Azlin F.
Osman
a,
Yosephine
Andriani
a,
Grant A.
Edwards
a,
Tara L.
Schiller
b,
Kevin S.
Jack
c,
Isabel C.
Morrow
ac,
Peter J.
Halley
ad and
Darren J.
Martin
*ad
aAustralian Institute for Bioengineering and Nanotechnology, The University of Queensland, QLD 4072, Australia. E-mail: darren.martin@uq.edu.au; Tel: +61 3346 3870; Fax: +61 3346 3973
bDepartment of Materials Engineering, Monash University, Clayton, 3800, Australia
cCentre for Microscopy and Microanalysis, The University of Queensland, QLD 4072, Australia
dSchool of Chemical Engineering, The University of Queensland, QLD 4072, Australia
First published on 9th August 2012
The viability of siloxane-based thermoplastic polyurethane (TPU) nanocomposites as a new insulation material for implantable and electrically active medical devices is investigated. We show that manipulating and controlling the specific interactions between the TPU segments and the engineered nanofiller greatly varies the TPU properties. The incorporation of dual modified organofluoromica as the nanofiller successfully enhanced the tensile strength, toughness and tear strength of the TPU. This is due to the presence of dual surfactants, which form regions of higher and lower surface energy on the layered silicate surface, thus enabling molecular interactions between the organofluoromica and both the hard and the soft TPU segment populations. We show that the addition of a second choline-based modifier with reactive –OH functionality may lead to the formation of positively charged TPU chain end groups as a result of trans-urethanization reactions during high temperature compounding, thus introducing labile “grip-slip-grip-slip” interactions between the TPU and the nanofiller. These molecular interactions assist in achieving a reduced level of stiffening, while at the same time enhance the toughening mechanism. The increase in the creep resistance and retardation in the stress relaxation of the TPU provides evidence that the dual modified organofluoromica also serves to enhance the dimensional stability of the TPU.
We have recently published an extensive study on a series of solvent cast biomedical PDMS-based TPU nanocomposites containing high and low aspect ratio organosilicates with an octadecyltrimethyl ammonium (bromide) (ODTMA) single surfactant modification.24 However, this solvent casting approach led to poor dispersion of the high aspect ratio organofluoromica in the TPU matrix.24 In this present work, E5-325 TPU-organofluoromica nanocomposites were produced by melt compounding, which offers advantages over solvent casting by way of eliminating solvent and enhancing the filler dispersion. In addition, both single and dual surfactants were used to modify the fluoromica, so that the effects of these surface modifications on organo-fluoromica-TPU interactions could be better optimized and understood. This patented surface modification26 involves the exchange of dual polar and non-polar molecules onto the layered silicate surface, thereby forming regions of higher and lower surface energy for better compatibilisation and performance in polyurethane systems. Therefore, it was anticipated that the presence of dual modifiers would provide an additional driving force for promoting TPU intercalation, and consequently improve the organofluoromica dispersion and reinforcement in the TPU matrix. Herein, we describe the complete mechanical and morphological characterization of the E5-325 TPU and associated nanocomposite series and communicate the performance benefits of dual nanofiller surface modification.
ElastEon E5-325 TPU consists of a 1000 g mol−1 poly(dimethylsiloxane) (PDMS) and 700 g mol−1 poly(hexamethylene oxide) (PHMO) mixed soft segment in a 98:2 (w/w) ratio, and a hard segment composed of alternating 4,4′-methylene diphenyl diisocyanate (MDI) and 1,4 butanediol (BDO) sequences. The hard segment concentration is 32.5 wt%. This TPU was supplied by AorTech Biomaterials Pty Ltd. Fluoromica (Somasif ME100), a synthetic mica (tetrasilicic trioctahedral fluoromica), was supplied by Kobo Products, Inc. It is a fine white powder with an average platelet size of approximately 650 nm with the chemical formula Na0.66Mg2.68(Si3.98Al0.02)O10.02F1.96.25 The surface modification of the fluoromica was performed by ion exchange using ODTMA (single surfactant) and 75% ODTMA/25% choline chloride (CC) (dual surfactants) via a previously published method.25,26 These organically modified fluoromica (organofluoromica) were used as nanofillers in the TPU.
Based on the ZP model, the scattering from two-phase systems can be represented as the product of the form factor, P(q), and the structure factor, S(q): 27–30
I(q) = AP(q) S(q) | (1) |
(2) |
S(q) represents the interference effects between X-rays scattered by different scattering bodies in the sample and depends on their relative positions and can be described by:27,30
(3) |
where
(4) |
(5) |
where
(6) |
In this case, f was calculated from the azimuthally averaged data, where 〈cos2Φ 〉 is the average cosine squared weighted by intensity I as a function of the radial angle, Φ. The value of f is equal to 1 and −0.5 when the orientation of the nanofillers is completely aligned perpendicular and parallel to the direction of strain, respectively, and is zero for random (isotropically) orientated nanofillers.
Fig. 1 TEM micrographs of (A&B) 2MEO, (C&D) 4MEO, (E–G) 2MEO–C and (G-H) 4MEO–C. |
The XRD signature of any polymer-silicate nanocomposite may be influenced by the average platelet size, degree of nanosilicate inter-platelet registration, orientation and increase in basal spacing of the nanosilicate due to intercalation by the host polymer. The increase in basal spacing depends on the amount of TPU intercalated in the galleries of the silicates,34–37 and, of course, is a function of nanosilicate loading, nanosilicate-TPU interactions and also nanocomposite processing history. XRD patterns of the ME (unmodified, O and O–C modified) are shown in Fig. 2a, and XRD patterns for their respective nanocomposites are shown in Fig. 2b. The modification of ME with both O and O–C substantially increases the basal spacing (d001) from 1.3 nm to 4.9 nm. The MEO and MEO–C organosilicates exhibited several well-defined diffraction peaks as compared to the pristine ME, which correspond to basal spacings of approximately 4.9 nm, 2.3 nm and 1.6 nm. The observed peaks are caused by the interstratified superstructure of the fluoromica, which originates from the charge heterogeneity.15,38 This charge heterogeneity may allow different amounts of surfactant chains to intercalate the fluoromica layers, and hence different surfactant arrangements are possible.15 The high aspect ratio fluoromica gives rise to long-range order in the resulting organosilicate tactoids (with respect to synthetic hectorites or natural montmorillonite nanosilicates, for example), as well as strong diffraction peaks. The XRD profile of the neat host TPU shows no significant peak, which is expected, as both soft and hard domains present within a PDMS-based TPU do not show any diffraction peaks between 2θ = 0.5° to 10°39,40 and this TPU segmental microphase periodicity is expected to be observed at larger spacings.24 Both MEO and MEO–C nanocomposites, regardless of the organosilicate wt% loading, exhibited a series of (00l) reflections with basal spacings of approximately 4.0 nm, 2.0 nm and 1.4 nm.
Fig. 2 XRD pattern of the a) pristine ME and after ODTMA (O) and ODTMA/CC (O–C) modification and b) the E5-325 TPU containing 2 and 4 wt% MEO and MEO–C. |
Fig. 3 Effect of nanofiller types and loadings on the stress–strain curve of E5-325 TPU. |
Material | Tensile strength/MPa | Young's modulus/MPa | Elongation at break (%) | Toughness/MPa | Tear strength/MPa | Tensile creep modulus (Et)/MPa |
---|---|---|---|---|---|---|
Nusil MED 4860 | 7.5 ± 0.1 | 2.7 ± 0.2 | 595 ± 39 | 35 ± 4 | 41 ± 14 | 1.51 ± 0.04 |
E5-325 | 20 ± 2 | 9.7 ± 0.7 | 941 ± 89 | 108 ± 17 | 60 ± 8 | 2.4 ± 0.1 |
2MEO | 20 ± 2 | 19.9 ± 0.7 | 847 ± 39 | 106 ± 10 | 65 ± 2 | 12.1 ± 0.1 |
4MEO | 17 ± 2 | 23.7 ± 3.9 | 704 ± 114 | 83 ± 16 | 71 ± 4 | 8.02 ± 0.02 |
2MEO–C | 23 ± 1 | 11.3 ± 0.5 | 998 ± 51 | 129 ± 13 | 68 ± 4 | 7.9 ± 2 |
4MEO–C | 23 ± 2 | 11.9 ± 0.6 | 1003 ± 61 | 129 ± 12 | 75 ± 7 | 6.8 ± 0.5 |
The stress–strain curve of the neat E5-325 is similar to that of other TPUs.11,37,41 It can be manifested as the evolution of the phase-separated soft and hard TPU segments during the deformation process. At low strains (0–50%), the linear region is due to the stretching of the elastic soft segments, in which the hard segments respond to this soft segment alignment. At high strains, the non-linear viscoelastic region is due to the more subtle TPU microstructural changes such as the partial extent of soft segment chains, and the hard domains orientation, rotation, deformation and fragmentation.24,42 This behaviour can be influenced by the nanofiller addition, depending on the degree of TPU-nanofiller interactions.24,28 In agreement, our results show that the incorporation of nanofillers with different functionality greatly influences the stress–strain behaviour of the E5-325 TPU. These morphological changes observed during the deformation process will be further described in the in situ strained SAXS section.
The following comparisons were made based on the data that shows a statistically significant difference. E5-325 TPU displays greater tensile properties when compared with Nusil MED4860, showing gains of 167% in tensile strength, 63% in elongation at break, 220% in toughness and 56% in tear strength. Adding dual modified ME (MEO–C) further increases the mechanical properties of this PDMS-based TPU. The best mechanical properties were achieved when 4 wt% MEO–C was added, giving rise to an increase of 15% in tensile strength, 19% in toughness and 25% in tear strength. Significantly, the use of dual surfactants resulted in enhancements in strength and toughness without the substantive accompanying increases in elastic modulus or stress at 100% strain observed when a single surfactant was employed. This suggests that the utilization of dual surfactants enables preferential interactions with both soft and hard segments.26 Furthermore, the aforementioned high temperature trans-urethanisation reactions and subtle molecular reorganization at the TPU-nanofiller interface lead to the formation of more lengthy aggregated arrays (Fig. 1E–H). This morphology and enhanced hard segment interaction is shown later in this communication to generate more efficient platelet orientation when subjected to tensile deformation (Fig. 12). There is also the possibility that the new population of choline-derived positive TPU end groups discussed previously provide labile, but mobile interactions with the nanofiller, which can be broken and re-formed, thereby retaining toughness without stiffening.
In contrast, the addition of MEO resulted in a drastic increase in Young's modulus, accompanied by a reduction in tensile strength and elongation at break. With the addition of 4 wt% MEO, the Young's modulus was substantially increased by ∼144%. In the absence of the additional CC surfactant, the hydrophobic MEO, when added at 2 and 4 wt%, dispersed and delaminated well in the TPU matrix (Fig. 3b), but tended to interact preferentially with the soft segments, causing a reduction in soft segment mobility, and a substantial increase in Young's modulus. As a result of this exfoliated morphology and better dispersed MEO, this nanocomposite system appeared to promote more “particle–particle” interactions during elastomer deformation rather than “particle-TPU interactions”, which can hinder both the TPU chain and nanofiller mobility.
Tensile creep and stress relaxation testing of selected samples was performed in order to measure their time-dependent dimensional stability under tensile deformation. The tensile-creep modulus (Et) was used to evaluate the creep resistance of the TPU and the single and dual modified nanocomposites. Et is the ratio of the applied stress to tensile-creep strain reached at total creep. Et values for Nusil MED 4860, E5-325 TPU, 4MEO and 4MEO–C measured at a stress of 2 MPa are shown in Table 1, while their representative tensile creep curves are displayed in Fig. 4. The tensile creep modulus of the E5-325 TPU is 59% greater than Nusil MED 4860. Adding 2 and 4 wt% MEO and MEO–C profoundly improves the creep resistance of this PDMS-based TPU. The highest Et was achieved by 2MEO with an increase in creep resistance of 404%, while 2MEO–C resulted in an increase of 229%. It appears that, at this relatively low constant tensile load of 2 MPa, the stiffer MEO system, with its greater affinity towards the PDMS soft segments and higher degree of nanofiller particle–particle interactions, has most successfully restrained the soft segment mobility and hence reduced the TPU extension under these conditions.
Fig. 4 Tensile-creep curves of Nusil MED 4860, E5-325, 4MEO and 4MEO–C at an applied stress of 2 MPa over a test period of 6 h. |
Stress relaxation data obtained at 50% strain is presented in Fig. 5. The data was normalized against the stress, σ(t′), at t = 5 s, and demonstrated a power-law dependence (σ(t)/σ(t′) = At−b as displayed in Fig. 5b. The fitting constants, A and b, from the power law analysis are provided in the inserted table in Fig. 5. Although in the initial stages the stiffer MEO system retains a higher stress, the fitted slopes of the curves reveal that the stress relaxation rate is reduced with the addition of MEO–C nanofiller, while the rate actually increased with the MEO inclusion. This clearly shows that the dual modified organosilicate (MEO–C) most successfully retarded the stress relaxation of the TPU. According to Sternstein and Zhu,43 at low strain, the stress relaxation rate in the nanocomposites is considered to be the result of strain-induced disentanglements and slippage of chain segments at the filler surface. Therefore, it can be inferred that the dual modified ME with a more preferential interaction with both TPU soft and hard segments is most effectively contributing to reduced strain induced molecular or segmental slippage, and hence a lower stress relaxation rate. The positive end groups and proposed “molecular bridging” that formed in the MEO–C nanocomposites might also serve to inhibit the molecular or segmental slippage during this relaxation under constant small strain. Conversely, poor organosilicate–hard segment interactions in the MEO nanocomposite resulted in greater segmental chain slippage at the interface. In addition, the enhanced particle–particle interactions in the MEO system placed polymer chains at the nanofiller-TPU interphase under increasing stress, inducing an increased rate of stress relaxation by promoting strain induced slippage.43,44
Fig. 5 Stress relaxation data obtained at 50% strain. |
To examine the possible reasons for these changes in TPU mechanical behavior with the organosilicate addition, we performed the DMTA, DSC and strained synchrotron SAXS analysis to provide information on the possible morphological changes and specific TPU-nanofiller molecular interactions.
Fig. 6 DMTA data as a function of temperature: a) damping factor (tanδ); b) storage modulus (E′) for E5-325 (neat host TPU and nanocomposites). |
We observe Tα1 in the −86 to −91 °C range, whereas Tα2 is observed in the −3 to 2 °C range. The Tα2 of the TPU increased slightly with the addition of the MEO–C and was associated with maintenance of the damping peak height, while Tα2 decreased with the addition of MEO, and peak height was reduced.
The storage modulus was observed to increase modestly with the addition of both 2 wt% and 4 wt% MEO–C. This is consistent with the tensile and creep results discussed previously. However, an even higher storage modulus was observable in the E5-325 with MEO. In agreement with our tensile test results, 2 and 4 wt% MEO more efficiently increased the modulus of E5-325 at room temperature with respect to the MEO–C counterparts. Interestingly, this largest increase in both Young's and storage moduli was achieved in conjunction with a reduction in soft microphase Tg, a phenomenon not observed with the MEO–C nanocomposites. This furthermore supports the notion of more specific soft segment interactions in the MEO system, including perhaps some plasticization of the PDMS-rich microphase with ODTMA alkyl tails.
Fig. 7 Typical DSC a) heating curves and b) cooling curves for the neat host E5-325 TPU and E5-325 TPU containing MEO and MEO–C. |
The MEO nanocomposites exhibit a slightly higher T1 temperature as compared to the neat TPU and MEO–C nanocomposites. Due to the very hydrophobic inter-gallery space in this system, which stretches over a much larger length scale than the typical TPU domain texture, one can imagine soft segments, single MDI residues and shorter hard segments being preferentially intercalated. This intercalated shorter hard segment population may be contributing to the increase in T1. A bimodal T2 endotherm in E5-325 with MEO–C can be attributed to the disruption of various degrees of short-range hard segment order, due to distribution in hard segment lengths and morphologies. A broad exotherm is observable immediately after the T2 endotherms in the TPU containing MEO–C, and this is probably related to conformational changes taking place in the hard phase. This new signature observed in both 2MEO–C and 4MEO–C systems supports our hypothesis that there is a potential covalent re-coupling of the single MDI and perhaps shorter hard segments (via trans-urethanisation) with –OH groups on CC, introducing labile positive end groups and subtle morphological changes.
Based on the cooling scan, crystallization exotherm peaks at 87 to 88 °C were observable for the E5-325 containing 2MEO–C and 4MEO–C respectively, meanwhile, for 2MEO and 4MEO, the crystallization exotherms occured at higher temperatures (90 and 92 °C, respectively). One possible reason for this is that because longer hard segments are effectively excluded from intercalating the MEO galleries, they are encouraged to nucleate sooner than in the MEO–C system where hard segment-nanofiller interactions and reactions are promoted.
Fig. 8 2D SAXS patterns at selected strains for E5-325 (neat host TPU and nanocomposites) obtained from the short sample-to-detector distance. The relaxed state refers to images taken 10 min after strain measurements were taken. |
Fig. 9 1D profiles of E5-325, 2MEO, 4MEO, 2MEO–C and 4MEO–C at 0%, 100% and 400% strain in the strain and transverse directions, obtained from the short sample-to-detector distance. |
In the initial unstrained state, all samples show isotropic SAXS patterns, indicating that the hard segment domains are randomly oriented, which then transforms to increasingly anisotropic patterns upon stretching. This TPU SAXS pattern is similar to that discussed previously.24 At 100% strain, the ring deforms to an ellipsoid with the long axis along the equator, attributed to an affine deformation of the two phase structures of TPU.24 The soft segments are mainly involved in the deformation of this material during the first 100% strain and TPU hard segments respond to the alignment of the soft segment chains.24,28 The SAXS pattern reveals scattering lobes on the meridian and a streak in the equatorial direction when the strain reaches 400%. As discussed previously,24 the equatorial streak appears as a result of reduced electron density contrast between the hard segment nanofibrils and the aligned soft segment chains and/or the decrease in coherent scattering due to the small size of the broken down hard domains. Conversely, the SAXS patterns for MEO and MEO–C nanocomposites demonstrate an isotropic outer ring prior to straining, which then progressively transforms into two arcs in the equator upon straining. This outer ring, is located in the high q region, and relates to the diffraction from registered high aspect ratio fluoromica tactoids, and is understandably more intense in MEO–C nanocomposites. In agreement with the TEM images, MEO–C nanocomposites exhibit larger, thicker elongated tactoids as compared to MEO nanocomposites. At 400% strain, the two arcs observed in the region of q = 0.160 to 0.165 Å−1 represent the oriented nanoparticles, and the respective scattering peak at q = 0.16 Å−1 can be observed in the SAXS profile (Fig. 9). Fig. 8 and 9 demonstrate the lower scattering intensity of 4MEO than that of the 4MEO–C nanocomposite in the q region between 0.02 and 0.15 Å−1. This lower scattering intensity, which is most notably seen at 400% strain, might be due to higher fractions of broken down hard segment lamellae in the 4MEO system. This greater fragmentation of hard segments in the 4MEO upon deformation could be the result of weaker hard segments interactions with the hydrophobic MEO nanofiller, particularly at the higher 4 wt% loading. In addition, the interparticle locking resulting from the superior MEO exfoliation also might induce frustrated platelet alignment, thus introducing more interphase void formation as illustrated in Fig. 12c and d. (We observed more substantial strained specimen whitening in the MEO nanocomposite system.)
For all strained samples, slightly intensified scattering was observable from SAXS patterns during the relaxation process, indicating a slight recovery of phase separation and hard segment ordering. The nanocomposite samples show a lesser degree of recovery due to reduced TPU microphase mobility, as a result of organosilicate inclusion.
Prior to straining, the 1D SAXS profiles presented in Fig. 9 for the neat host TPU and nanocomposites reveal prominent peaks at ∼q = 0.07 Å−1 corresponding to the segmental microphase periodicity in the materials. Further stretching results in the decrease in the scattering intensity, both in strain and transverse directions, which is due to a disruption of the hard segment domain fraction. Another reflection at ∼q = 0.16 Å−1, corresponding to diffraction from the fluoromica tactoids with basal spacings of 4.0 nm was observable in the undeformed MEO and MEO–C nanocomposite samples. The scattering peaks appear stronger and shift slightly towards higher q in the transverse direction as the strain increases from 0% to 400%, suggesting greater organosilicate alignment and reduced spacing between the individual platelets. The XRD analysis of the unstrained 2MEO and 4MEO, as shown in Fig. 2b, also confirms the presence of a similar organosilicate basal spacing. The change in hard domain spacing with deformation for E5-325 TPU and nanocomposites was also studied from the SAXS patterns. SAXS data averaged over 10° segments in the strain and transverse directions was analyzed using the ZP model, which was previously used to successfully fit SAXS data from E5-325 TPUs subjected to uniaxial deformation.24 The morphological data obtained from the fitting, combined with the direct visualization of the 1D scattering profiles is tabulated in Table 2 and presented graphically in Fig. 10.
Fig. 10 Estimated average hard domain spacing (d) of neat host TPU and nanocomposites determined from the ZP model in the strain and transverse direction. |
Strain (%) | Material | Strain Direction | Transverse Direction | ||||
---|---|---|---|---|---|---|---|
d/nm | R/nm | σ/d | d/nm | R/nm | σ/d | ||
a Expected uncertainty due to sample variations and curve fitting. | |||||||
Uncertaintya | ±0.5 | ±0.5 | ±0.05 | ±0.5 | ±0.5 | ±0.05 | |
0 | E5-325 | 8.9 | 2.8 | 0.44 | 8.9 | 2.7 | 0.44 |
2MEO | 8.5 | 2.8 | 0.47 | 8.4 | 2.8 | 0.54 | |
4MEO | 8.4 | 2.9 | 0.64 | 8.3 | 2.9 | 0.50 | |
2MEO–C | 8.7 | 2.8 | 0.45 | 8.4 | 2.8 | 0.47 | |
4MEO–C | 8.4 | 2.8 | 0.46 | 8.5 | 2.8 | 0.54 | |
50 | E5-325 | 11.5 | 2.9 | 0.53 | 7.6 | 2.7 | 0.42 |
2MEO | 10.9 | 2.9 | 0.63 | 7.6 | 2.8 | 0.53 | |
4MEO | 9.8 | 2.9 | 0.64 | 8.0 | 2.9 | 0.53 | |
2MEO–C | 10.5 | 2.9 | 0.68 | 7.7 | 2.8 | 0.49 | |
4MEO–C | 10.8 | 2.9 | 0.65 | 7.7 | 2.7 | 0.52 | |
100 | E5-325 | 13.8 | 3.0 | 0.64 | 7.0 | 2.7 | 0.43 |
2MEO | 12.6 | 2.9 | 0.90 | 7.1 | 2.8 | 0.56 | |
4MEO | 10.8 | 2.9 | 0.80 | 7.6 | 2.8 | 0.58 | |
2MEO–C | 10.1 | 2.9 | 0.78 | 7.2 | 2.7 | 0.54 | |
4MEO–C | 12.1 | 2.8 | 0.99 | 7.1 | 2.7 | 0.55 | |
200 | E5-325 | 13.2 | 2.6 | 0.94 | 6.5 | 2.6 | 0.43 |
2MEO | 10.7 | 2.7 | 0.89 | 6.4 | 2.6 | 0.62 | |
4MEO | 9.8 | 2.7 | 0.84 | 7.4 | 2.5 | 0.90 | |
2MEO–C | 10.1 | 2.7 | 0.84 | 6.8 | 2.5 | 0.67 | |
4MEO–C | 10.3 | 2.7 | 0.83 | 6.3 | 2.5 | 0.61 | |
400 | E5-325 | 9.9 | 2.5 | 0.73 | 5.2 | 2.5 | 0.42 |
2MEO | 9.8 | 2.4 | 0.94 | — | — | — | |
4MEO | 9.0 | 2.5 | 0.95 | — | — | — | |
2MEO–C | 9.2 | 2.5 | 0.83 | — | — | — | |
4MEO–C | 10.0 | 2.5 | 0.82 | — | — | — |
For the neat TPU, the average hard domain spacing (d) is observed to increase sharply in the strain direction at the initial stages of deformation. The randomly oriented hard domains subjected to tensile deformation, exhibited an increase of d-spacing from 8.9 nm to 13.8 nm upon 100% strain as the TPU hard blocks respond to the alignment of the soft segment chains.24 Subsequent deformation up to 400% strain resulted in a decrease in the hard domain spacing to 9.9 nm, as the hard segment aggregates were disrupted, broken down to smaller widths, and partially aligned in the direction of stretch. Further stretching resulted in no further change to the d spacing of the neat host TPU. Conversely, straining all of the TPU nanocomposites up to 200% is associated with smaller shifts of d in the strain direction, most notably seen in the stiffest 4MEO system. On the basis of TEM analysis, MEO demonstrated the best nanofiller dispersion and exfoliation. At higher loading of 4 wt% MEO, the “interparticle locking” resulting from the MEO exfoliation restricted the soft segment alignment under straining, as well as the hard segment mobility, and increased the modulus of the TPU. This led to a better retention of hard domain spacing upon deformation. It is also interesting to note that 2MEO–C system demonstrated an overall better retention of hard domain spacing than the 2MEO system, and it displayed the smallest d at 100% strain. This is particularly impressive given the overall property profile of this material. This is due to the positive end groups introducing a ‘grip-slip-grip’ interaction between the TPU segmental chain and the nanofiller, preventing the hard domains spacing further apart.
For all samples, a decrease in d is observed in the transverse direction with increasing strain. As expected, 4MEO, which is the stiffest material and involves the highest degree of interparticle interactions, showed the most substantial retention of d in the transverse direction. The ZP model was unable to fit the scattering data of high strained MEO nanocomposites in the transverse direction because the scattering peak had diminished. In all cases, the hard segment radius of gyration (R) did not change significantly over the entire strain range.
Fig. 11 Selected 2D SAXS patterns at various strains for E5-325, 4MEO and 4MEO–C, obtained from the long sample-to-detector distance. |
In the initial unstrained state, all samples show isotropic SAXS patterns. For the neat TPU, the absence of nanofiller led to slightly anisotropic SAXS patterns upon straining up to 400%. In comparison, both MEO and MEO–C nanocomposites demonstrate different scattering geometry upon straining. A diamond-shaped pattern suggests that there was alignment of fluoromica tactoids in both the strain and transverse direction, presumably due to delamination of matrix and nanofiller in some instances, and effective load transfer and tactoid orientation in others.24,28 However, MEO–C nanocomposites exhibit a slightly sharper silicate scattering along the equatorial axis, indicating greater platelet alignment in the strain direction. TEM images displayed in Fig. 12 were taken after a few days of unloading from 400% strain. It is believed that tactoids in 4MEO–C act as well-bonded and moveable “concertina-like” ensembles, which are aligned in the direction of the strain during straining. However, after unloading, these tactoids undergo a reduction in alignment in the strain direction and rotated in ∼45° from the strain direction, due to the TPU relaxation (Fig. 12a&b). Meanwhile, the single modified fluoromica system, which is associated with a more exfoliated and interlocked nanofiller dispersion, certainly shows concrete evidence of more frustrated platelet orientation. Tensile stresses built at the interface of poorly aligned tactoids are believed to contribute to void formation (Fig. 12c) that was also evidenced via stress whitening (figure not shown). This nanocomposite morphology during loading and unloading was previously described by Finnigan et al. in the TPU-organofluoromica system.44
Fig. 12 TEM micrographs of 4MEO–C (a & b) and 4MEO (c & d), which were taken after the TPU was relaxed after stretching to 400% strain. The arrows indicate the pre-strain direction. The dual modified tactoids in 4MEO–C are seen to ensemble in a “concertina-like” arrangement. 4MEO shows void formation due to interparticle locking and interphase failure. |
The Herman orientation function (f) was calculated at q = 0.0031, where the scattering is dominated by the organosilicates. These values were illustrated in Fig. 13. The degree of platelets alignment, especially at high strains (>50%) was significantly greater for 4MEO–C than for 4MEO nanocomposites. This suggests that the dual modified fluoromica aligns more preferentially in the strain direction than the single modified counterpart at a similar concentration during uniaxial tension of TPU, and consequently provides greater improvement in toughness. This was also supported by the TEM analysis displayed in Fig. 12. This further elucidates why the presence of MEO in the TPU led to a more pronounced stiffening effect and a reduction in tensile strength and elongation at break.
Fig. 13 Herman orientation functions versus strains for 4MEO and 4MEO–C nanocomposites. |
This journal is © The Royal Society of Chemistry 2012 |