat Polymer Sheet Shuwei Wang, Guo Li, Zhao-Tie Liu, Zhong-Wen Liu, Jinqiang Jiang, and Yue Zhao ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.9b10071 • Publication Date (Web): 24 Jul 2019 Downloaded from pubs.acs.org on July 24, 2019

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ACS Applied Materials & Interfaces

Controlled 3D Shape Transformation Activated by Room Temperature Stretching and Release of Flat Polymer Sheet

Shuwei Wang,a# Guo Li,*a# Zhaotie Liu,a Zhongwen Liu,a Jinqiang Jiang,*a Yue Zhao,*b a

Key Laboratory of Syngas Conversion of Shaanxi Province, Key Laboratory of Applied Surface

and Colloid Chemistry, Ministry of Education, School of Chemistry and Chemical Engineering, Shaanxi Normal University, Xi’an, Shaanxi Province 710062, China

b

Département de chimie, Université de Sherbrooke, Sherbrooke, Québec J1K 2R1, Canada

KEYWORDS: polymers; deformation-responsive polymers; shape change; shape memory; controlled plastic deformation; surface patterning

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ABSTRACT Shape transformation of polymeric materials, including hydrogels, liquid crystalline and semicrystalline polymers, can be realized by exposing the shape-changing materials to the effect of a variety of stimuli such as temperature, light, pH, magnetic and electric field. Herein, we demonstrate a novel and different approach that allows a flat sheet or strip of a polymer to transform into a predesigned 3D shape or structure by simply stretching the polymer at room temperature and then releasing it from the external stress; that is, a 2D-to-3D shape change is activated by mechanical deformation under ambient conditions. This particular type of stimuli-controlled shape-changing polymers is based on suppressing plastic deformation in selected regions of the flat polymer sheet prior to stretching and release. We validated the design principle by using a polymer blend composed of poly(ethylene oxide) (PEO), poly(acrylic acid) (PAA) and tannic acid (TA), whose plastic deformation can be locally inhibited by surface treatment using aqueous solution of copper sulfate pentahydrate (Cu2+ ink) that crosslinks PAA chains through Cu2+-carboxylate coordination and, consequently, increases the material’s Young’s modulus and yield strength. After room temperature stretching and release, elastic deformation in the Cu2+ ink-treated regions leads to 3D shape transformation that is controlled by the patterned surface treatment. This facile and effective “stretch-and-release” approach widens the scope of preparation and application for shape-changing polymers.

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INTRODUCTION

The ability to tailor or transform the shape or geometry of polymeric materials enables a broad range of applications including switchable adhesives, optical devices, soft robotics and reconfigurable metamaterials.1-4 Current processing techniques, like conventional molding or 3D printing, can produce 3D objects of complex shapes, but the cost in equipment and time can rise sharply with the size and shape complexity of the products.5 Alternatively, stimuli-responsive shape-changing polymers, like shape memory polymers (SMPs), liquid crystal elastomers and hydrogels, have been developed to exhibit intricate shape morphing behaviors.6-18 These polymers can be processed into regular shapes by conventional molding methods and later transform into more complicated shapes under the stimulation of an external trigger. Generally, the complexity of shape change is determined by spatially controllable stimulation or some sort of heterogeneous structure in the responsive polymer. For instance, light and focused ultrasound have been utilized to induce complex deformation of SMPs via spatially controlled shape recovery, while hydrogels exhibiting solvent- or temperature-responsive shape change could be developed by altering the cross-linking density distribution in the network structure.19-23 Although there are many SMPs and hydrogels capable of complex stimuli-controlled shape change or morphing, the most explored stimuli are limited to a few ones, such as light, heat, magnetic field, focused ultrasound, electric power, pH and solvent, each of which has advantages and drawbacks.24-30 Therefore, it is continuing interest to develop new approaches and stimuli for shape changing polymers in order to extend their application scope.

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In this work, we demonstrate a new approach that enables a flat sheet or strip of a polymer to transform into a predesigned 3D structure by simply stretching the strip at room temperature and then releasing it from external stress. This facile “stretch-and-relax” approach makes use of polymers undergoing plastic deformation when imposed to external stresses. This type of polymers can be twisted, bended or stretched at room temperature, and the deformation persists after releasing the external force owing to re-constructed glassy state or crystalline network. The underlying mechanisms mainly involve forced motion and orientation of chain segments (amorphous polymers) or crystalline domains (crystalline polymers) along the strain direction.31-33 The plastic deformation remains stable until the temperature is raised above Tg (glass transition) or Tm (crystal melting) to activate strain recovery that is driven by entropic elasticity. On the contrary, when a strip of an elastomer is stretched at room temperature, the deformation is elastic and lost once the external force is removed, and the strip recovers its initial length. Our approach to leading a flat piece of polymer to change to a targeted 3D structure by just room temperature stretching and relaxing is based on the following consideration. While bending, folding, twisting and stretching are among the commonly applied types of deformation, only stretching leads to prominent plastic deformation. Prior to stretching, if additional chain crosslinking is introduced to selected regions of the polymer, forced motion of polymer chains would be hindered and the plastic deformation would be suppressed in those regions. Consequently, when the polymer

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is stretched, it cannot undergo a uniform plastic deformation, instead the plastic deformation occurs only in the untreated regions while the regions with additional crosslinking would be deformed elastically; when the external force is removed, the elastic recovery in the selected regions could result in a 2D-to-3D geometry change. For instance, a strip can bent if the middle area on one side is treated with additional crosslinking.

To validate the above designing principle, as a proof-of-concept study, we prepared a plastically deformable polymer blend and realized its 2D-to-3D shape transformation using the “stretch-and-relax” approach (Figure 1). The used polymer is prepared by solvent processing of a blend composed of poly(ethylene oxide) (PEO), poly(acrylic acid) (PAA) and tannic acid (TA), in which PEO is a crystallizable polymer and its crystalline domains are responsible for plastic deformation, hygroscopic PAA provides carboxyl groups as the sites for additional crosslinking, and can absorb a certain amount of water molecules from air that act as plasticizer to increase elasticity of the blend.34 TA molecules serve as physical crosslinkers through formation of H-bonding to adjust the mechanical properties of the blend.35-37 To introduce additional crosslinks in selected regions, the blend sample can be treated with an aqueous solution of copper chloride (referred to as Cu2+ ink), since each Cu2+ ion can bind with two carboxyl groups in PAA. Figure 1 schematically illustrates the whole shape-changing process and the structural evolution. As shown, by depositing the Cu2+ ink on one side of a strip of the polymer, a Cu2+-coordinated PAA network can

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be formed with a gradually decreased crosslinking density from surface to the inner part along the thickness direction; the room temperature stretching and release results in bending of the strip towards the side with addition crosslinking. Moreover, since reducing Cu2+ ions to Cu+ ions can induce the dissociation of Cu2+-carboxylate coordination, aqueous solution of sodium metabisulfite can be used to selectively treat the Cu2+-coordinated regions, and thus adjust the obtained shape of the film after stretching and relaxing.38 The underlying mechanism of stretch-and-relax induced shape transformation is investigated by dynamic mechanical analysis (DMA), tensile tests and wide-angle X-ray diffraction (WAXD). Various examples of 2D-to-3D shape transformations are presented. Using the same approach, 3D micro-patterns are also created on the surface of films with stretching/heating changeable morphology. We also demonstrate one possible application of this stretching-activated shape-changing polymer to reveal its application potential especially in engineering fields.

Figure 1. Schematic illustration of shape transformation induced by stretching and release of a flat polymer sheet at room temperature. The used polymer is a blend of poly(ethylene oxide) (PEO), poly(acrylic acid) (PAA) and tannic acid, while the Cu2+-carboxylate coordination crosslinks PAA chains in patterned regions.

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EXPERIMENTAL SECTION

Materials. Ammonium persulfate (APS), tannic acid, sodium metabisulfite, N,N,N',N'-tetramethylethylenediamine (TEMED) and polyethylene oxide (PEO, Mw: 1,000 kDa) were purchased from Chengdu Aike Reagent Co., LTD. (China) and used without further purification. Acrylic acid (AA) was purchased from Shanghai Aladdin Bio-Chem Technology Co., LTD. Copper sulfate pentahydrate was bought from Sinopharm Chemical Reagent Co., Ltd. Preparation of Poly(acrylic acid). Poly(acrylic acid) (PAA) was synthesized using the same free radical polymerization conditions as reported previously.39 In a typical run, AA, APS (0.5 wt% of AA) and TEMED (1 wt% of AA) were sequentially added into a round-bottom flask with a certain amount of ultrapure water to prepare the reaction solution. The solution was continuously stirred for 24 hours at room temperature for polymerization, before dialysis against water for 3 days to remove non-reacted monomers and initiators from the polymer. The product was finally obtained after freeze-drying treatment. The molecular weight was determined by Gel permeation chromatography (Mn: 531 kDa). Preparation of the blend. Certain amounts of PEO, PAA and TA were sequentially added into a round-bottom flask filled with dimethyl formamide (DMF), which was stirred for 12 hours at 90 oC to ensure that the components were completely dissolved. Afterwards the solution was filled in homemade plastic molds and the temperature

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was increased to 90 oC for solvent evaporation. The obtained film was further vacuum-dried in 80 oC oven overnight to remove residual solvent before being placed in air for at least 48 hours for absorption of water from air before use. Surface Treatment of the Films. Desired amount of copper sulfate pentahydrate was added in deionized water to prepare the Cu2+ ink with certain concentrations. Then, a commercially available Chinese brush was dipped in the solution to conduct surface treatment at selected regions of the film, followed by solvent evaporation at room temperature. Characterizations. The molecular weight of synthesized PAA was determined using an Agilent PLgel 5 μm MIXED-C with DMF as the eluent. X-ray diffraction (XRD) analysis was carried out on a Rigaku D/max2550VB3þ/PC using Cu (40 kV, 40 mA). X-ray photoelectron spectroscopy (XPS) data were obtained on an AXIS ULTRA system. The morphologies of the 3D micro-patterns are examined by confocal laser scanning microscopy (Keyence, VK-X250, Osaka, Japan) with the following settings: Laser wavelength 408 nm, power 0.95mW, fully automatic scanning. Thermal phase transitions were measured on a Q1000 DSC, the scanning rate was fixed at 5 oC/min. Dynamic mechanical analysis was performed on a Q800 DMA. Tensile behaviors of the blends were investigated using a Suns UTM2103 universal tensile test machine applying a 100 N load cell.

RESULTS AND DISCUSSION

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Figure 2. (a) Tensile behaviors of the PAA/PEO (80/20, w/w) blend with different TA contents at 50% relative humidity. (b) Photos showing stretching-induced plastic elongation and heating-activated shape recovery of the blend with 0.3 wt% TA. (c) The effect of elongation on shape fixity ratio (Rf) and shape recovery ratio (Rr) obtained in the room temperature shape memory process with the blend containing 0.3 wt% TA. (d) WAXD patterns of the blend with 0.3 wt% TA before stretching, after stretching and after heating to 80 oC before cooling to room temperature.

We first prepared and studied a series of samples to optimize the mechanical properties and ductility. A previous work found that the PEO/PAA (80/20 w/w) blend can be stretched at room temperature to a strain beyond 800% and retain most of the elongation after removal of the external stress.28 However, the sample becomes harder

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and brittle in relatively dry environments (less than 55% relative humidity). In the present study, the PEO/PAA (80/20) blend was chosen, and in order to increase its tolerance to humidity change and further improve the mechanical properties, tannic acid (TA), a naturally-derived polyphenol that can form multiple interactions with a wide range of polymers and small molecules, is added in the blend to form H-bonds with both PEO and PAA.40-43 The hydrophilic nature of TA also permits the blends to absorb more water from environment that act as plasticizers (Figure S1), resulting in moderate compromise of crystallinity for PEO (Figure S2). Figure 2a shows the effect of TA content in the PEO/PAA blend on its tensile behaviors at 50% relative humidity (RH). It can be clearly seen that with a very small amount of TA, the tensile strain is prominently increased: from 300% with no TA to 600% with 0.1 wt% of TA, to 820% with 0.2 wt% of TA, and to 1220% with 0.3 wt%. Further increasing the TA content to 0.5 wt%, however, results in lowered tensile strength and fracture strain. Therefore, in what follows, the PEO/PAA (80/20) blend with 0.3 wt% of added TA was used. Its stretching-induced plastic deformation and thermally activated strain recovery can be quantified using a process that is shown in Figure 2b, where from top to bottom, a sample (initial length: L0) is stretched to a certain elongation (Lm, 600% in the picture) at room temperature. After stress removal, a large part of tensile strain is retained and the sample with an elongated shape is obtained (Lt). The fixed strain can be released upon heating the sample above Tm of PEO to obtain a recovered shape (Lr). This process is also known as room temperature shape-memory process, which can be quantified using shape fixity ratio (Rf) that is determined as the percentage ratio of

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retained elongation (Lt–L0) to the total deformation (Lm–L0), and shape recovery ratio (Rr) is the ratio of the strain that recovered (Lt–Lr) to the total deformation (Lm–L0).7 As seen in Figure 2c, with increasing the strain from 100% to 1200%, the value of Rf first increases from about 72% to 79%, and then decreases to 66%, while the value of Rr slightly decreases from about 100% to 94%. The structural evolution of PEO crystals during the room temperature shape-memory process was revealed from WAXD (Figure 2d). Before stretching, there are two brag reflection peaks located at 2θ = 19.1° and 23.2°, the former is assigned to (120) planes and the latter contains the overlapping reflections from (132), (112), (212), (124), (204), and (004) planes, which have similar d-spacing of ~0.39 nm.44 After stretching the sample to 600% strain, the relative intensity of the two peak (I(120)/I(032)) is notably increased. This indicates that after a disruption, orientation and re-crystallization process activated by elongation,45 the PEO crystals are well oriented along the stretching direction, because the (120) plane is parallel to the direction of chain extension, and also the direction that has the fastest growth rate with high intermolecular interactions.46,47 Subsequently, after heating the sample to 80 oC followed by cooling, the peak at 23.2° reappears with almost the same intensity as that before stretching, indicating melting of oriented PEO crystals, relaxation of tensile strain driven by entropic elasticity, and isotropic recrystallization of PEO in the strain-free condition.

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Figure 3. (a) Images showing a suspended strip before and after chemical ink treatment. (b) XPS results of the blend treated by Cu2+ ink (ST2+) and subsequently reduced (STre). (c) DMA results showing the storage modulus (G’) vs. temperature for ST0, ST2+, and STre. (d) Tensile behaviors of ST0, ST2+, and STre. (e) Stretching-induced inhomogeneous plastic deformation of a strip after the treatment of Cu2+ ink in selected regions. (f) 2D-WAXD patterns of ST0 and ST2+ under different tensile stresses.

We then investigated how the introduction of Cu2+-carboxylate coordination as additional crosslinking in the polymer blend affects its mechanical properties and plastic deformation. After a thin and flat film is obtained through solution-casting and drying, its surface can be treated using aqueous solution of copper sulfate

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pentahydrate (Cu2+ ink) for coordination between Cu2+ ions and carboxylate groups on PAA chains. Our hypothesis is that upon stretching, the formation of ionically crosslinked PAA network would hinder the motion and elongation of PEO chains in the crystals and thus suppress plastic deformation. Surface treatment in selected regions can be realized by “writing” a pattern with a brush soaked with the Cu2+ ink, followed by heating to 60 oC for solvent evaporation. Such a patterning process can also be conducted using an inkjet printer equipped with Cu2+ ink, which can improve the precision of patterning. The penetration of Cu2+ ink into the polymer blend induces the coordination between Cu2+ ions and carboxyl groups and thus the formation of ionically crosslinked PAA network. Moreover, this network can be dissociated by reducing Cu2+ ions to their +1 oxidation state via treating the patterns using a reducing solution, such as sodium metabisulfite or sodium citrate, which should restore and unlock the chain mobility. The change of mechanical properties of the film caused by treatment with the Cu2+ ink can be noticed visually. As can be seen from the pictures in Figure 3a, the polymer strip before Cu2+ ink treatment is soft and flexible, if the central position is fixed in air, the two ends are bent downwards due to gravity. After the whole strip is treated by Cu2+ ink, its color turns from light-gray to blue, and the film becomes much harder and can well resist gravity and maintain its straight shape when suspended. After erasing the Cu2+ pattern on the right half of the film using aqueous solution of sodium metabisulfite (reducing ink), the ionic crosslinking is dissociated and the film becomes soft and flexible again, with the color recovered to light-gray. As a result, the reduced right half drops down again when the

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film is suspended. In what follows, the three states of the polymer network, i.e., native, with Cu2+-coordination, and after dissociation of coordination, are denoted as ST0, ST2+ and STre, respectively. XPS results shown in Figure 3b clearly reveal the effective reduction of Cu2+ ions by the reducing ink, as the characteristic peak of Cu2+ at 935.41 eV is almost disappeared while that of Cu+ at 932.51 eV emerges. The Cu2+ ink/reducing ink treatment induced mechanical property change of the polymer is also unveiled by dynamic mechanical analysis (Figure 3c) and tensile tests (Figure 3d). Figure 3c shows that at room temperature, the storage modulus (G’) of ST2+ is much higher than that of ST0 and STre, which is clearly due to the formation of Cu2+-carboxylate coordination. The even lower G’ of STre compared to that of ST0 is possibly caused by the products of the reduction reaction that act as plasticizers. From Figure 3d, it can be seen that the yield strength of ST2+ is 80.6 MPa, which is several times higher than that of ST0 (9.1 MPa) and STre (23.2 MPa). Since it is known that the yield strength must be exceeded to activate plastic deformation, when a polymer film partly treated using Cu2+ ink is subjected to increasing tensile stress, the untreated ST0 regions with much lower yield strength should undergo plastic deformation first while the treated TS2+ regions with high yield strength would only be elastically deformed during this process. This is confirmed by the simple experiment shown in Figure 3e. A polymer strip (5.0 cm in length) is first surface-treated (on the both sides) using Cu2+ ink in alternating regions along the length direction. After elongating the strip to 10.2 cm, only the ST0 regions (light-gray color) are narrowed and elongated due to plastic deformation, while no elongation can

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be noticed by naked eye in the ST2+ regions (blue color). Upon heating, the plastic deformation can be recovered owing to the melting of PEO crystalline domains. Figure 3f compares the 2D-WAXD patterns recorded for untreated (ST0) and treated (ST2+) regions of the polymer strip under various tensile stresses. Two diffraction rings corresponding to the (120) reflection (inner) and (032) reflection (outer) are visible in both ST0 and ST2+ regions before application of tensile stress, indicating the absence of macroscopic orientation. With the gradual increase of the tensile stress from 8.6 to 13.6 MPa, the rings become arcs and finally spots in ST0 regions, indicating that PEO crystals are well oriented along the stretching direction due to large chain extension arising from plastic deformation. By contrast, no changes occur in the ST2+ regions even under the tensile stress of 13.6 MPa, indicating the absence of PEO crystal orientation as a result of the suppressed plastic deformation due to the high yield strength (80.6 MPa).

Figure 4. (a) Photos showing the surface patterning of a rectangular flat sheet and its bending occurred after room temperature stretching and releasing from external stress.

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Plots of achieved bending angle as a function of (b) tensile strain (original thickness: 140 µm; Cu2+ concentration: 0.1 mol/L), (c) concentration of Cu2+ ink (initial thickness: 140 µm; strain: 200%), and (d) sample thickness (Cu2+ concentration: 0.1 mol/L; strain: 200%).

To realize controlled 2D-to-3D geometry change through the “stretch-and-release” process, the basic out-of-plane bending deformation was first investigated. This is achieved by treating the film on only one surface; the diffusion-controlled penetration of Cu2+ ink and the preferential binding of Cu2+ ions with carboxyl groups near the treated surface result in the formation of Cu2+-carboxylate coordination with a continuously decreased ionic crosslinking density along the thickness direction. Once stretched, the plastic deformation in the treated regions is suppressed to different extents as determined by the crosslinking density, their contribution to the elongation should be mostly elastic. Consequently, removal of the external stress results in bending deformation toward the treated side. For the sake of clarity, hereafter we use “pattern” to describe a film that is treated with the Cu2+ ink only on one side of surface with the purpose to create a gradient of ionic crosslinking density along the thickness direction. It can be envisioned that some parameters, like the concentration of Cu2+ ions in the ink, tensile strain and thickness of the film, would affect the bending behavior. In order to make the bending deformation predictable, we systematically investigated how these parameters affect the bending angle of the film after stretching and releasing. The experiment in Figure 4a was performed to show

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how the stretch and release can result in bending of the sample: a polymer sheet with a rectangular shape was surface patterned using Cu2+ ink in the central position (~ 2 mm in length); afterwards, the treated region of the sample was stretched and the external stress was then released, while the bending angle was recorded. We first investigated how tensile strain affects the bending angle. Rectangular sheets with a thickness of 140 µm were patterned using 0.1 mol/L Cu2+ ink before deformation, after stretching to various strains followed by relaxation, different bending angles were observed. As is shown in Figure 4b, higher strain leads to larger bending angle. This is understandable, since a higher elongation means a larger elastic deformation, which decreases along the thickness direction, in the patterned region. Secondly, Cu2+ ink with different concentrations were prepared and applied to do the surface patterning for the samples having a thickness of 140 µm before stretching the samples to 200% strain and then relaxing them. The results in Figure 4c show that sample patterned using higher concentration of Cu2+ ink exhibits greater bending deformation. This can be explained by that more concentrated Cu2+ ink results in higher ionic crosslinking in the treated region that can more effectively resist plastic deformation and give rise to larger elastic force in the patterned region. We also prepared a series of films with different thicknesses, patterned them using the same 0.1 mol/L Cu2+ ink, and then stretched them to 200% strain before removal of the external stress. The results on bending angle shown in Figure 4d reveal that thicker sample exhibit less prominent bending deformation. Since bending is mainly the consequence of competition between the ST2+ regions that suppress plastic deformation and the ST0

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regions that allow plastic deformation, there must exists an optimized balance of the two parts to get the maximum bending angle. Within our surveyed range of thicknesses, the actuation portion (elastic deformation in the treated region) seems to decrease with increasing the sample thickness, resulting in reduced bending angle. These three parameters can be considered as controlling factors and utilized in versatile combinations to achieve controlled 3D transformation through simply stretching and releasing flat sheets of the polymer.

Based on above results, a plausible mechanism of mechanical deformation-induced shape transformation of flat polymer sheets can be proposed. After surface patterning using Cu2+ ink, Cu2+-carboxylate coordination is formed in the patterned region. However, due to diffusion-controlled penetration of Cu2+ ions, not only the coordination density should continuously decrease from the surface to the inner part along the thickness direction, but also an interfacial region on surface with decreasing coordination density towards the two ends may exist. The formed coordination increases the material’s modulus and yield strength to varying extent, depending on the actual coordination density. Upon elongation, plastic and elastic deformation combines to determine the response of the polymer sheet after releasing the external stress. Essentially, the non-patterned side undergoes plastic deformation while the patterned side is more elastically deformed, so that upon stress relaxation the imbalanced elastic force bends the sheet towards the patterned side.

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Figure 5. (a) - (d) Examples showing controlled 3D transformation of flat polymer sheets patterned with Cu2+ ink through the room temperature stretching and release process. The dashed lines in the scheme indicate where a line cut is imposed using a blade. (e) Confocal laser scanning microscopy images showing the surface morphology change of a micro-patterned film subjected to the “stretch-and-release” process.

On the basis of the above investigation, customized flat sheets of the polymer with patterned ST2+ regions can be transformed into predesigned 3D shapes or structures simply by stretching and releasing the polymer at room temperature. To demonstrate this peculiar stimuli-controlled shape changing process, several examples are

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presented in Figure 5. In the first case (Figure 5a), strip-shaped ST2+ regions with altered oblique angle are distributed along the length direction of a rectangular film; after the “stretch-and-release” process, the flat sheet transforms into a cylindrical spiral with altered diameter along the stretching direction. In the second case (Figure 5b), the ST2+ region on one surface has a continuously changed density formed by controlling the time of patterning. Upon stretching and relaxing, the film turns into a concentric helix. In the third case (Figure 5c), a flat sheet is cut to have petal-like sections that are patterned using Cu2+ inks with different concentrations, after stretching and relaxing all sections, the sheet transforms into a flower-like structure with the petals curled differently as determined by the pattering. Moreover, combination of patterning and cutting of flat sheet can generate even more complicated 3D structures, as shown by the example in Figure 5d. Upon heating to above PEO crystal melting temperature, all these 3D structures can recover to their corresponding 2D shape through strain relaxation.

Besides macroscopic shape transformation, the morphology change on the patterned surface was also investigated. Surface patterning was conducted with a mask in between the film and the Cu2+ ink-depositing brush to form micro patterns on the flat surface (Figure S4). It can be seen from the confocal laser scanning microscopy images in Figure 5e that after mask-assisted surface patterning, 3D micro-patterns are formed with the ST2+ regions as ridges and ST0 regions as valleys. The formation of 3D micro-patterns can be explained by the absorption of moisture in the ST2+ regions

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that leads to swelling and increase in the thickness of the film, while the ionic Cu2+-carboxyl crosslinking retains the surface morphology after water evaporation. The morphology changes after stretching and release can be noticed from the altered distance between two ridges; after stretching the film to about 100% strain and releasing it from stress, the distance increases from 523.4 µm to 849.1 µm. Upon heating, the distance is recovered to 490.2 µm. The even lower value compared to that before stretching is probably due to relaxation of residual stress in the initial film.

Figure 6. Photos showing the reconfigurable 2D-to-3D shape transformation induced by stretching and release, which is achieved by selective erasure of Cu2+ ink patterned region using a reducing ink (aqueous solution of sodium metabisulfite).

As mentioned above, the chemical pattern formed using Cu2+ ink can be erased by reducing Cu2+ ions to Cu+ ions using a reducing ink (aqueous solution of sodium metabisulfite). This ability allows a flat polymer sheet initially patterned to achieve a certain 3D shape transformation to be reconfigured so that it displays a different shape change upon stretching and release. This re-configurability is interesting because it can endow a given polymer sheet with, in principle, an infinite number of shape changes by tuning the pattern of the ST2+ regions. An example is demonstrated in

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Figure 6, where a polymer strip is entirely patterned using Cu2+ ink on one side. Upon stretching and release, the film curls toward the patterned side and becomes a spiral, while subsequent heating induces the recovery of the film back to its original flat shape. By applying the reducing ink to treat partly the ST2+ region, the Cu2+-carboxylate coordination is dissociated as revealed by the color change, and the resulting STre region only shows homogeneous elongation upon stretching, while the unaffected ST2+ regions still curl. By increasing the portion of erased pattern, the 3D shape transformation through the “stretch-and-release” process varies. It should be noted that reduced Cu+ ions can be slowly oxidized by oxygen in the air to reform the Cu2+-carboxylate coordination, the effect of reversible formation of ST2+ regions on the shape-changing behaviors will be investigated in future studies.

Figure 7. Photos showing a Cu2+ ink patterned strip that forms a hook at one end after room temperature stretching and release, which is then used to hang, lift and release an object.

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Finally, compared to the widely used stimuli for polymer shape change, such as heat, light, magnetic and electric field, the stimulus of mechanical interaction between human and materials is peculiar, and may be convenient and cost-effective in some conditions. Developing stretching-responsive shape-changing material is worth exploitation. Shown in Figure 7 is a sort of stretching-induced hook that, when needed, can accomplish a physical work. The hook was prepared by surface patterning a polymer strip (1.8 cm in length) on one end (0.4 cm in length). After fixing the non-patterned end in air, the polymer strip was elongated by hand and then released from the stress to induce shape transformation, comprising the non-patterned part evenly elongated and the patterned end curled toward the object. After hooking the object, the evenly elongated part of the strip was heated using a blow dryer to elevate the object to a higher position. Afterwards, the object was horizontally moved to the target location and subsequently released by gently heating the bended part (hook) of the strip.

CONCLUSIONS In conclusion, we have made the first demonstration of controlled 3D shape transformation from flat polymer sheets activated by simply stretching and releasing the polymer at room temperature. We show that solution-processed PEO/PAA/TA blends exhibit superior ductility (> 1200% tensile strain) and prominent stretching-induced plastic deformation. After introducing Cu2+-carboxylate coordination by surface treatment, modulus and yield strength are much increased due

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to the ionic crosslinking, resulting in suppression of plastic deformation in these regions. The 2D-to-3D shape transformation enabled by the “stretch-and-release” process is realized by forming an ionic crosslinking density gradient along the thickness direction through patterning the film with Cu2+ ink on only one surface. This leads to out-of-plane bending deformation in the patterned regions due to imbalanced plastic and elastic deformations on both sides of the film upon stretching and release. After investigation on the effect of several experimental parameters (Cu2+ ink concentration, strain and sample thickness) on the bending deformation in order to make the shape transformation predictable, we showcased the power of this easy approach for controlled 2D-to-3D shape transformation. Moreover, the erasure of the Cu2+-carboxylate pattern by using a reducing ink was utilized to demonstrate the re-configurability of a patterned polymer film, so that a single piece of flat polymer film can be reconfigured to display different 3D shapes induced by the stretch-and-release process. The exploration of direct mechanical interaction between human and materials as a stimulus can widen the scope of interest in the field of shape-changing polymers.

ASSOCIATED CONTENT

Supporting Information. Water absorption behavior of the polymer blends with different TA contents; DSC curves of the blends with different TA contents; Effect of TA content on shape fixity and shape recovery ratios of the blends; a scheme showing

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the process to create micro-patterns on the surface of the blend; FTIR spectra of the polymer blends with or without TA addition; DSC curves of partly patterned sample during room temperature shape memory process; XPS results.

AUTHOR INFORMATION Corresponding Authors * E-mail: [email protected] (G.L.)

* E-mail: [email protected] (J.J.)

* E-mail: [email protected] (Y.Z.)

Author Contributions #

S.W. and G.L. contribute equally to this work.

Notes The authors declare no competing financial interest.

ACKNOWLEDGMENT The authors thank the funding from the Natural Science Foundation of China (Grant NO. 51803115), the Nature Science Foundation of Shaanxi Province (Grant No. 2019JQ-528), the Fundamental Research Funds for the Central Universities (Grant Nos. GK201601005, GK201801003, GK201802009, GK201901001 and GK201803031), the China Postdoctoral Science Foundation Funded Project (2017M623106), and Young Talent Fund of University Association for Science and

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Technology in Shaanxi, China (20180602). Y. Z. acknowledges the financial support from the Natural Sciences and Engineering Research Council of Canada (NSERC) and le Fonds de recherche du Québec: Nature et technologies (FRQNT). The authors acknowledge Mr. Chun Zhang for his help in XRD tests.

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Table of Content

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Controlled 3D Shape Transformation Activated by Room Temperature

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