Reversibly Actuating Solid Janus Polymeric Fibers - ACS Applied

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Reversibly actuating solid Janus polymeric fibers Leonid Ionov, Georgi Stoychev, Dieter Jehnichen, and Jens Uwe Sommer ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.6b13084 • Publication Date (Web): 23 Nov 2016 Downloaded from http://pubs.acs.org on November 26, 2016

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ACS Applied Materials & Interfaces

Reversibly actuating solid Janus polymeric fibers Leonid Ionov1 *, Georgi Stoychev1, Dieter Jehnichen2, Jens Uwe Sommer2 1

College of Engineering, College of Family and Consumer Sciences, University of Georgia, Athens, GA 30602, USA

2

Leibniz-Institut für Polymerforschung Dresden e.V ., Hohe Str. 6, 01069 Dresden, Germany

KEYWORDS: Actuators, polymers, Janus, fibers

ABSTRACT It is commonly assumed that the substantial element of reversibly actuating soft polymeric materials, is chemical crosslinking, which is needed to provide elasticity required for the reversible actuation.

On the example of melt spun and 3D printed Janus fibers, we

demonstrate here for first time that crosslinking is not an obligatory prerequisite for reversible actuation of solid entangled polymers since the entanglement network itself can build-up elasticity during crystallization. Indeed, we show that not-crosslinked polymers, which typically demonstrate plastic deformation in melt, possess enough elastic behavior to actuate reversibly. The Janus polymeric structure bend because of contraction of the polymer and due to entanglements and formation of nano-crystallites upon cooling. Actuation upon melting is simply due to relaxation of the stressed non-fusible component. This approach opens perspectives for design of solid active materials and actuator for robotics, biotechnology and smart textile applications. The great advantage of our principle is that it allows design of self-moving

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materials, which are able to actuate in both water and air and, which are not crosslinked. We demonstrate application of actuating fibers for design of walkers, structures with switchable length, width and thickness which can be used for smart textile applications.

Introduction. Actuating polymers find application in medicine

1-3

, as sensors

4-5

, as imaging devices

6-7

, as

active elements of microfluidic devices 8-12 , in robotics 13-16, as elements of smart textiles 17 etc. There are many kinds of stimuli-responsive polymers, which are able to actuate in different environment and when different stimuli are applied

18-19

. One big group are the hydrogels

20-22

,

which are crosslinked polymers imbibed with water. They are able to change their volume upon swelling and de-swelling stimulated by change of temperature23-24, pH25-26 or light27-29. The hydrogels are able to actuate only in the presence of a solvent. Due to the high water content, hydrogels are very soft and are, therefore, extensively used as actuators for bio-applications, such as cell encapsulation 30 . Another group is shape-memory polymers (SMPs) 31-41, which are either chemically or physically cross-linked networks and consist of one or two polymers. Physical cross-linking is usually due to local crystallization of one of the polymers. Chemical or physical cross-linking determine the permanent shape. Heating the polymer brings it to the viscoelastic state where it can be easily deformed. Being deformed at elevated temperature and cooled down it retains its shape. Subsequent heating brings it back to the viscoelastic state and it recovers to its permanent shape. Dielectric elastomers

42

are also crosslinked rubbers placed between two

electrodes. Applying an electric potential to the electrodes leads to electrostatic attraction between them that deforms the elastomer. Removing the electric potential restores the shape43. Liquid crystalline elastomers

44

are highly oriented crosslinked polymers with mesogen groups.

Increasing the temperature leads to isotropization of the polymer, which in terms results in an


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anisotropic shape change – contraction along the orientation axis and expansion in the perpendicular direction 45-47. The initial shape is then restored upon cooling. Reversibly actuating shape memory polymers and liquid crystalline elastomers are very promising alternatives to the volume changing hydrogels and do not require water or any other solvent for actuation. Furthermore, such materials can exert stresses in the order of MPa. Chemical crosslinking of the material is a necessary requirement for polymer-based actuators. It prevents hydrogels from simply dissolving in the solvent. Without cross-linking, elastomers would deform irreversibly under external loads; the same is true for liquid-crystalline elastomers. One system that can lack chemical cross-linking is the shape-memory polymers, where physical cross-linking is usually sufficient. However, actuation based on the shape-memory effect of such systems is only possible if the physical cross-linking is not destroyed, for example by high temperature. Therefore, cross-linking, whether it be of chemical of physical nature, is an indispensable precondition for any polymer-based actuator design. In this work, we show for first time that crosslinking is not an obligatory prerequisite for reversible actuation of solid entangled polymers since the entanglement network itself can buildup elasticity during crystallization. This opens novel perspectives for the design of solid reversibly actuating materials with thermally healing properties using common processes such as extrusion and 3D printing. We show this approach on the example of non-crosslinked Janus sideby-side filaments made by melt spinning, as well as by 3D printing. The essential component of such materials is fusible polycaprolactone (PCL) with relatively high molecular weight (80 kDa). The second polymer is any polymer with a high softening point such as ABS (acrylonitrile butadiene styrene) or PLA (polylactide).


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Materials and methods Materials. Polycaprolactone (Mn ~ 80 kDa) was purchased from SigmaAldrich and was used without further modification. ABS pellets were received from Noztek and were used without further modification. For the 3D printer, blue-colored Hatchbox ABS filament (diameter 1.75 mm) was used. Fabrication of the fibers. The fibers were fabricated using homemade device consisting of two extruders (Noztek Pro, Noztek) connected to self-made spinneret. Polymers were extruded at 180 °C. The Janus fiber was collected on a rotating spindle. 3D printing. Janus structures were printed on a RepRap i3 Pro dual extruder 3Dprinter. Both ABS and PCl were printed at 240 °C through a 0.3 mm nozzle. PCl filament for the printer was prepared on a Noztek Pro extruder at 95 °C. DSC. Differential Scanning Calorimetry was performed on a Metler Toledo DSC821 measuring module. Samples were prepared by loading 5-10 mg of finely cut polymer pieces in a closed aluminium crucible. The polymers were scanned in three steps – heating from 0 °C to 180 °C, then cooling down to 0 °C, and then heating to 180 °C again. The heating/cooling rate was 10 K/min for all samples and all steps. DMA. Dynamic Mechanical Analysis was performed on a Perkin Elmer DMA 8000 in the dual cantilever mode. The samples were kept clamped for several minutes at room temperature before measurements to allow for any residual strain due to the clamping to dissipate. Storage and loss moduli were monitored under 0.5 % strain and 1 Hz. The temperature was ramped to 80 °C at 2 K/min. Rheology. Rheological measurements were performed on an Anton Paar MCR 302 rheometer, equipped with a heating bed. The polymer sample was first melted at 180 °C on the heated bed,


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then the measuring probe was lowered. The parallel plates geometry was used. The measuring distance was 0.3 mm for all samples. Multi-wave oscillatory mode was used. The storage and loss moduli of the materials were obtained for frequencies between 0.1 and 100 Hz. The temperature profile was acquired by cooling to 30 °C in 2 degrees steps; measurements were taken at each step. X-ray scattering investigations. The X-ray scattering experiments were executed by means of the multi-range device Ganesha 300 XL+ (SAXSLAB ApS, http://saxslab.com, Denmark/USA). We used Cu-Kα radiation (µ-focus tube 50 kV, 600 mA; monochromatization with bifocal Göbel mirror). Scattering intensities were accumulated by 2D-detector Pilatus 300K (pixel size 172 × 172 µm2). Path of rays, sample and detector are completely under vacuum (p < 5·10-2 mbar). For the actual investigations we applied a 2-slit configuration a beamstop with 2 mm diameter for two scattering range with following limits of parameters (Table 1) Table 1. Parameters for SAXS and WAXS (min...max): q (nm-1)

d (nm)

2Θ (°)

SAXS

0.05…2.0

125…3.14

0.007…2.8 ~1041

WAXS

0.9…24.4

7.0…0.257 1.27…34.8 ~101

lsample-det (mm)

tacc (s) 7200 1800

The experiments were realized in asymmetric transmission (beam perpendicular to sample surface (free-standing). The primary data were corrected to absorption. Results were presented as 2D scattering patterns lg I(qx, qy); radial profiles lg I(q) [orientation neglected]; azimuthal profiles I(ϕ) [here taken at q = 5 nm-1 (WAXS) and 0.4 nm-1 (SAXS), respectively].


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For partially crystalline polymers (here PCL) the lamellae thickness can be estimated from the d-value (Bragg’s law) of the first (strong) scattering maximum in SAXS. The orientation of the morphology (in SAXS) as well as that of crystalline structure (in WAXS) can be characterize using Herman’s orientation factor f = (32-1)/2). Here, θ denotes is the angle between the liquid-crystal molecular axis and the local director. The crystallinity was determined according to the peak area method with the crystallinity defined as the ratio of the integral intensity of the crystalline scattering to the total scattering (crystalline + amorphous). The calculations were performed using underground-corrected WAXS curves, in which all available scattering maxima (reflections and amorphous halo inside an appropriate scattering range) will be fitted by means of Pseudo-Voigt functions.

Results and discussion Investigation of actuation. We prepared Janus filaments using both simultaneous melt spinning by simultaneous extrusion and step-by-step FFF 3D printing of two polymers - PCL and ABS or PCL and PLA. The thickness of the drawn filament is 70 - 1000 µm depending on the drawing rate (Figure 1a,b,c). The thickness of the 3D printed structures is ca 1.2 -1.6 mm (Figure 1e). The ratio between polymers is nearly 1:1 in both cases. We observed that ABS-PCL fibers bend in the direction of the ABS side when it is heated to 65 °C (Figure 1f). Cooling down results in bending of the spun and printed filaments backwards and it adopts its original shape (Figure 1c, Figure 1g, d). The bending is fully reversible and could be repeated many times with the same amplitude (Video_S1, Video_S2, Video_S3, Video_S4). Moreover, we demonstrated


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that, in addition to actuation upon heating, fibers also reversibly actuated upon cooling to 4 °C (Video_S5). Thus, actuation is possible in the temperature range between 4 °C and ca 65 °C. We have found that a bilayer exposed to hot water (80 °C) bends and keeps its shape for more than 1 hour meaning that the actuation has not a kinetic character. The actuation behavior was also independent of the age of the sample - freshly prepared samples and samples, which were 3 months old, demonstrated the same actuation behavior. We observed actuation in both dry and wet state. The actuation rate also depends on the thickness of the fibers and 1 mm thick fibers actuate much slower (10-20 s) than 0.1 mm ones (2-3 s). Actuation upon heating is faster than upon cooling. As curvature of bending scales inversely with filament thickness, we estimated actuation velocity as ratio k h / t ≈ 1·10-4 m/(3·10-2m x 3s) ≈ 0.1 s-1 (where k is the curvature, h is the filament thickness and t is the time), which is much faster that other actuators (10 −5 –10 −2 s −1

) and is comparable to the most fast ones (0.1 s −1 ). 48

The actuation of the fibers is robust. In order to demonstrate this, we fixed the fiber by clamps to inhibit actuation. After heating above the melting temperature of PCl and subsequent cooling to room temperature the bilayer remained straight after release of the clamps. Increase of the temperature did not lead to actuation. However, cooling led to actuation again. In another experiment, we kinked the fiber and melted the PCL and it was still able to reversibly actuate (Video_S6). The fibers demonstrate also actuation behavior after their heating to more than 95°C when ABS starts to soften and deform (Figure S1). The Janus fibers are mechanically very robust - their Young’s moduli at room temperature and at 70 °C were 7 GPa and 5 GPa, respectively (Figure S2).

Thus, PCL based Janus structures possess a number of advantages. They

demonstrate fully reversible bending in dry state within the temperature range 4 - 80°C. They are


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rigid and their actuation is robust. They can also be fully biodegradable, if PLA is used instead of ABS.

Figure 1. Actuation behavior of ABS-PCL Janus filaments. (a)- Typical SEM image of a crosssection of a filament - two sides are clearly distinguishable; optical microscopy image (b,e) and camera shots (c,f) of the filament at room temperature and at >65 °C, respectively; (d,g) –printed Janus filament at room temperature and >65 °C, respectively. (Video_S1, Video_S2, Video_S3)

We characterized individual polymers to elucidate the origin of the actuation. ABS is amorphous polymer with Tg around 110 °C and its structure as well as length of its monofilament remains unchanged upon heating and cooling (Figure S2,S3). Its Young’s modulus is ca 10 GPa and remains almost unchanged up to its Tg (Figure S2). PCL is semicrystalline polymer with melting and crystallization points around 60 °C and 30 °C, respectively (Figure 2a). Young’s modulus of PCL is ca 2 GPa at room temperature and falls down to 0.5 GPa right before melting (Figure 2b). The storage modulus of PCL is larger than its loss modulus up to 120 °C at frequency 10 Hz, i.e. the molten polymer has elastic properties due to interpolymer chain entanglements (Figure 2c). The storage modulus is also larger than the loss modulus on all time scales at temperatures below


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the crystallization temperature of PCL. The length of the monofilament PCL fibers depends on temperature. The fiber contracts during first 5 cycles of change of temperature between 22°C and 70°C. Further cycling of temperature results in cycling change of PCL fiber length change of length (ca 2-3 %) (Figure 2d). Polymer chains in the “as prepared” PCL fibers are oriented along the fiber (Figure 2e1-e6, f1-f6) due to fiber stretching during extrusion. The lamellas with average thickness of ca 16 nm are also oriented along the fiber. The orientational order parameter (f) estimated from SAXS and WAXS data are f = 0.6 ± 0.1 and f = 0.43 ± 0.1, respectively. The crystallinity degree is ca 50 ± 5 %. Cycles of melting and re-crystallization of PCL have no essential impact on its degree of crystallinity and the size of lamellas, despite the fact that the polymer becomes isotropic when melted the first time after extrusion. Thus, ABS does not possess any thermoresponsive properties but PCL demonstrated alteration of the size and elastic modulus when the temperature was switched between 20 °C and 70 °C. The glass and melting transition of polymers in fibers correlate well with that reported in literature 49,50. The measured mechanical moduli are in general slightly higher than that reported in the literature

51,52

. We

explain this difference by orientation of polymers chains in the drawn fibers, the chains in dogbone samples used for standard testing are not oriented. The volume expansion of PCL upon melting cannot, however, be the origin of reversible actuation of Janus fiber because molten PCL does not possess elastic properties on the time scale of actuation.


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Figure 2. Properties of polycaprolactone: (a) DSC; (b) DMA; (c) rheological properties (d) relative change of length of monofilament PCL fiber upon cyclical switching of temperature between 22°C and 70°C; (e1-e6) SAXS and (f1-f6) WAXS results of PCL and Janus ABS-PCL fibers ”as prepared” and after actuation: 2D scattering patterns lg I(qx, qy); azimuthal profiles I(ϕ) [taken at q = 5 nm-1 (WAXS) and 0.4 nm-1 (SAXS), respectively].

In order to answer the question of why ABS-PCL Janus fibers actuate, we investigated the structure of polymers in Janus fibers and the dynamics of their change during actuation. In Janus fibers, the structure of the ABS side is the same as in the mono-component fiber and it does not change during the actuation. The structure of PCL in Janus filaments is, however, different from


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its structure in a mono-component filament. Although the size of lamellas in both cases is the same (ca 20 nm), lamellas of PCL in Janus filaments are isotropic (Figure 2e6). On the other hand, the PCL chains in Janus filaments are slightly oriented along the fiber length and its orientation is retained after melting of the polymer and crystallization (Figure 2f6). The apparent degrees of orientation of PCL chains53 in Janus filaments, which is underestimated because of scattering from ABS, is ca f = 0.26 ± 0.02 and f = 0.28 ± 0.02 before and after melting, respectively. The amorphous molten and the crystalline solid PCL can be clearly distinguished by their shade: the crystalline polymer is opaque, while the amorphous one is almost fully transparent. This difference allows for easy correlation between the shape of the fiber and the state of the PCL. We observed that the polymer filament is coiled and fully opaque before heating. The filament unfolds upon increase of temperature and it becomes slightly less opaque i.e. the degree of crystallinity of PCL decreases (Figure 3a, 4b). Thus, the actuation occurs during melting. Interestingly, the final melting of the PCL, when the PCL becomes fully transparent (Figure 3c), and the beginning of crystallization, when it changes from transparent to slightly opaque, do not lead to any bending (Figure 3d, e). The bending occurs during the crystallization (Figure 3f) and stops at the point when the filament becomes exactly as opaque as before its melting i.e. when the polymer has reached its maximal degree of crystallinity. The same was confirmed by experiment with very slow cooling down of the fiber from 75 °C to 22 °C (Figure 3i). The filament does not actuate at temperature above 50 °C, actuates slightly during cooling down to 38°C that corresponds to onset temperature of crystallization of PCL. The strongest actuation occurs during cooling to ca 34 °C. Cooling down to room temperature results in further actuation. Thus, actuation occurs not at the moment of total melting and beginning of


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crystallization but during decrease and increase of crystallinity when the temperature increasers or decreases, respectively. Based on the observation of actuation of Janus filaments and on the structure of the polymers, the following scenario of behavior can be suggested. During extrusion, ABS solidifies first, while PCL remains molten. As a result, PCL chains have time to relax and no orientation of lamellas is observed. Cooling down leads to crystallization and contraction of the PCL. Exactly in this moment the actuation starts. PCL demonstrates elastic properties because of crystallites, which play a role of physical crosslinking points, i.e. PCL acts like a physically crosslinked rubber (Figure 3j). These crystallites do not contact each other but are surrounded by fluid amorphous phase in the beginning of the crystallization. We believe that the fiber actuates because the polymer chains, as rheology showed, are strongly entangled and these crystallites are not independent but linked to each other via entanglements. The generation of stress (σ = 0.8 MPa) during the crystallization of monofilament PCL fiber fixed in clamps of a tensile testing machine is an indirect evidence of the formation of a network with elastic properties during crystallization. This force bends the ABS leading to actuation. Exactly this stress leads also to orientation of the PCL chains in Janus filaments upon crystallization. The actuation must stop as soon as the crystalline phase forms a network. The heating of a curled filaments leads to melting of the PCL and that allows the bent ABS side to relax and straiten. This scenario explains the difference in actuation speeds upon cooling and heating. Cooling actuation rate is low because of slow crystallization, while the melting actuation is fast because it is simply a relaxation.

Modelling of actuation. We modelled the behavior of PCL during crystallization by considering the elastic model sketched in Figure 3j. In particular, we consider the deformation


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between the faces of the crystalline domains, which shall by close enough to prevent disentanglement. In order to calculate the stress during crystallization, we define the effective deformation ratio of the elastic matrix in one direction of the sample by 1/ 3

L − (φυ ) L0 λ − (φυ ) = λ= L0 − φ 1 / 3 L0 1 − φ 1/ 3

1/ 3

(1)

where L denotes the extension of the sample in the given direction with L0 being the nondeformed extension before crystallization starts. The degree of crystallization is denoted by φ, and ν denotes the shrinkage of the specific volume of the crystalline phase with respect to that of the melt phase. The (overall) deformation ratio of the sample in this direction is denoted by λ . This equation can be derived by considering M crystalline domains having an average extension of R0 = (φV0 ⁄ M)1 ⁄ 3 without taking into account shrinkage, and R = R0v1 ⁄ 3 after shrinkage For simplicity we consider cubic shape of the domains such as drawn in Figure 3j. On the other hand their average center-to-center distance is given by D0 = (V0 ⁄ M)1 ⁄ 3. The non-deformed elastic length is thus given by D0 − R0. For the case of (affine) deformation in one direction, D is changing according to the external deformation. . Thus we can define the elastic deformation by λ = (D − R) ⁄ (D0 − R0) which is Eq.(1). For the case of φ = 0.5 and v = 0.8 one obtains for the fixed sample (L = L0) an effective elastic deformation between the crystalline domains of λ = 1.275. The free energy of the elastic matrix can be written as ଵ

‫ = ܨ‬ଶ ݃‫ܯ‬൫ߣଶ௫ + ߣଶ௬ + ߣଶ௭ ൯

(2)

Here, ߣ௫,௬,௭ denote the three spatial deformation ratios of the sample (we consider diagonal modes of deformation only). The symbol M denotes the sum of the number of crosslinks and entanglements in the crystalline state and g denotes a topological prefactor, which takes into


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account the connectivity of the network and can be considered as of order unity. We note the relation for the shear modulus G0 = gM0/V0 with V0 and M0 being the volume and the number of elastic links respectively of the non-crystalline sample. In order to understand the experiment for the pure PCL-filament fixed by clamps during crystallization, we assume a uniaxial deformation in the direction of the stress applied by the clamps. The overall volume shrinkage, υ ,  can be expressed by the relation

λλ2⊥ = 1 − φ (1 − υ ) = υ

(3)

where λ⊥ denotes the deformation ratio in the directions perpendicular to the fixation. We should take into account that during crystallization disentanglement processes might occur and that the number of crosslinks, formed by crystalline domains may change. In any case, crystallization will transform elastic monomers into non-elastic. However, as long there is no change in the number of crosslinks/entanglements this will not cause a change in the free energy as long as no deformation of the elastic strands take place since the shortened chains are again fixed in their averaged equilibrium conformations. Elastic deformation occurs only due to shrinkage of the crystalline domain by fixing the sample’s extension. One might take into account a change in the number of crosslinks/entanglements by disentanglement during crystallization: M = M0(1 − ψφv). A value of ψ = 1 correspond to complete depletion of elastic active links by the crystalline phase while ψ = 0 corresponds to depletion of entangled sequences from the crystalline phase. Inserting eq.(3) in eqs.(1) and (2) we obtain for the stress (σ) by fixing λ  = 1.

1 − ψφυ

σ =G υ

1/ 2

(1 − φ )

1/ 3 2

[1 − (φυ )

1/ 3

(

1/ 3

− υ 1 / 2 υ 1 / 2 − (φυ )

)] (4)


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which reduces to σ (φ ) ≅ Gφ (1 − υ ) for the case of φ

Reversibly Actuating Solid Janus Polymeric Fibers - ACS Applied

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