2.1. Synthesis
All segmented multiblock (co)polymers were synthesized by the reaction of an aminobenzoyl-terminated polycaprolactone (
PCL4200 or
PCL8200) with terephthaloyl chloride in the presence of chlorotrimethylsilane at room temperature (
Scheme 1). The condensation is known to involve an
N,
N’-bis(trimethylsilyl)-substituted aromatic diamine intermediate which is susceptible to hydrolysis, an undesired side reaction that can be prevented by making the silylated intermediate
in situ [
14,
15]. A range of copolymers was subsequently made by replacing part of
PCL4200 or
PCL8200 with short-chain telechelics
T1–
T4 (
Scheme 2). These telechelics served as "chain extenders" and, like the chain extenders in polyurethane chemistry, they were expected to raise the hard segment content and thus the modulus of the copolymer (
Scheme 3). They were chosen because of their different segment length (increasing from
T1 to
T3/
T4) and their ready availability. We were also interested to determine whether shorter or longer segments within the chain extender would be more beneficial. Details of the composition of the polymers studied in this paper are summarized in
Table 1.
Scheme 1.
Structure of PCL4200/PCL8200 and typical polymerization procedure (soft segments are highlighted in blue, hard segments in red).
Scheme 1.
Structure of PCL4200/PCL8200 and typical polymerization procedure (soft segments are highlighted in blue, hard segments in red).
Scheme 2.
Structures of short-chain telechelics T1–T4 (“chain extenders”) used in copolymerizations.
Scheme 2.
Structures of short-chain telechelics T1–T4 (“chain extenders”) used in copolymerizations.
Scheme 3.
Schematic representation of a segmented polymer chain (a) without and (b) with the use of a chain extender illustrating the increase in hard segment content as a result of a copolymerization with a short-chain telechelic.
Scheme 3.
Schematic representation of a segmented polymer chain (a) without and (b) with the use of a chain extender illustrating the increase in hard segment content as a result of a copolymerization with a short-chain telechelic.
Table 1.
Details of polymer composition, molar mass averages, thermal transitions and breaking strain.
Table 1.
Details of polymer composition, molar mass averages, thermal transitions and breaking strain.
Polymer | Composition | Mn (g mol–1) | Mw/Mn | Tg (°C) | Tms (°C) | Tflow (°C) | %Strain at break |
---|
P1 | PCL4200 | 85,000 | 1.41 | –53 | 53 | 76 a | 630 |
P2 | PCL4200:T1 (2:1) | 110,000 | 1.84 | –52 | 46 | 179 | 850 |
P3 | PCL4200:T1 (5:1) | 74,000 | 1.95 | –55 | 47 | 66 | 640 |
P4 | PCL4200:T2 (2:1) | 63,000 | 1.54 | –52 | 52 | n.d. b | 8 |
P5 | PCL4200:T2 (5:1) | 85,000 | 1.39 | –58 | 46 | n.d. | 640 |
P6 | PCL4200:T3 (2:1) | 100,000 | 1.36 | –58 | 47 | 120 | 240 |
P7 | PCL4200:T3 (5:1) | 135,000 | 1.37 | –59 | 45 | 107 | 620 |
P8 | PCL4200:T4 (2:1) | 170,000 | 1.41 | –59 | 49 | 96 | Brittle |
P9 | PCL4200:T4 (5:1) | 234,000 | 1.41 | –57 | 50 | 83 | n.d. |
P10 | PCL8200 | 170,000 | 1.41 | –57 | 57 | ~60 | 890 |
P11 | PCL8200:T1 (1:1) | 110,000 | 1.34 | –57 | 55 | n.d. | 1,320 |
P12 | PCL8200:T1 (1:2) | n.d. | n.d. | –55 | 50 | >150 | 695 |
P13 | PCL8200:T3 (1:1) | 150,000 | 1.47 | –57 | 52 | 130 | 1,030 |
P14 | PCL8200:T3 (1:2) | 85,000 | 1.37 | –58 | 53 | 124 | 785 |
P15 | PCL8200:T3 (1:3) | 110,000 | 1.55 | –60 | 52 | 127 | Brittle |
P16 | PCL8200:T3 (1:4) | 90,000 | 1.68 | –60 | 54 | 147 | 510 |
P17 | PCL8200:T3 (1:5) | 70,000 | 1.60 | n.d. | 53 | 153 | Brittle |
P18 | PCL8200:T4 (1:4) | 110,000 | 1.35 | –48 | 48 | >150 | 1,540 |
2.5. Dynamic Mechanical Thermal Analysis
Dynamic mechanical analysis served as a key characterization technique for screening the series of (co)polymers for potential shape-memory properties [
13]. We will first discuss the data for polymers containing the shorter polycaprolactone
PCL4200 soft segments.
Figure 3a shows the temperature-dependence of the storage modulus of
P2–
P9. In all cases, the storage modulus decreased steadily above –60 °C, the
Tg of the soft segment. An additional drop in storage modulus was observed at around
Tms.
For optimum performance, the drop in modulus for a shape-memory polymer should not only be large and sharp, it should ideally be followed by an extended rubbery plateau. The DMTA measurement helped not only in determining the presence of a rubbery plateau but also in identifying an upper service temperature. The flow temperature,
Tflow, was estimated as the temperature at which the storage modulus fell below 1 MPa (
Table 1). While all polymers showed the desired steep decline in modulus at around 50 °C, few had a flow temperature above 100 °C. Increasing length and amount of the soft segment noticeably lowered
Tflow which was ~70 °C for the
PCL4200-based homopolymer
P1 whereas for the
PCL8200-based
P10 (54 °C) it more or less coincided with the melting of the soft segments. This observation was not surprising and confirmed Kim
et al. who noted that the increase in soft segment content drastically lowers the rubbery modulus [
7]. As a consequence of the low flow temperature and a virtually non-existent rubbery plateau, polymers
P1 and
P10 deformed irreversibly upon heating above
Tms, and neither homopolymer exhibited a shape-memory effect.
The incorporation of a chain extender (
T1–
T4) was expected to increase both the hard segment and the flow temperature of the polymers, and this was indeed the case (
Table 1). The rubbery plateau widened and, not surprisingly, the rubbery modulus increased as well, particularly with the short-chain
T1 as the chain extender. The latter was also the most effective in raising the flow temperature again well above
Tms. A flow temperature above 100 °C was considered sufficient to avoid irreversible deformation upon heating to slightly above
Tms, a key feature of the shape-memory test cycle (see
Section 2.6). Three copolymers (
P2,
P6 and
P7) in the first series fulfilled this criterium.
Figure 3.
Plot of (a) the storage modulus and (b) tan δ as a function of temperature for polymers P2–P9.
Figure 3.
Plot of (a) the storage modulus and (b) tan δ as a function of temperature for polymers P2–P9.
Tan delta, the ratio of loss modulus (not shown) to storage modulus, provides an indication of how easily elastic energy is stored, or lost in the form of heat, during a stretch–relaxation cycle as occurs in a DMTA experiment or a cyclic tensile test (the standard test for shape-memory polymers). A good elastomer should possess a tan δ well below 0.1, whereas a tan δ > 0.5 is likely to lead to significant losses of elastic energy in form of heat (
i.e., damping).
Figure 3b shows a plot of tan δ as a function of temperature for
P2–
P9. A maximum is observed around the glass transition. Most polymers in this series exhibited a steep rise in tan δ above 50 °C, that is, at about
Tms. Only
P2 and
P6 had tan δ values below 0.1 up to 80 °C. All remaining polymers were prone to a high damping behavior already at 60 °C, the temperature chosen for the shape recovery experiments in the subsequent cyclic tensile tests.
Figure 4.
Plot of (a) the storage modulus and (b) tan δ as a function of temperature for polymers P10–P16. Arrows indicate the onset of irreversible sample deformation.
Figure 4.
Plot of (a) the storage modulus and (b) tan δ as a function of temperature for polymers P10–P16. Arrows indicate the onset of irreversible sample deformation.
Understanding the dependence of the shape-memory properties to the tan delta curve would make future work easier because unsuitable polymers could be sorted out prior to the time-consuming cyclic measurements.
Figure 4 shows the temperature-dependence of the storage modulus of the polymers containing the longer polycaprolactone
PCL8200 soft segment. In all cases, the storage modulus displayed a significant drop at
Tms, as desired for an SMP. DMTA measurements also revealed irregularities in the modulus
vs. temperature plots such as a step or slip, which coincided with irreversible sample deformation, for all polymers except
P12 and
P16. Several polymers did not show any rubbery plateau at all.
2.6. Cyclic Thermomechanical Tests
Polymers P2, P6, P7, P11, P12, P16 and P18 were selected for cyclic stress–strain tests because (i) their storage modulus vs. temperature curves showed the required extended rubbery plateau and (ii) tan delta was in most cases still comparatively low and indicated elastic behavior. Samples of these polymers were elongated to a strain of 300% at room temperature (cold drawing) and then heated to 60 °C (i.e., slightly above Tms) to initiate shape recovery.
Figure 5 shows the results of such cyclic tests. The first two cycles deviated most from each other. During the first cycle, the polymer chains were forced to partially disentangle within the material and, as in
Figure 2, the stress–strain curve looked more like that of a thermoplastic. After having been through one cycle, all polymers adopted the typical S-shaped stress–strain curve characteristic for an elastic material. The reproducibility of the test cycle is quantitatively expressed by the strain recovery rate (
Table 2). It took at least 3 cycles to achieve a steady recovery rate that remained more or less unchanged for subsequent cycles. From that point on, reproducibility was generally quite good except for
P12 where we suspected that the amount of hard segment was above its optimum. Among the other
PCL8200-containing polymers,
P11 and
P18 recovered to 99% already after 3 cycles whereas
P16 required five cycles until the shapes of the curves almost matched each other.
Strain fixity provides information about the stability of the temporary shape. Ideally, strain fixity would be 100%, and lower values indicate that a sample relaxes to some extent after drawing even without being triggered by heat to do so. This is not uncommon during cold drawing, but can be overcome by drawing at higher temperature together with immediate freezing of the temporary shape in a cooling bath. We have opted for a cold-drawing procedure in our experiments as it avoids an extra step and is also more representative for practical applications.
P2,
P6,
P7 and
P16 were the polymers with the highest strain fixity rates of °80% (
Table 2). It is interesting to note that the inclusion of a comparatively long chain extender based on a polytetrahydrofuran (
P6,
P7,
P16) or polycarbonate (
P18) was not detrimental to the shape-memory properties as might have been anticipated.
P16 was the most flexible of the shape-memory polymers which could be attributed to the incorporation of the flexible polytetrahydrofuran segments which provided some elastic properties leading to recovery of deformations already at room temperature. We also note that SMPs containing
PCL8200-based soft segments showed poor reproducibility in the cyclic tests even with increasing numbers of cycles which we attribute to loss of elastic energy in form of heat (
Figure 5e,f). It confirmed our previous observation that tan delta was high and rising above
Tms for these polymers.
To provide additional comparison with the literature, we performed a full shape-memory cycle for a selected sample which involved heating above Tms, deformation above Tms, cooling, storage in the deformed state, then reheating to trigger shape recovery. Only polymer P2 could be investigated in this way as most other polymers, particularly those containing PCL8200 soft segments, could not sustain large forces at elevated temperature (60 °C) and were prone to break or deform irreversibly. This is not unexpected given the large tan delta values for these polymers at temperatures exceeding 50 °C. An additional difficulty arose due to stress in the deformed sample decreasing notably with time at 60 °C, which resulted in an artificial improvement in εu the longer one waited. To minimise this effect in the cyclic experiment, the cooling process was limited to 15 minutes.
Figure 5.
Cyclic tensile tests for polymer (a) P2, (b) P6, (c) P7, (d) P11, (e) P16 and (f) P18.
Figure 5.
Cyclic tensile tests for polymer (a) P2, (b) P6, (c) P7, (d) P11, (e) P16 and (f) P18.
Table 2.
Shape recovery rate and strain fixity during the first five thermomechanical cycles (N = number of cycle) for selected polymers.
Table 2.
Shape recovery rate and strain fixity during the first five thermomechanical cycles (N = number of cycle) for selected polymers.
N | Shape recovery rate (%) | Strain fixity (%) |
---|
P2 | P6 | P7 | P11 | P12 | P16 | P18 | P2 | P6 | P7 | P11 | P12 | P16 | P18 |
---|
1 | - | - | - | - | - | - | - | 77.1 | 76.6 | 79.2 | 73.7 | 84.6 | 69.0 | 61.7 |
2 | 86.7 | 77.9 | 85.5 | 95.5 | 85.5 | 68.5 | 90.0 | 80.5 | 79.6 | 82.8 | 82.5 | 80.9 | 81.8 | 64.7 |
3 | 96.0 | 93.6 | 97.3 | 99.1 | 91.5 | 93.5 | 100 | 81.3 | 79.6 | 83.1 | 75.6 | 82.7 | 82.2 | 66.1 |
4 | 99.9 | 97.1 | 99.7 | 97.3 | 100 | 94.9 | 98.8 | 81.7 | 79.5 | 83.1 | 76.9 | 84.6 | 80.6 | 69.0 |
5 | 99.2 | 98.0 | 100 | 100 | 95.8 | 99.7 | 98.8 | 81.7 | 79.8 | 85.1 | 76.0 | 78.6 | 80.8 | 72.3 |
Figure 6 shows the cyclic stress–strain behaviour of polymer
P2 when a sample was stretched at 60 °C. The result is very similar to
Figure 5a. Since loading at 60 °C includes a break when relaxation takes place, improved reproducibility of ε
u led to an increase in the strain fixity rate to above 86%, about 5% higher than for loading at room temperature (
Table 1). Rabani
13 and Kim
7 showed previously that a reduction in the strain value from 300% to 200% also raises the strain fixity rate. From these initial results we conclude that, if it is possible to deform the sample at room temperature, it may be a better way to obtain compareable and more representative values for the shape memory properties of a polymer.
Figure 6.
Cyclic stress-strain behaviour of P2 when the sample was deformed at 60 °C.
Figure 6.
Cyclic stress-strain behaviour of P2 when the sample was deformed at 60 °C.
To confirm their suitability as shape-memory polymers, samples were finally tested manually. Films of these polymers could easily be drawn by hand at room temperature which resulted in the temporary (deformed) shape. The samples subsequently recovered their original shape immediately after they were warmed to 60 °C (
Figure 7).
Figure 7.
(a)–(c) A thin strip of shape-memory polymer P11 was stretched at ambient temperature to several times its original length. (d)–(f) Immersion of the stretched film in a warm water bath (kept at 60 °C) resulted in the film recovering its original length.
Figure 7.
(a)–(c) A thin strip of shape-memory polymer P11 was stretched at ambient temperature to several times its original length. (d)–(f) Immersion of the stretched film in a warm water bath (kept at 60 °C) resulted in the film recovering its original length.