Entanglement-based shape memory polyurethanes: Synthesis and characterization

Polymer 53 (2012) 5924e5934

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Entanglement-based shape memory polyurethanes: Synthesis and characterization Xinzhu Gu, Patrick T. Mather* Syracuse Biomaterials Institute, Department of Biomedical and Chemical Engineering, Syracuse University, Syracuse, NY 13244, USA

a r t i c l e i n f o

a b s t r a c t

Article history: Received 26 July 2012 Received in revised form 25 September 2012 Accepted 29 September 2012 Available online 8 October 2012

In this paper, we report on the shape memory behavior of a family of hardblock-free multiblock thermoplastic polyurethanes (TPUs) consisting of poly(ε-caprolactone) (PCL) and poly(ethylene glycol) (PEG). Due to high molecular weight (Mn > 200 kDa), high degrees of entanglement were achieved. Instead of conventional “hard” blocks, entanglements served as the physical crosslinks in this system, slowing stress relaxation (suspending flow) above the melting temperatures of the soft blocks long enough to allow facile elastic deformation for shape fixing. Moreover, highly deformed samples (>800%) imparted by plastic deformation at room temperature completely recovered their as-cast shape within 1 min when heating above the transition temperature. Upon tensile deformation, the constituent chains and domains of both PCL and PEG phases became oriented, while heating-induced recovery reversed this orientation, as evidenced by wide-angle and small-angle X-ray diffraction studies. The observed large recoverable deformation and fast actuation make these materials strong candidates for applications in such medical devices as self-tightening sutures. Ó 2012 Elsevier Ltd. All rights reserved.

Keywords: Polyurethanes Shape memory polymers Multiblock copolymers

1. Introduction Shape memory polymers (SMPs) are a class of smart polymeric materials that have the ability be manipulated to retain or “fix” a temporary shape, and later recover to its “memorized” original (permanent) shape upon an external stimulus such as heat, light, electromagnetic induction, and solvents [1,2]. Nowadays, SMPs have drawn much attention because they possess a lot of advantages over shape memory alloys such as high recoverable strains, flexible transition temperatures, easy processing, and low manufacturing cost [3]. The polymers designed to exhibit shape memory effect require two components on the molecular level: crosslinks (fixed phase) to determine the permanent shape, and switching segments with transition temperatures (Tt) to fix the temporary shape. SMPs are categorized into two groups according to the nature of crosslinks: chemically crosslinked SMPs and physically crosslinked SMPs. Based on the nature of switching segments, SMPs are further subdivided into two categories: SMPs with amorphous switching segment where Tt is the glass transition temperature, and SMPs with crystalline switching segment, where Tt is melting temperature.

* Corresponding author. Tel.: þ1 315 443 8760; fax: þ1 315 443 9175. E-mail address: [email protected] (P.T. Mather). 0032-3861/$ e see front matter Ó 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.polymer.2012.09.056

Among physically crosslinked block copolymers utilized as SMPs, thermoplastic segmented polyurethanes have been widely investigated for their shape memory effects [4e7]. Conventionally, polyurethanes are phase-separated multiblock copolymers consisting of alternating sequences of hard and soft segments. The “hard” blocks, which show the highest thermal transition (Th), act as physical crosslinks by way of polar interaction, hydrogen bonding, and/or crystallization, and are therefore responsible for the “permanent” shape. Above this temperature the polymer melts to erase the permanent shape and can be processed by conventional thermal processing techniques such as compression molding, extrusion, or injection molding. Meanwhile, the “soft” blocks serve to encompass the thermally reversible phase and the vitrification or crystallization of these soft segments is able to fix the temporary shape. After deforming the material above Tt, but below Th, the polymer networks often show “superelasticity”, where the polymer chains between crosslink points can deform freely. Following such elastic deformation, the temporary shape can be fixed by cooling the polymer below Tt, effectively storing the elastic energy exerted during the prior deformation. Finally, recovery of the original shape can be achieved (more or less, depending on the SMP) by reheating up to Tt, releasing the stored energy. Polyurethanes feature the advantage of easily tuning the mechanical properties and transition temperature by manipulating their compositions. Besides the conventional physically crosslinked SMPs, there is a category of amorphous polymers with ultra-high molecular

X. Gu, P.T. Mather / Polymer 53 (2012) 5924e5934

weight >106 g/mol that also exhibit shape memory properties due to their lack of flow above glass transition. The most widely known materials with these characteristics include polynorbornene (PN, NorsorexÒ) [8,9], which was developed in the late 1970s, and high molecular weight poly(methyl methacrylate) (PMMA) [10]. Such polymers feature a significant number of entanglements per chain and these entanglements function as physical crosslinks, which forms a three dimensional network that gives excellent elasticity above Tg. Such polymers demonstrate fast and complete shape recovery due to high entanglement density, evidenced by a flat rubbery plateau measured rheologically [1]. On the other hand, their high molecular weight limits processibility, generally forcing the use of a solvent. The biodegradable and biocompatible aliphatic polyesters, such as poly(3-caprolactone) (PCL), poly(ethylene glycol) (PEG), poly(lactic acid) (PLA), poly(glycolic acid) (PGA), and their copolymers and blends have been studied extensively for their potential applications as biomaterials and as environmentally friendly materials [11,12]. Among them, PCLePEG block copolymers were developed for their unique properties, such as amphiphilicity, permeability, selfassembly [13], biocompatibility [14,15], and controllable biodegradability [16], and they have been widely investigated for tissue repair [17], immobilization of biomolecules [18] and controlled drug delivery [19e21]. In this work, we hypothesized that entanglements, instead of conventional “hard” blocks, could serve as the physical crosslink in this system, therefore engender memory for multiblock polyurethanes. Thus, a family of high molecular weight hardblock-free multiblock thermoplastic polyurethanes (TPUs) consisting of PCL and PEG was synthesized and characterized. Both PCL and PEG blocks can be regarded as “soft” blocks in the multiblock polyurethanes in the sense that they share very similar melting temperatures and allow large deformation with memory of the stress-free state being derived from entanglements, not crystallization of one of the blocks. Here, we programmed the temporary shape by cold drawing below the transition temperature (here, Tm) of PEG and PCL blocks through a combination of elastic and plastic deformations [22,23]. The benefits of cold drawing programming include more convenient shape fixing process and enhanced recovery stress. Shape recovery was induced by heating above the transition temperature. This type of shape memory was also documented as “reversible plasticity shape memory (RPSM)” [24]. In contrast, conventional shape memory involves deformation above a critical point and cooling to fix the temporary shape, and later recovering only the elastic deformation region by heating. Particularly, the highly deformed sample (>800%) showed nearly complete recovery within 1 min when heating above the transition temperature. To our knowledge, such high recoverable strains have not been reported previously for a thermoplastic material.

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purification. As a representative example, we describe the detailed procedure to prepare [PEG]50e[PCL]50. In a 250 mL three-neck flask, 5.0 g (0.5 mmol) PCL diol and 5.0 g (0.5 mmol) PEG were dissolved in 100 mL dried toluene (Fisher, ACS Certified). Under the protection of gaseous nitrogen purge, the flask was heated to 50  C. After heating at 50  C for half an hour, the solution became transparent, indicating that both PCL diol and PEG had completely dissolved in toluene. Then 0.183 mL (1 mmol) of purified LDI (r ¼ 1.157 g/mL) and 3e5 drops of dibutyltin dilaurate catalyst (SigmaeAldrich) were added through a syringe into the 10 wt-% toluene solution. The reaction mixture was further heated to 90  C and kept at 90  C for ca. 8 h. The resulting toluene solution was then precipitated into a 6-fold excess of n-hexane, and washed with deionized water. The obtained products were then dried under vacuum for one week at room temperature for further processing and characterization. We have adopted a nomenclature to distinguish samples by composition where the subscripts are the feed weight percent for each segment. Therefore, [PCL]50e[PEG]50, for example, designates a multiblock copolymer consisting of PCL (10 kg/mol) and PEG (10 kg/mol) blocks with a feed weight percent ratio of 50:50. Since the molecular weight of both PEG and PCL blocks is 10 kg/mol and remains same for all the polymers, that characteristic is not designated in the sample nomenclature of this paper. The PCLePEG TPUs thus obtained were compression molded into films using a Carver 3851-0 press with heating platens. The polymers were sandwiched between two Teflon sheets and a 0.35 mm thick Teflon spacer was placed in between the sheets to control the thickness of the final films. A compressive stress of 0.4 MPa was applied at 90  C and held for 30 s. Then the platens were cooled by cooling water and the compressive stress was released at ca. 20  C, after approximately 3 min of cooling. The resulting films were flexible and 0.35 mm in thickness, the latter being determined by a digital caliper. The sample thickness was maintained at the same value for all the following characterizations unless otherwise mentioned. 2.2. Molecular characterization The molecular weight and molecular weight distribution of the TPUs were determined by gel permeation chromatography (GPC) with a Waters Isocratic HPLC System equipped with a temperature controlled differential refractometer (Waters 2414). Multi-angle laser light scattering was employed (Wyatt miniDAWN) using three angles (45 , 90 , 135 ) for in-line absolute molecular weight determination. THF solutions (w2 mg/mL) were passed through a 0.2 mm PTFE filter and injected at 35  C using THF as mobile phase. 1 H NMR spectra of the samples were obtained by employing a Bruker 300 MHz Avance spectrometer. Samples were made from CDCl3 solutions and run at room temperature. 2.3. Thermal analysis

2. Experimental section 2.1. Synthesis of PCLePEG TPUs Multiblock thermoplastic polyurethanes (TPUs) were synthesized from polycaprolactone (PCL) diol, polyethylene glycol (PEG), and lysine methyl-ester diisocyanate (LDI). A one-step method was used to synthesize the TPUs. Before each reaction, LDI (Kyowa Hakko Chemical Co., Ltd.) was purified by vacuum distillation. PEG (10 kg/mol, Fluka, Inc.) was purified by preparing a THF solution, precipitating into n-hexane several times and then filtering. These steps were repeated several times. The purified powder was then dried in a vacuum oven at room temperature overnight. PCL diol (10 kg/mol, SigmaeAldrich, Inc.) was used without further

The glass transition temperatures (Tg), melting temperatures (Tm), crystallization temperatures (Tc), enthalpy of melting (DHm), and enthalpy of crystallization (DHc) were determined using differential scanning calorimetry (DSC, TA Instruments Q200) under a continuous nitrogen purge (50 mL/min). The samples were made by encapsulating approximately 3e5 mg of polymer in a TA aluminum pan. All samples were heated to 160  C at a rate of 10  C/min (to erase any previous thermal history), then cooled to 85  C at 5  C/min, and heated once more to 160  C at 10  C/ min. Values reported for Tg, Tm, and DHm are taken from the second heating trace. High molecular weight PCL (80 kg/mol, Aldrich) and PEG (400 kg/mol, Aldrich) were purchased and used as comparisons for the thermal analysis.

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2.4. Microstructural characterization

mechanical testing and shape memory characterization unless otherwise mentioned. After loading each film specimen at room temperature under tensile stress, they were cooled to 80  C, thermally equilibrated, and finally ramped to 160  C at 3  C/min. The tensile storage modulus (E0 ) and loss tangent (tan d) values are reported. The ultimate tensile properties of PCLePEG TPUs were evaluated by uniaxial tensile testing employing Linkam TST-350 tensile testing system (Linkam Scientific Instruments, Ltd.) with a 20 N (0.01 N resolution) or 200 N (0.1 N resolution) load cell at room temperature. The samples were extensionally deformed at crosshead speed of 200 mm/s (192%/min) until they reached the extensional limit of the instrument, which corresponds to a displacement of 77.5 mm (giving a strain of 1240%). The samples were then fixed at this strain for approximately 3 min to observe any stress relaxation of the materials. The strain was decreased until the force dropped to zero by closing the clamps at 192%/min. This experiment was replicated three times for each composition. Young’s modulus was calculated by computing the initial slope of the stress vs. strain curve (0 < ε < 10%) using linear regression. The yield stress was determined by the maximum stress following the proportional limit, and the yield strain was the corresponding strain. Strain after relaxation was recorded as the strain at the point where the force became zero (after elongation).

In order to determine the microstructures of the thermoplastic polyurethane (TPU) films, wide-angle X-ray scattering (WAXS) and small-angle X-ray scattering (SAXS) was conducted simultaneously by employing a Rigaku S-MAX3000 (Woodlands, TX) instrument in transmission mode at room temperature. A MicroMax-002þ HighIntensity Microfocus Sealed Tube X-ray Generator operated at 45 kV and 0.88 mA was used to produce a beam of collimated Cu Ka X-ray radiation with the wavelength of 1.5405 Å. Data collection was controlled by SAXS 16.0.0.2 software (Rigaku Americas Corporation, Auburn Hills, MI). Scattered and diffracted X-ray radiation was collected by two detectors: an electronic detector for the small-angle region, and an image plate for the wide-angle scattered X-rays. The distance between sample and image plate was fixed at 120 mm for WAXS collection. The image plates (FujiFilm BAS-MS 2025) were read using an FLA-7000 IP reader operated with FLA-7000 software (version 1.0.1) from Fuji Photo Film Company. The resulting scattering angular range was 5 < 2q < 40 for a d-spacing range of 2e18 Å. The distance between sample and SAXS detector was fixed at 1550 mm allowing a d-spacing range of 2e400 nm to be probed. 2.5. Mechanical properties Dynamic mechanical properties of all TPUs were investigated using a dynamic mechanical analyzer (DMA). A TA Instruments Q800 apparatus was employed in tensile mode with a preload force of 1 mN, amplitude of 15 mm (tensile strain < 0.3%), static stress/ dynamic stress amplitude ratio (“force tracking”) of 108%, and an oscillation frequency of 1 Hz. The polymer films were first cut from the hot-press films into a dogbone geometry (ASTM D638-03 Type IV, scaled down by a factor of 4), featuring a width of 1.5 mm, gauge length of 6.25 mm, and total length of 28.75 mm, using a custommade dogbone cutting die (TestResources, Inc., Shakopee, MN). This dogbone geometry was employed for all the following

2.6. Shape memory characterization Reversible plasticity shape memory (RPSM) was studied for each sample. The sample was loaded in Linkam with a load cell (20 N or 200 N) and stretched at 200 mm/s (192%/min) at RT until it reached the extensional limit of the instrument. Then the force was released by closing the clamps at 200 mm/s (192%/min). After being removed from the Linkam apparatus, each sample was heated to 70  C and maintained isothermally for 1 min in a convection oven, resulting in shape recovery, before cooling to RT. The performance measures,

O

(x)

H

OH O

+

(y)

O

+

O

R1

H

n

O

m

O

O O

m

H

O

(x+y) NCO

OCN

Toluene, 90oC, 6 hr

Precipitation

Dibutyltin Dilaurate

Hexane

Washing Deionized water

O

O

O O

CH2

CH2

O

C n

O

C NH

CH2

CH

NH

O

O O

C

O

CH2

C 5

4

O O

R1 m

O

C

C

O CH2

O 5m

C

NH

CH2

CH

O NH

C

4

x

Scheme 1. Preparation of PCLePEG multiblock TPUs by reacting the PEG and PCL diol with a lysine-derived diisocyanate.

y

X. Gu, P.T. Mather / Polymer 53 (2012) 5924e5934 Table 1 Chemical properties of PCLePEG multiblock TPUs. The subscript numbers represent the block weight %. The molecular weight for both PCL and PEG segments are 10 kg/ mol. Materials

PCL/PEG actual ratioa, mol%

Mnb (kg/mol)

PDIb

[PCL]100 [PCL]70e[PEG]30 [PCL]60e[PEG]40 [PCL]50e[PEG]50 [PCL]40e[PEG]60 [PCL]30e[PEG]70 [PEG]100

100/0 70/30 58/42 48/52 38/62 29/71 0/100

445 363 439 357 211 212 262

1.4 1.1 1.2 1.1 1.2 1.4 1.5

a b

5927

Table 2 Summary of thermal transitions for PCLePEG multiblock TPUs evaluated using DSC. The heating rate was 10  C/min, and the cooling rate was 5  C/min. Materials

Tg ( C)

Tm ( C)

DHm (J/g)

Tc ( C)

DHc (J/g)

PEG400k [PEG]100 [PCL]30e[PEG]70 [PCL]40e[PEG]60 [PCL]50e[PEG]50 [PCL]60e[PEG]40 [PCL]70e[PEG]30 [PCL]100 PCL80k

54.4 48.9 55.0 48.6 53.0 49.4 51.2 59.0 64.1

65.4 55.0 51.8 51.7 50.2 52.8 51.7 51.4 55.7

130.4 98.8 85.9 83.9 81.0 78.3 73.6 55.7 60.4

47.3 39.9 38.7 24.7/31.9 24.9 25.6 19.2/23.9 29.3 30.2

49.5 85.2 82.3 80.1 77.2 72.3 67.9 48.8 60.1

Based on 1H NMR. Based on light scattering from GPC.

Deformation (%), Fixing ratio (Rf) and Recovery ratio (Rr), were calculated for each composition using Eqs. (1)e(3):

Deformation ð%Þ ¼

 Rr ð%Þ ¼

(1)

 Deformation % $100 1240 %

(2)

 Ls  Lr $100 Ls  Lo

(3)

 Rf ð%Þ ¼

  Ls  Lo $100 Lg

where Ls is the length of the sample after stretch, Lo is the length of original dogbone sample, Lg is the gauge length of the dogbone sample which is 6.25 mm here, and Lr is the length of the sample after recovery, respectively. Then tensile tests were performed using the Linkam TST-350 apparatus under the same experimental condition as before, applied to the fully recovered samples to compare the mechanical properties of the original sample and the recovered sample. The microstructures of deformed and recovered samples were investigated by WAXS and SAXS under the same experiment condition as described in Section 2.4. The RPSM effect was quantified using dynamic mechanical analyzer (Q800 DMA, TA Instruments) following a previously established four-step thermomechanical cycling method [24]. The sample was first stretched to a displacement of 9.5 mm (152%) at

a

a rate of 1 mm/min (16%/min) at 25  C, yielding a strain of 3 m. Then the strain was decreased at 0.2 mm/min (3.2%/min) until the force was unloaded to 1 mN, after which a large percentage of strain remained (3 u) for all the samples. Shape recovery was then induced by heating to 65  C at a rate of 5  C/min under this same small load (1 mN) to a recovered strain, 3 r. Finally, the sample was cooled back to 25  C at 5  C/min to complete the SM cycle. The fixing (Rf) and recovery (Rr) ratios were calculated for each sample using Eqs. (4) and (5), where each of the strains is defined above. This thermomechanical cycle, known as the “one-way shape memory (1WSM) cycle”, was then performed consecutively for two more times on the same sample.

 Rf ð%Þ ¼

 εu $100 εm

(4)

 εu  εr $100 εu

(5)

 Rr ð%Þ ¼

3. Results and discussion 3.1. Molecular characterization The synthesis of PCLePEG multiblock hybrid thermoplastic polyurethanes is shown schematically in Scheme 1 and follows the methodology previously reported for other segmented polyurethanes [25,26]. The urethane linkages were formed through the addition reaction between isocyanate groups of the lysine methyl-

b 1.0 W/g

Heat Flow (Exo Up)

Heat Flow (Exo Up)

(iii)

(ii)

(iii)

(ii)

(i) (i) 1.0 W/g

-80

-40

0

40

Temperature

80 (oC)

120

160

-80

-40

0

40

Temperature

80

120

160

(oC)

Fig. 1. (a) DSC second heating and (b) first cooling curves of representative PCLePEG multiblock TPUs: (i) [PCL]100, (ii) [PCL]50e[PEG]50, and (iii) [PEG]100. The heating rate was 10  C/ min, and the cooling rate was 5  C/min.

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X. Gu, P.T. Mather / Polymer 53 (2012) 5924e5934

110

50

32

44 42

80

40 70

38 36

Δ Ηm

34

Δ Ηc χc

50

L110, PCL L120, PEG

20

16

32

40

30 0

24

20

40

60

80

100

b

12

30

240

Young's modulus Yield stress

Fig. 2. Trends of heating and cooling enthalpies based on DSC data, and overall degree of crystallinity (cC) based on WAXS data as a function of PCL wt-% content for PCLePEG TPUs.

ester diisocyanate (LDI) and the hydroxyl groups of either poly(ethylene glycol) (PEG) or poly(3-caprolactone) (PCL) diol. Seven PCLePEG TPUs of varying compositions were synthesized in this study. Here, we varied the feed ratio of PCL/PEG to include 100/0, 70/30, 60/40, 50/50, 40/60, 30/70 and 0/100. The numbereaverage molecular weight and molecular weight distributions of the hybrid polyurethanes are detailed in Table 1. The molecular weight of the synthesized polymers varied in the range of 211 kg/mole445 kg/ mol, and their corresponding molecular weight polydispersities

Young's Modulus (MPa)

PCL Content (wt-%)

25

200 20 160 15

120

Yield stress (MPa)

60

28

Lhkl (nm)

46 90

Degree of Crystallinity (%)

100

Δ H (J/g)

a

48

10

80

40

5 0

20

40

60

80

100

PCL Content (wt-%) Fig. 4. (a) Crystallite sizes (Lhkl) of the principal diffraction peak of PCL segments at 2q ¼ 21.5 and of PEG segments at 2q ¼ 19.3 , and (b) Young’s modulus and yield stress of PCLePEG TPUs as a function of PCL wt-% content.

d-spacing (Å) 8.8

5.9

4.4

3.5

3.0

2.5

varied from 1.1 to 1.5. High molecular weight polymers were achieved, giving an expectation of high degrees of entanglement [27,28]. To our knowledge, such high MW is not commonly reported for thermoplastic polyurethanes [29,30]. 1H NMR analysis was used to quantitatively determine the molar ratio of PCL/PEG by

(vii) 104

5

(vi) 4

(iv)

(iii)

(ii)

Storage Modulus (MPa)

(v)

102 101

3 Tan δ

Intensity (a.u.)

103

100 2

10-1 10-2

1 10

-3

(i) 10-4

10

15

20

25

30

35

2 θ (deg) Fig. 3. WAXS patterns of PCLePEG multiblock TPUs: (i) [PCL]100, (ii) [PCL]70e[PEG]30, (iii) [PCL]60e[PEG]40, (iv) [PCL]50e[PEG]50, (v) [PCL]40e[PEG]60, (vi) [PCL]30e[PEG]70, and (vii) [PEG]100. The X-ray wavelength (l) is 1.5405 Å.

-80

0 -40

0

40

80

120

160

Temperature (oC) Fig. 5. Tensile storage modulus measurement for [PEG]100. The heating rate was 3  C/ min. The tan delta curve is also shown, indicating the ratio of the viscous to elastic contribution of the sample tested with a peak at Tg and a dramatic rise at Tm.

X. Gu, P.T. Mather / Polymer 53 (2012) 5924e5934

fact that only one Tg was detected indicates apparent compatibility of both PCL and PEG segments in the amorphous state. The melting temperature of [PEG]100 was 55.0  C, which was slightly higher than that of [PCL]100 (51.4  C). A single melting peak, ranging from 50.2 to 52.8  C, was found during the second melting process of all of the PCLePEG block copolymers. We postulate that the melting points of PEG and PCL segments are very close to each other, which results in the overlap of the melting peaks. This postulation will be supported by X-ray diffraction results, discussed below. No melting transition above this temperature was observed for these hardblock-free TPUs. The melting and crystallization temperatures of PCL80k and PEG400k were much higher than the corresponding high MW TPUs. The imperfect crystal structure with lower melting temperatures is thought to result from the segmented structures of multiblock TPUs. Among the TPUs, a clear trend in DHm and DHc was evident: as the percentage of PCL increased, the melting and cooling enthalpies gradually decreased in a manner approaching pure PCL TPUs (Fig. 2).

80 [PCL]50-[PEG]50 [PCL]100 [PEG]100 60

Stress (MPa)

5929

40

20

0 0

200

400

600

800

1000

1200

1400

Strain (%) Fig. 6. Stressestrain response of PCLePEG TPUs. Samples were deformed at 192%/min at room temperature.

comparing the integration value of proton signal of PCL at d ¼ 2.32 ppm (e(C]O)e(CH2)4eCH2eOe) to that of PEG at d ¼ 3.65 ppm (eCH2eCH2eOe). Before comparison, all of the integration values were normalized for a single proton. For all seven reactions, the PCL/PEG feed ratio, as indicated by the final number in each polymer’s naming scheme, was in good agreement with what was observed in the resulting product (Table 1). 3.2. Thermal analysis The thermal behavior of each composition was studied by DSC experiments. First cooling and second heating traces are shown in Fig. 1 for the following three representative compositions: [PCL]100, [PCL]50e[PEG]50, and [PEG]100. The results for remaining compositions are reported in Fig. S1. Tg, Tm, Tc, DHm, and DHc for all TPUs as well as the commercial PCL80k and PEG400k samples for comparison are summarized in Table 2. All TPUs showed one glass transition temperature (Tg) ranging from 48.6  C to 55.0  C. The

3.3. Microstructural characterization To further clarify the microstructure of our PCLePEG TPUs after processing, wide-angle X-ray scattering studies (WAXS) were conducted on hot-pressed films, revealing the dependences of microstructure on PCL/PEG molar ratios, shown in Fig. 3. [PCL]100 exhibited unoriented (isotropic) crystalline rings characteristic of the orthorhombic unit cell structure with two strong reflection rings centered at d-spacing of 4.1 Å (2q ¼ 21.5 ), & 3.7 Å (2q ¼ 23.8 ), that are indexed as the (110) and (200) respectively, as well as weak reflection rings at 5.6 Å (2q ¼ 15.8 ), and 3.0 Å (2q ¼ 29.9 ), that are indexed as (102) and (210) [31]. For [PEG]100, there were two diffraction peaks centered at d-spacing of 4.6 Å (2q ¼ 19.3 ), and 3.8 Å (2q ¼ 23.5 ), which are attributed to the 120 and diverse planes of PEG monoclinic unit cell, respectively [32]. Both the PCL and PEG components showed well developed crystalline structures in the five copolymers. As shown in Fig. 3, with the increase of PEG content (ii / vi), the diffraction intensity of reflection peaks for PCL crystals at 4.1 Å decreased while those related to reflection from PEG crystal planes at 4.6 Å increased and sharpened. The peak of PCL at 3.7 Å overlapped with that of PEG at 3.8 Å to yield a large and broad peak. Beneath diffraction peaks typical of PCL and PEG crystalline phases, a broad zone characteristic of the amorphous phase was observed. In total, these WAXS results revealed that crystalline phases of both the hydrophilic PEG blocks and

Fig. 7. Shape memory tests of PCLePEG TPUs hot-pressed film: (a) [PCL]100, (b) [PCL]50e[PEG]50, and (c) [PEG]100. Samples on the left of each pictures: original dogbone films; samples shown in the middle: stretched films by Linkam; samples on the right: films after stretch at room temperature and heat recovery at 70  C for 1 min.

X. Gu, P.T. Mather / Polymer 53 (2012) 5924e5934

hydrophobic PCL blocks independently exist in the multiblock TPUs, indicating micro-phase separation driven by thermodynamic incompatibility. This finding is consistent with what has been reported in other literature for similar polymers [33e35]. Furthermore, crystalline peaks were separated from the amorphous halo, and the overall degree of crystallinity (cC) was calculated by the following formula:

AC AC þ AA

(6)

where AC and AA are the areas below the crystallite peaks and the amorphous halo respectively. As shown in Fig. 2, the overall degree of crystallinity of [PEG]100 was 48%, and it gradually decreased to 32% in a manner approaching [PCL]100. We expected that the higher degree of crystallinity would lead to higher fixing ability of these SMPs, which will be discussed below. As the composition changed, the crystallite dimensions for both components changed as well. Here, the Scherrer equation (Eq. (7)) was used to calculate the apparent mean crystalline sizes, Lhkl, perpendicular to the (hkl) plane, where K is the Scherrer constant (here K was assigned a value of 1 for all samples), l the wavelength of the X-ray radiation used (1.5405 Å), b0 the full-width at half maximum (FWHM) in radians, and q is the scattering angle (Bragg angle).

Lhkl ¼

Kl b0 cos q

Dynamic mechanical analysis (DMA) was conducted on all compositions in tensile mode to gain an understanding of the linear viscoelastic thermomechanical properties of our materials, particularly TPU rubbery plateau (elastic moduli) and transition temperatures. After mounting each sample in the DMA tensile grips at ambient temperature, they were cooled to 80  C and ramped up to 160  C (or until the sample yielded) at 3  C/min while being subjected to a small oscillatory strain at a frequency of 1 Hz while

a

(7)

As shown in Fig. 4a, the size of PCL crystallites in [PCL]100 was 28.8 nm, and it gradually decreased to 14.9 nm with the increase of PEG component in the copolymers. On the other hand, the PEG crystallite dimensions gradually increased from 21.9 nm to 29.7 nm as its weight percent in the copolymers increased. The observed trends indicate competing crystallization in the copolymers, which may be explained in the following way. When the PCL component is dominant in the copolymers, as in [PCL]70e[PEG]30, the PCL unit crystallizes first, which has a tendency to suppress the growth of PEG crystallites. Similarly, when PEG is the main component in the copolymers, PEG crystallization dominates and effectively hinders the crystallization of the PCL phase, which results in smaller PCL crystals with lower melting temperature. (Note: Fig. 4b will be discussed further below and compared with the data of Fig. 4a at that point.) The microstructures in the TPUs were further explored employing small-angle X-ray scattering (SAXS). As shown in Fig. S4, each of the two homopolymers, [PCL]100 and [PEG]100, featured a SAXS peak at q ¼ 0.038 Å1, indicating a correlation length of 16.5 nm. These peaks are assigned as the long periods of the crystalline lamellae of the homopolymers. In the case of copolymers, the SAXS patterns revealed domains with two distinct periodicities, which were around 15 nm and 35 nm. These periodicities were nearly independent of the ratio between two segments. The

Table 3 Summary of deformation (%), fixing and recovery ratios of PCLePEG TPUs, which were deformed at room temperature and recovered at 70  C for 1 min. Material

Deformation (%)

[PCL]100 [PCL]70e[PEG]30 [PCL]60e[PEG]40 [PCL]50e[PEG]50 [PCL]40e[PEG]60 [PCL]30e[PEG]70 [PEG]100

819 865 763 787 819 859 851

      

22 20 8 33 35 20 13

Rf (%) 66 70 61 64 66 70 69

      

Rr (%) 2 2 1 3 3 2 1

100.3 99.2 100.4 99.7 99.4 101.8 99.7

80

60

40

20

b

0 25 20

Stress (MPa)

cC ð%Þ ¼

 $100

3.4. Mechanical properties

15 10 5

c

0 25 20

Stress (MPa)



one around 15 nm is thought to come from the crystal lamella, while the one at 35 nm is attributed to the block copolymer lamellae.

Stress (MPa)

5930

15 10 5 0

      

0.8 0.6 0.8 0.8 0.2 2.6 1.2

0

200

400

600

800

1000

1200

1400

Strain (%) Fig. 8. Stressestrain response of original PCLePEG TPUs dogbone films and the same films after deformation to 800% and recovery at 70  C. (a) [PCL]100, (b) [PCL]50e[PEG]50, and (c) [PEG]100. The recovered sample stayed at room temperature for 12 h before the tensile tests. (d) original samples, (- - -) samples after stretch and heat recovery.

X. Gu, P.T. Mather / Polymer 53 (2012) 5924e5934

measurements of tensile storage modulus (E0 ) and loss tangent (tan d) were made. The result for one representative composition is shown in Fig. 5. The remaining compositions showed similar behavior and their results are shown in supporting information. All samples showed three distinct thermomechanical transitions: the glass transitions (Tg) are located near 50  C, while the melting transitions (Tm) are evident at approximately 50  C. The third transitions appeared around 120  C, above which the samples showed fluid-like behavior. E0 was at or above 1 GPa below Tg, then it decreased gradually to ca. 100 MPa between Tg and Tm, and then dropped dramatically to ca. 1 MPa upon heating above Tm. The observed rubbery modulus plateau above Tm is quite broad in temperature for all TPUs, which is of great importance for shape memory applications. It is apparent that the entanglements function as effective physical crosslinks, yielding the targeted elastic mechanical behavior. Ultimate tensile mechanical properties were probed, revealing the PCLePEG TPUs to be mechanically robust, as originally desired. Representative stressestrain curves of three samples are provided in Fig. 6, and the tensile properties are summarized in Table S1. The polymers deformed initially in an elastic manner, followed by yielding with clearly visible necking observed at the yield point. Following this yield point, the materials could be stretched further with plastic-like deformation followed by strain-hardening to an extension of 1240% without breaking, reaching the maximum extension limit of our instrument for samples with the specific dimension utilized. The samples were held at this strain for w3 min to observe some stress relaxation of the materials. The strain was then decreased until the tensile stress reached zero. The Young’s modulus and yield stress for all TPUs are plotted in Fig. 4b. It was found that these characteristics of the two homopolymers were

5931

significantly higher than the five copolymers. Moreover, the trend in Young’s modulus and yield stress appeared to be correlated with average crystallite size, showing a minimum for copolymers and increasing for homopolymers when the crystallite size was maximized. Between the two homopolymers, [PCL]100 exhibited higher mechanical strength than [PEG]100. Indeed, it is widely reported that PCL is stiffer than PEG of comparable molecular weight and this is attributed to differences in cohesive strength of PCL and PEG crystals. 3.5. Shape memory characterization The robust semicrystalline multiblock TPUs were further investigated for their Reversible Plasticity Shape Memory (RPSM) properties by external heating. The reversible plasticity of our PCLePEG TPUs was first investigated using the Linkam TST-350 tensile testing system. Samples were deformed to 1240% at RT, and the strain was then released to w800%, where the stress dropped to 0 MPa, to observe the initial elastic shape recovery (this will be referred as the “stretched state”). This allows deformations between permanent and temporary shape to be w800%. Samples were then heated to 70  C for 1 min to trigger nearly complete shape recovery (this will be referred as the “recovered state”). Fig. 7 shows the picture of original dogbone samples, stretched samples and heat-recovered samples for three compositions, and the shape memory characteristics figures-of-merit (Deformation (%), Rf % and Rr %) are reported in Table 3. It was found that all compositions completely recovered within 1 min, with recovery ratios higher than 99.2% (Real-time movie of sample recovery can be found in Fig. S10). Based on our DMA result (Fig. 5), the recovery temperature (here 70  C) is within the

Fig. 9. (a) 2D WAXS patterns of representative PCLePEG TPUs in the stretched state; (b) 2D WAXS patterns of PCLePEG TPUs after recovery; (c) 2D SAXS patterns of PCLePEG TPUs in the stretched state; and (d) 2D SAXS patterns of PCLePEG TPUs after recovery. (i) [PCL]100, (ii) [PCL]50e[PEG]50, and (iii) [PEG]100. Stretch direction is horizontal.

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X. Gu, P.T. Mather / Polymer 53 (2012) 5924e5934

range of rubbery plateau (below the temperature where we first see fluidity), so the materials are elastomeric during recovery. Instead of traditional hard blocks in polyurethanes, entanglements that exist above the melting temperature serve as physical crosslinks in our systems, thus promoting recovery while preventing flow. The recoverable strain achieved in this manuscript (800%), is much higher than the maximum recoverable deformation reported in the literature for thermoplastic shape memory polymers (400%) [36,37]. It is important to note that the strains used in our shape memory characterizations are from crosshead displacement, not extensometry. As such, the reported strains are likely larger than actual strains within the Type IV dogbone guage length [37]. In examining the fixing ability, it was observed that Rf % was independent of the overall degree of crystallinity. We reason that the fixing ability of our SMPs gradually increases and eventually reaches a plateau as cC increases. The cC of our PCLe PEG TPUs range from 30% to 50%, which are thought to be high enough to ensure that the fixing ability of SMPs reach the plateau region, resulting in invariance in the Rf % of the TPUs. The ultimate mechanical properties of all heat-recovered samples were further investigated using tensile testing and compared with original samples. Importantly, we observed only minor loss in mechanical properties after the deforming and heat recovering process (Fig. 8). Supplementary video related to this article can be found at http://dx.doi.org/10.1016/j.polymer.2012.09.056. Wide-angle X-ray scattering (WAXS) and small-angle X-ray scattering (SAXS) were employed to elucidate the microstructures of PCLePEG TPUs at the deformed and recovered states. WAXS of three representative compositions at stretched state and recovered state are shown in Fig. 9a and b, respectively. The patterns for the copolymers at the original state showed three strong isotropic (unoriented) rings (Fig. S3(a)). The inner ring is attributed to the (120) reflection of PEG monoclinic unit cell, having a d-spacing of 4.6 Å; the middle ring belongs to (110) reflection of PCL orthorhombic unit cell, having a d-spacing of 4.1 Å; and the outer ring is the superposition of (200) planes of PCL crystallites and diverse PEG reflections (d-spacing of 3.7e3.8 Å). For [PCL]100 deformed to a fixed strain of 800%, the two isotropic layers reflections split into equatorial reflections (Fig. 9a-i). For [PEG]100 in the stretched state, the lower angle reflections (120) were located on the equator, while the higher angle reflections appeared at off-equatorial azimuthal angles (Fig. 9a-iii). The WAXS pattern for [PCL]50e [PEG]50 (Fig. 9a-ii) was a combination of that of two homopolymers at the stretched state, indicating both PCL and PEG phases were highly oriented by stretching. Upon recovery, nearly all of the orientation of the homopolymers and copolymers was removed, as shown in Fig. 9b. The SAXS patterns of the stretched samples (Fig. 9c) show equatorial anisotropic diffuse scattering orthogonal to the stretching direction, suggesting the creation of nano-voids or nano-fibrils elongated along the stretching axis [38,39]. A 4point pattern was observed on the meridian, which is typically observed for cold drawn poly(ethylene) and is interpreted as chevron-type tilts of crystallites due to local strain relaxation [38]. All these features disappeared upon heating (Fig. 9d). While the reversal in complex nanostructure observed is quite clear and reproducible, in-depth SAXS analysis is beyond the scope of this paper. Precise, in-situ quantification of the RPSM effect was performed using the DMA apparatus, though restricted to smaller strains than those reported in Figs. 7e9. Shape memory loops were performed on all compositions, and representative RPSM cycles are reported for the following compositions: [PCL]100, [PCL]50e[PEG]50, and [PEG]100 (Fig. 10). The remaining compositions showed similar behavior and their results are shown in Fig. S13. As shown in Fig. 10,

each sample was stretched at 25  C by gradually ramping the tensile strain to w150% (step 1e2). The materials experienced both elastic and plastic deformation during this deformation step, which is evident in the projection curves on the stressestrain plane (dashed green curve). The strain was decreased to observe the initial elastic shape recovery (step 2e3), and the samples then heated to 65  C to induce complete shape recovery (step 3e4). Finally, the sample was cooled back to 25  C to complete the

Fig. 10. Reversible plasticity shape memory cycles (RPSM) of representative PCLePEG TPUs hot-pressed films: (a) [PCL]100, (b) [PCL]50e[PEG]50, and (c) [PEG]100. Three cycles are shown: (d) first cycle, (e e) second cycle, and (- - -) third cycle. Strain vs. temperature curve (red) and stress vs. strain curve (green) are also shown for each composition. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

X. Gu, P.T. Mather / Polymer 53 (2012) 5924e5934 Table 4 Fixing and recovery ratios calculated for PCLePEG TPUs. All the numbers represent percentages (%). Materials

Cycle 1

Cycle 2

Cycle 3

Rf (%)

Rr (%)

Rf (%)

Rr (%)

Rf (%)

Rr (%)

[PCL]100 [PCL]70e[PEG]30 [PCL]60e[PEG]40 [PCL]50e[PEG]50 [PCL]40e[PEG]60 [PCL]30e[PEG]70 [PEG]100

67.4 56.0 50.4 47.9 62.2 62.2 75.5

106.4 96.9 118.9 127.0 93.7 111.3 91.7

67.8 58.7 53.2 51.2 56.0 59.0 75.9

90.4 89.0 100.9 95.5 91.3 101.2 91.6

67.3 58.8 53.7 49.3 55.5 59.2 75.3

92.6 90.4 102.3 98.0 93.0 97.8 90.5

cycle (step 4e5). The recovery event is evident in the projection curve on the strainetemperature plane (dashed red curve). This thermomechanical cycle (steps 1e5) was then performed consecutively for two more times on the same sample. The shape memory fixing and recovery percentages, Rf and Rr, calculated from the three shape memory cycles for all samples are shown in Table 4. Quantitatively, Rf for each composition were quite consistent though all three cycles tested, ranging from w49% to w75% across the different compositions. It is noted that under RPSM condition, the fixity is usually compromised compared to the Rf approaching 100% for conventional SMP due to the elastic recovery during the unloading step [40]. The recovery ratios of the first cycle (Rr1) for some compositions were higher than 100%. It is postulated that the residual stress resulting from the hot-press molding lead to >100% recovery, and the stress was reflected as oriented domains in X-ray patterns shown in Fig. S3. After removing the residual stress in cycle 1, Rr for cycle 2 and cycle 3 decreased, but still higher than 90%. 4. Conclusions A family of high molecular weight PCLePEG TPUs was synthesized from PCL diol (10 kg/mol) and PEG (10 kg/mol) chainextended with lysine-diisocyanate (LDI) using a one-step technique. In this study, the ratios between the two components were varied from 0/100 to 100/0. X-ray measurements indicate microphase separation between hydrophilic PEG blocks and hydrophobic PCL blocks. The entanglements that exist above the melting temperatures of the soft blocks served as physical crosslinks in this system, yielding a temporary three dimensional network that gives good elasticity above the melting transition, therefore prevent flow. When heating upon the melting temperature, the highly deformed samples (>800%) rapidly recovered. Both PEG and PCL phases were oriented tremendously by stretching. Considering the large deformation, fast actuation and good recovery, we envision potential of this material in medical applications encompassing self-tightening sutures, sensors, and actuators. Acknowledgments We gratefully acknowledge the financial support of the New York State Office of Science, Technology and Academic Research (NYSTAR) (CON01587) and partial support from Baxter Corporation. We also thank Dr. Angel Romo-Uribe (Universidad Nacional Autónoma de México) and Dr. Kazuki Ishida (Syracuse University) for insightful discussions concerning X-ray analysis. Appendix A. Supplementary data Supplementary data related to this article can be found at http:// dx.doi.org/10.1016/j.polymer.2012.09.056.

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