Selective Mechanical Reinforcement of Thermoplastic Polyurethane

N. Coleman*. School of Physics, Trinity College Dublin, Dublin 2, Ireland. J. Phys. Chem. C , 2010, 114 (26), pp 11401–11408. DOI: 10.1021/jp102...
7 downloads 12 Views 1MB Size
Download PDF
J. Phys. Chem. C 2010, 114, 11401–11408

11401

Selective Mechanical Reinforcement of Thermoplastic Polyurethane by Targeted Insertion of Functionalized SWCNTs Umar Khan, Fiona M. Blighe, and Jonathan N. Coleman* School of Physics, Trinity College Dublin, Dublin 2, Ireland ReceiVed: April 1, 2010; ReVised Manuscript ReceiVed: June 2, 2010

We have prepared composites from a thermoplastic polyurethane reinforced with functionalized single walled nanotubes. Nanotubes with two types of functional groups were used: water-soluble tubes functionalized with polyethyleneglycol or poly(amino benzene sulfonic acid) and tetrahydrafuran-soluble tubes functionalized with octadecylamine. Composites prepared with water- or tetrahydrafuran-soluble tubes showed markedly different properties. Addition of water-soluble tubes tended to result in crystallization of the polyurethane soft segments, whereas addition of the tetrahydrafuran-soluble tubes promoted crystallization of the polyurethane hard segments. We interpret this as evidence of selective insertion of tubes in either hard or soft segments depending on the surface chemistry of the (functionalized) nanotube and the chemical structure of the segment. This interpretation is supported by differences in the mechanical properties of the composites. The waterbased composites tend to be stiffer and display higher plateau stress, consistent with reinforcement of the soft segments. However, the tetrahydrafuran cast composites tend to maintain their strength and ductility at higher nanotube loading levels, whereas the water-based composites become weak and brittle above ∼10 vol % nanotubes. This is consistent with the water-based nanotubes impeding the extension and motion of the soft segments, resulting in loss of ductility. In contrast, the tetrahydrafuran-soluble nanotubes become segregated in the hard segments and so do not negatively impact on the mechanical properties at high nanotube content. This controlled reinforcement has allowed us to prepare composites with modulus, plateau stress, strength, and ductility of up to 250 MPa, 8 MPa, 60 MPa and 750%, respectively, significantly better than neat polyurethane. Introduction Carbon nanotubes (CNTs) have excellent mechanical, electrical, and thermal properties,1-7 generating much interest into their use as fillers in polymer-based composites.8 Much work has focused on utilizing them as reinforcements to mechanically enhance polymers.9-12 One important parameter in composite reinforcement is good dispersion of the filler.7 Pristine nanotubes can be dispersed efficiently in solvents at very low concentration but at higher concentrations they tend form bundles.13,14 These bundles are held together by weak physical interactions and therefore are less effective fillers than individual nanotubes. However this problem can be overcome by the use of functionalized CNTs. Very good dispersions of highly exfoliated, functionalized CNTs can be achieved at very high concentrations compared with pristine CNTs.15-18 Another requirement for good mechanical reinforcement is good stress transfer leading to strong interfacial bonding.7 Excellent interfacial bonding can also be achieved by using functionalized CNTs. For nanotubes functionalized with polymer strands, we would expect the grafted chains to entangle with the matrix polymer (assuming chemical compatibility) creating a very strong interface. Lower molecular weight functional groups will bond with the polymer matrix via van der Waals interactions. Thus, as they address both dispersion and stress transfer requirements, functionalized nanotubes are ideal for mechanical reinforcement applications. However, we note that the interaction with both the solvent and the matrix polymer will be very sensitive to the nature of the functional group. As such, great care must be taken when * To whom correspondence should be addressed. E-mail: [email protected].

choosing the functional group (or grafted polymer chain). However, as we will show, the correct choice of functional group can allow targeted insertion of the nanotubes into specific regions of the matrix. This can allow fine control over the properties of the final composite. Much of the work to date on polymer-nanotube composites for mechanical applications has used rigid thermoplastic polymers such as polyvinylalcohol19-21 or polystyrene22 as the matrix. One problem with mechanical reinforcement of such composites is that the composite tends to lose ductility and toughness at high volume fraction.23 It has been suggested that an alternative approach would be to load nanotubes into elastomers.23 To allow solution processing, it would be advantageous to use thermoplastic elastomers such as thermoplastic polyurethane. Thermoplastic elastomers are segmented copolymers where certain segments tend to form interchain aggregates. These aggregates link together separate chains and so act as noncovalent cross-links. We propose that insertion of nanotubes at high enough loading level could increase the stiffness and yield strength of optimized elastomer-nanotube composites toward that of rigid thermoplastics. However, if even a fraction of the ductility of the original elastomer could be retained, one might have a composite significantly tougher than composites based on brittle thermoplastics. However to achieve this, it will be critical to understand how to minimize the reduction in ductility, while maximizing the increase in stiffness on adding nanotubes. Among elastomers, thermoplastic polyurethane (TPU) is a versatile matrix as it possess very high ductility, good stress-strain recovery, and good strength.24 Unlike regular rigid

10.1021/jp102938q  2010 American Chemical Society Published on Web 06/16/2010

11402

J. Phys. Chem. C, Vol. 114, No. 26, 2010

thermoplastics, when TPU is stretched beyond the elastic limit, significant permanent deformation does not occur. On release of the applied strain, TPU tends to recover to close to its original dimensions with little permanent set.25 This shape recovery is due the presence of phase-separated7 hard segments (HS) which act as physical cross-links for the continuous phase of soft segments (SS). The problem with TPU is that it demonstrates low stiffness and low stress in the plateau region of its stress strain curve. A number of papers have shown that TPU can be reinforced by CNTs.26-30 However, this work has generally used unfunctionalized SWCNTs which are not designed to interact with the matrix in any organized way. However, the phase separation of hard segments and soft segments in TPU allows one to consider targeted reinforcement of either HS or SS phase. The HS and SS generally have different hydrophobicities.31-34 This strongly suggests that by selecting appropriate functionalized CNTs, either HS or SS can be independently reinforced for a specific application. Recently Liff et al. have demonstrated the fabrication of nanoclay-TPU composites by restricting the nano filler to the hard segment only.32 This suggests that both the hard and soft segment’s mechanical properties can be tuned by the addition of appropriately functionalized CNTs. In the present work, keeping in mind the expected difference in hydrophobicities of HS and SS, three different types of functionalized-SWCNTs were used in two different solvents, water and tetrahydrafuran (THF), to fabricate composites. All three composites exhibited different properties. Water-based samples show higher stiffness at all volume fractions (Vf) and good strength and ductility at low Vf but suffer a sharp reduction in strength and ductility at high Vf. However, while THF-based samples also exhibit increases in stiffness, strength, and ductility at low CNTs Vf, these samples do not show the sharp decline in mechanical properties observed for water-based samples. These results are compatible with the hypothesis that the watersoluble nanotubes interact predominately with the soft segments while the THF soluble nanotubes interact with the hard segments. These composites have potential application for use as actuators/artificial muscles,24,35 extremely flexible electrodes,36 and sensors.37 Results and Discussion We have studied the thermal and mechanical properties of thermoplastic polyurethane filled with three different types of single walled nanotubes (see experimental). Each nanotube type was functionalized with a different group: polyethylenegycol (PEG-SWCNTs), poly(amino benzene sulfonic acid) (PABSSWCNTs), and octadecylamine (ODA-SWCNTs). The first two tube types are hydrophilic and are soluble in water, whereas the latter type is hydrophobic and is soluble in organic solvents, notably THF. The water-soluble nanotubes were blended with a dispersion of surfactant stabilized polyurethane in water while the ODA-SWCNTs were mixed with dissolved PU in THF. Films were made by drop casting, dried, and cut into the appropriate shape for testing. Differential Scanning Calorimetry. DSC curves for a subset of the samples measured are shown in Figure 1. We note that the water-based samples were dried vigorously to ensure complete solvent removal (as confirmed by thermogravimetry). In all cases, the glass transition was very broad and centered ∼-25 °C (the glass transition is hard to see in figure 1 and is best observed by plotting the derivative of the DSC trace). No real difference was observed between the Tg of the composite and the polymer-only samples. This suggests that the function-

Khan et al.

Figure 1. DSC curves for TPU and functionalized SWCNT/TPU composites. Tm1 is soft segment melting peak and Tm2 corresponds to HS melting nucleated by ODA-SWCNTs.

alized nanotubes do not act as plasticizers. In addition, a crystalline melting peak was observed at Tm1 ∼ 41 °C, corresponding to the TPU soft segment melting. SS melting in this temperature range has been reported by a number of other researchers.24,28,30,31,38-40 The melting enthalpies were very similar for the ODA-SWCNT composites and the TPU/THF polymer-only sample (8.1 J/g for TPU/THF versus 8.2 J/g for 1.55 vol % ODA-SWCNTs). This suggests that the ODASWCNTs do not influence the degree of soft segment crystallization. However, the melting enthalpies of the water dispersed PABS-SWCNT and PEG-SWCNT composites were consistently higher than the TPU/H2O polymer-only sample. For example, we found melting enthalpies of 2.2 J/g for the dried water-based TPU, 3.3 J/g for the dried PEG-SWCNT/TPU composite (Vf ) 1.55%) and 4.8 J/g for the dried PABS-SWCNT/TPU composite (Vf ) 1.55%). This data shows that the soft segment melting enthalpy (Tm1) is significantly enhanced by addition of both PEG- and PABS-SWCNTs but not by the ODA-SWCNTs. We suggest that the water-based nanotube systems tend to nucleate SS polymer crystallinity in a manner impossible for the ODASWCNT. One explanation for this behavior would be the presence of a preferential interaction between the water dispersed CNTs and the SS. The DSC study also revealed a second endothermic peak (Tm2) at 114 °C in ODA-SWCNT composites, whereas no such peak was observed in water-based composites. This second melting peak was only observed for volume fractions above 1.55 vol %. The melting enthalpy (Figure 2) associated with this peak scales linearly with nanotube content, suggesting crystallites nucleated by the presence of nanotubes.41 Nucleation effects of CNTs in polymers have been reported in a number of papers.42-44 Liff et al.32 observed a DSC peak for TPU/nanotube samples in the same temperature range, which they associated with TPU hard segment melting. We suggest that this new peak is due to the melting of nanotube-nucleated HS crystallites. Again, this may be explained by a preferential interaction between ODASWCNTs and the HS of TPU.

Reinforcement of Thermoplastic Polyurethane

J. Phys. Chem. C, Vol. 114, No. 26, 2010 11403

Figure 2. Differential scanning calorimetry data showing a linear increase in melting enthalpies of TPU hard segment (Tm2) with ODASWCNT content.

Figure 4. Dynamic mechanical measurement of tan δ as a function of temperature for TPU and TPU filled with functionalized nanotubes. This data shows shifting of Tg to higher temperature with the addition of water dispersible CNTs. This indicates SS interaction with these CNTs.

Figure 3. Dynamic mechanical measurement of the storage modulus of TPU and TPU filled with functionalized nanotubes as a function of temperature.

We note that these DSC experiments immediately suggest a preferable interaction between water-soluble PEG-SWCNTs and PABS-SWCNTs and the TPU soft segments. Conversely, the data for the ODA-SWCNTs is compatible with a preferable interaction between the nanotubes and the TPU hard segments. This implies that, for the TPU used here, the SS are hydrophilic, whereas the HS are hydrophobic. (We note that for the TPU used here the manufacturer declined to provide detailed information on the chemical nature of HS and SS.) This is an interesting result as it suggests that control of the functional group allows one to target certain parts of the polymer for nanotube insertion. Over the course of the rest of this paper we will consider alternative evidence for this hypothesis. Dynamic Mechanical Analysis. Dynamical mechanical analysis was also carried out with the results shown in Figures 3 and 4. In general, most composite samples show increased storage modulus compared to their respective polymer only samples as shown in Figure 3. At room temperature, all composites (except ODA-SWCNT, 0.1 vol %) showed storage moduli higher than the polymer-only samples. For PABSSWCNT composites (3.1 vol %), the storage modulus increased from 41 to 111 MPa, whereas for PEG-SWCNT composites, it

reached 102 MPa. This indicates good interaction and stress transfer between the CNTs and the polymer at room temperature. For the 3.1 vol % ODA-SWCNT composite, the storage modulus was increased from 79 to 112 MPa (compared to polymer). For all composites, the storage modulus for the 3.1 vol % samples remains high compared to the polymer-only at all temperatures above room temperature. This shows that addition of nanotubes can slightly increase the usable temperature range of these samples.32 Note that the difference between the TPU/THF and TPU/water moduli will be discussed later. Shown in Figure 4 is a plot of tan δ as a function of temperature. This quantity is the ratio of the loss modulus to the storage modulus, and is a measure of the fractional energy loss per cycle. The peak in tan δ coincides with the dynamic glass transition. In general, glass transition temperatures measured using DMTA are usually higher than those measured by DSC. We observe such a difference here with Tg ∼36 °C higher than that measured by DSC. Such a large difference between Tg’s measured by DMTA and DSC has been reported for other thermoplastic polyurethanes.45 For both TPU/THF and TPU/H2O samples, the dynamic Tg was measured to be ∼11 °C (compared to -25 °C as measured by DSC). For the ODA-SWCNT samples, the Tg was unaffected by the introduction of CNTs. However, for the PEG-SWCNT and PABS-SWCNT composites, Tg increased by 6-14 °C compared to TPU. We understand these trends as follows. The glass transition is the temperature above which the SS start to become mobile. The results discussed above suggest the water dispersed CNTs to physically bond to the SS, thus restricting their mobility. This results in an increase in the Tg of the system.46-48 On the other hand, the ODA-SWCNTs appear to show only limited interaction with the SS as their presence does not influence Tg. This would be expected if the ODA-SWCNTs preferentially interact with the HS. This data reinforces the

11404

J. Phys. Chem. C, Vol. 114, No. 26, 2010

Khan et al.

Figure 5. Representative stress-strain curves of TPU and its composites made with functionalized SWCNTs.

hypothesis that the water dispersed CNTs interact with the SS, whereas the ODA-SWCNTs interact preferentially with the HS. Tensile Mechanical Testing. Shown in Figure 5 are the representative stress-strain curves for the TPU and associated composites, which can be divided into three parts.32,39 At low strain the curves are approximately linear with slope equal to the Young’s modulus. At higher strains, the curves flatten to give a plateau region which is associated with the continued stretching of the SS and the rotation and alignment of the HS segments. The plateau region is followed by a high-slope region which we associate with strain hardening (SH) and strain induced crystallization (SIC). At high strain values, we observe failure of the sample. In general, addition of nanotubes tends to increase the Young’s modulus and increase the stress for all strains. We also observed some loss in ductility for higher nanotube content samples. For all three CNT types, the Young’s moduli of the composites increased significantly compared to the unfilled polymer. The polymer-only samples displayed moduli of 12 and 6.3 MPa for the water and THF cast systems, respectively. We note that these are likely due to morphological differences between these systems. For example, PU dissolves in THF, whereas in water it is dispersed as globules. This will result in morphological differences in the final films. In addition, the different sample preparation techniques may result in different degrees of phase mixing or SS crystallization between THF and water samples. Shown in Figure 6A is a graph of the composite modulus as a function of volume fraction for all samples studied. PEG-SWCNT composites showed the largest increase in Young’s modulus of all of the composites. For PEG-SWCNT composites, by 12.4 vol %, the Young’s modulus reached its maximum value of 286 MPa, an increase of ×23 relative to the polymer (TPU/H2O, Y ) 12 MPa). The maximum modulus observed for the PABS-SWCNT composites was 230 MPa at Vf ) 23%. For composites filled with ODA-SWCNTs, the Young’s moduli were lower, reaching 75 MPa for the 12 vol % composite which compares with 6.3 MPa for the polymer (TPU/THF). The higher levels of reinforcement observed for the water-based composites are consistent with the reinforcement of the soft segments by the nanotubes. However the reinforcement of the (already stiff) hard segments by the ODA-SWCNTs would be expected to result in lower reinforcement values as is observed. The reinforcement values measured here are significantly higher than those shown in other similar work on functionalized CNTs-PU. Buffa et al. fabricated functionalized-SWCNTpolyurethane composites. In their work they achieved an increase of ∼3 times in the Young’s modulus at 20 wt % loading.49 In

Figure 6. (A) Young’s modulus and (B) ultimate tensile strength of PU-functionalized nanotube composites as a function of nanotube content.

another attempt, Gopal et al.30 demonstrated an increase in Young’s modulus of 7.4 times at this same CNT loading. They used functionalized MWCNTs in their polyurethane work, but the ductility was greatly reduced by a factor of ∼7. In a recent study (nanoclay-TPU composites), Liff et al. achieved high reinforcement. They observed a 23 fold increase in Young’s modulus from 38 to 856 MPa at 20 wt % in their TPU-nanoclay work by restricting the nanoclay to the HS. This work suggests that, by restricting nanofiller to the HS, significant reinforcement can be achieved without sacrificing ductility at high filler loading.32 We note that another contributing factor in the achievement of good reinforcement in this work can be the enhancement in the crystallinity,7,50 as revealed by DSC and discussed earlier. In addition to enhancing the modulus, it would be useful to increase the stress in the plateau region. For the purpose of discussion we define the “plateau stress” as the stress at 50% strain for all samples. The plateau stress is shown in Figure 7A for all samples. The plateau stress increases only weakly for the ODA-SWCNT samples, from 1.5 MPa for the TPU/THF sample to 4 MPa for the 12 vol % sample. The increase is much more significant for the water-based composites. For both PEGSWCNT and PABS-SWCNT containing composites, the plateau stress increases in a very similar fashion from 2 MPa for TPU/ H2O to 7.6 and 6.7 MPa respectively for the 6 vol % sample. We note that the evidence previously presented suggests that the ODA-SWCNTs selectively interact with the HS while the PEG and PABS-SWCNTs selectively interact with the SS. That the ODA-SWCNTs do not significantly influence the plateau stress confirm that the plateau region is associated with stretching of the SS and is not significantly influenced by the mechanical properties of the HS. We expect that the interaction of the water-soluble nanotubes with the SS tends to hinder the mobility of these segments making it more difficult for them to

Reinforcement of Thermoplastic Polyurethane

J. Phys. Chem. C, Vol. 114, No. 26, 2010 11405

Figure 8. Slopes of high strain region as a function of nanotube volume fraction for TPU-functionalized nanotube composites. This data shows the effect of nanotube type on strain hardening.

Figure 7. Impact of addition of functionalized nanotubes on stress in the plateau region. (A) Plateau stress (stress at 50% strain) as a function of nanotube volume fraction. (B-D) Stress at 100%, 200%, and 300% strain, respectively, as a function of nanotube volume fraction.

stretch and align. This results in an increase in the plateau stress. These results are echoed in the data for the stress at 100% (Figure 7B), 200% (Figure 7C) and 300% strains (Figure 7D), which show substantially the same trends. Generally speaking, in TPU stress-strain curves, the steepest region, which follows the plateau region, is usually associated with strain induced crystallization (SIC) or strain hardening (SH).32,39 Strain induced alignment promotes the formation of crystallites which reinforce the material, hence increasing the slope, dσ/dε. This can be influenced by the presence of additives. For example, higher slopes due to the incorporation of silicate particles into TPU-silicate nanocomposites can be seen in the work of Finnigan et al.51 We can monitor this process by measuring dσ/dε at a fixed strain. This data is shown in Figure 8 for a stain of 450%. No real change in dσ/dε at ε ) 450% is observed for the ODA-SWCNTs filled samples. In contrast, the water-based composites display a significant increase in dσ/dε at ε ) 450% and hence SIC. The highest values of dσ/dε at ε ) 450% were observed for PABS-SWCNTs composites with Vf in the range 0.01-0.04. This data is consistent with our understanding that the ODA-SWCNTs tend to interact with the HS while the PEG and PABS-SWCNTs interact with the SS if we assume that strain hardening is associated with the soft segments only. The final portion of a stress-strain curve is defined by material failure. One way to quantify this is through the stress at break or the ultimate tensile strength (UTS). For all samples, the UTS is plotted as a function of volume fraction in Figure 7B. For the ODA-SWCNT composites, significant increases in the UTS from the TPU/THF value of 26 MPa are observed. The maximum values observed were >60 MPa for volume fractions around 10-3-10-2. As the volume fraction is increased above 10-2, the UTS falls off slightly to 35 MPa for the Vf )

0.12 sample. For the PEG-SWCNTs and PABS-SWCNTs samples, the strength increase is less impressive; from 41 MPa for TPU/H2O to 55-60 MPa at volume fractions around 10-2 for both water-based composites. Interestingly, the UTS collapses more completely for the water-based composites compared to the ODA-SWCNT filler materials, falling to ∼10 MPa for Vf ∼ 0.25. The dramatic enhancement of the UTS for the ODA-SWCNT samples can be attributed to their interaction with the HS. We propose that by selectively reinforcing the HS, they prevent their breakup under strain,39 and ultimately increase the UTS. In addition, the UTS data is consistent with the idea that up to Vf ∼ 10-2 all of the THF soluble CNTs remain embedded in the HS, but above Vf∼10-2 CNTs begin to influence the SS, resulting in a drop in UTS at higher volume fractions. In contrast, by interacting with the SS, the water-soluble nanotubes impart a much lower increase in UTS. In addition, their presence in the SS caused a larger decrease in UTS at high volume fractions compared to the ODA-SWCNTs. We can also examine the effect of the nanotubes on the strain at failure (Figure 9). The polymer only samples exhibit strain at break values of ∼500% and ∼650% for the TPU/THF and TPU/H2O samples, respectively, values typical for a segmented PU.28,32,39,52 For low volume fractions, ODA-SWCNT composites show significantly increased strain at break values, reaching 750% for the Vf ∼10-3 sample. At higher volume fractions, εB falls off, saturating at ∼600-650% for volume fractions of ∼0.1 This is surprising as εb usually decreases sharply at high CNT filler loadings.28,30 It has been shown that the strain at break in segmented TPU remains high if the nano filler remains in or interacts only with the HS.32 This agrees well with the previously presented evidence for the restriction of the ODA-SWCNTs to the HS. It may also be that these CNTs facilitate the mobility of the SS and hence can actually increase ductility. In water-based composites, no real increase in ductility is observed for low volume fractions. However for volume fractions above 0.01, a sharp reduction in strain at break was recorded, with εB collapsing to ∼10% for the 0.25 volume fraction sample. Large strains before breaking in polyurethane are due to SS extension. Therefore, any additive which binds to the SS not only reinforces the sample by restricting the SS mobility but has a negative impact on the strain at break value. Thus a sharp decline in εb is associated with interaction of watersoluble CNT with the SS.

11406

J. Phys. Chem. C, Vol. 114, No. 26, 2010

Figure 9. (A) Strain at break and (B) toughness of PU-functionalized nanotube composites as a function of nanotube content.

An important parameter for elastomers is the toughness. This is the energy required to break the samples, as calculated from the area under the stress strain curve. The stiffness, plateau strength, plateau region slope, and the stiffness of the strain induced crystallization region can all influence the toughness within a particular TPU composites system. However, the main contributors are the strain at break and UTS values. As the strain at break and UTS values are in the order, ODA-SWCNT > PEGSWCNT/PABS-SWCNT composites, so the toughness will follow a similar trend. At low volume fractions, the ODA-SWCNT composites exhibited significantly higher toughness compared to the polymer only samples. For these composites the toughness initially increased sharply by a factor of ×4.2, from 36 (TPU/ THF) to 151 MJ/m3 (0.1 vol %). The increase in the ODASWCNT samples’ toughness saturated around 80 MJ/m3 for volume fractions around 0.1. However, in the water-based composites at low volume fractions, a relatively lower increase in toughness was observed. Compared to a TPU/H2O value of T ) 85 MJ/m3, the toughness increased to maxima of 113 and 117 MJ/m3 for PEG and PABS-SWCNTs, respectively. However, the toughness sharply decreased at high volume fractions, in both cases falling below 50 MJ/m3 for volume fractions above 0.1. These trends are virtually identical to those observed for the strain at break and are due to the same factors. We note that the composite prepared in this work demonstrated mechanical properties, i.e., Young’s modulus, plateau stress, and UTS, in the same range as reported values for TPU with hard segment content of 47-65 wt %.31,34 Such materials display the highest moduli and plateau stress reported for TPUs. However, the elevated HS content typically results in a significant reduction in ductility. For example a TPU with HS content of 47% has been reported to display Y ) 200 MPa and plateau stress of σε)50% ) 16 MPa. However this was achieved

Khan et al. by reducing the ductility to 200%.31 Some of the samples prepared in this work display similar moduli and plateau stress. For example the 7 vol % PEG-SWCNT sample has Y ) 180 MPa and σε)50% ) 8 MPa but retains ductility of 450%. Interaction of Functionalized Nanotubes with TPU Segments. We have made a case for a specific interaction between ODA-SWCNTs and the TPU HS and between PEG-SWCNTs and PABS-SWCNTs and the TPU SS. It is important to discuss the nature of this interaction. This discussion is limited by our lack of knowledge of the TPU structure; the supplier declined to provide information on the exact HS and SS structure. However we can consider the type of segments typically used in TPU synthesis. The soft segments generally consist of polyols such as polyethers and polyesters which can be hydrophilic, and so it is not surprising that they may be compatible with water-soluble nanotubes. The HS generally consist of groups such as MBI-BD, TDI-BD, HDI-BD, HMDI-BD, and IPDIBD (where MBI ) methylene bisphenyl isocyanate, TDI ) touluene diisocynate, HDI ) hexamethylene diisocyanate, HMDI ) methylenebis (cyclohexyl isocyanate), IPDI ) isophorone diisocynate, and BD ) butanediol). In the case of nanotubes, it has been shown that the interaction between functionalized nanotubes and solvents is dominated by the interaction between functional group and solvent.18 Thus, we can rate the compatibility of the HS with the ODA-SWCNTs by looking at the compatibility between these groups and ODA. Chemical compatibility can be quantified via the Hildebrand solubility parameter, δT. Mixing of different molecules at the nanoscale is favored if both have similar δT. The solubility parameter of ODA is 19.5 MPa1/2.18 We have used HSPiP software (www.hansen-solubility.com) to calculate the solubility parameters for the typical HS groups listed above using the Stefanis-Panayiotou algorithm. The resulting values are given in Table 1. In all cases, the calculated solubility parameters of typical HS segments are higher than that of ODA. The HS group with the closest δT is IPDI-BD (22 ( 1 MPa1/2) which is 2.5 ( 1 MPa1/2 higher than that of ODA. We note that ODA-SWCNTs are soluble in solvents with δT values which differ by up to 2 MPa1/2 from ODA. This means that IPDI-BD should be just compatible with ODA-SWCNTs although the other HS groups are likely to be less compatible. This suggests that the HS in the TPU studied here is in fact a low δT group. In future work, the best approach would be to design the function group on the nanotube to be compatible with the segment of interest (HS or SS) in the TPU of choice, depending on the application. Such chemistry should be relatively straightforward. We expect that this will open the way to tailored reinforcement of TPU with functionalized nanotubes. Conclusion Generally, the introduction of large amounts of CNTs to TPU is accompanied by a loss in mechanical properties due to the CNTs interactions with SS.28,30 In this work we were able to prepare very high volume fraction composites with tunable mechanical properties by targeted reinforcement of either SS or HS of TPU. The targeted SS and HS interaction of water dispersed CNTs and ODA-SWCNTs respectively are evident from the DSC and DMTA data. The increase in crystallinity of the SS on the introduction of water dispersed CNTs to TPU implies their affinity toward the SS. However, for ODA-SWCNT samples, no change in SS crystallinity is observed, while the nucleation of the HS crystallites in the composites indicates a specific interaction between ODA-SWCNT and the HS. The higher dynamic Tg of the composites compared to the polymer,

Reinforcement of Thermoplastic Polyurethane

J. Phys. Chem. C, Vol. 114, No. 26, 2010 11407

TABLE 1: Hildebrand Solubility Parameters of Possible Hard Segments Compared to ODA HS group δ

T

1/2

(MPa )

MBI-BD

TDI-BD

HDI-BD

HMDI-BD

IPDI-BD

ODA

26 ( 1

25.5 ( 1

24.5 ( 1

23.7 ( 1

22 ( 1

19.5

in the water-based samples, suggests the restriction of molecular mobility due to the CNTs bonding to the SS. While no change in Tg of the ODA-SWCNT composites, underlines that there is no interaction between the SS and the ODA-SWCNTs. In addition, the large difference in scale between nanotube length and the dimensions associated with different polymer phases remains problematic. As the soft segments form a continuous phase, it is straightforward to consider a nanotube to be completely embedded in the SS phase. However the length scale associated with hard segment phase is much smaller, typically 10s of nm.53,54 Given that the nanotubes are several hundred nm long, it is hard to see how the nanotubes can be solely contained in the HS phase. One possibility is that average HS size increases dramatically on the introduction of nanotube. However, further work is required to understand this point. The concept of targeted insertion of nanotubes into SS or HS has implications for the composite mechanical properties. The increases in stiffness and stress in all three part of stress-strain curve, coupled with the sharp decrease in ductility with increasing SWCNT content in water-based composites can be attributed to the CNTs interactions with SS. However ODASWCNT composites have slightly lower stiffness and saturation rather than sharp decline of other mechanical properties, precisely because these ODA-SWCNTs do not affect the SS. To conclude, when nanotubes are mixed with a thermoplastic elastomer such as TPU, the tubes have only three choices: (A) to disperse in the SS, (B) to disperse in the HS, or (C) to aggregate. Functionalization of the nanotubes can prevent (C) and depending on the chemical structure of the functionality can drive the system to either (A) or (B). We suggest that such treatment will facilitate the development of composites with tailored mechanical properties. Experimental Section Thermoplastic polyurethane/H2O (product code U2-01) emulsion was provided by Hydrosize55 with an average particle size ∼3 µm. The dispersion was diluted to 50 mg/mL with deionized water. Single walled nanotubes functionalized with polyethylenegycol (PEG-SWCNTs), poly(amino benzene sulfonic acid) (PABS-SWCNTs), and octadecylamine (ODASWCNTs) were purchased from Carbon Solutions, Inc.56 According to the suppliers, the average molecular weights of the functional groups were ∼600 (PEG), 466 (PABS), and 268 g/mol (ODA). For both PEG-SWCNTs and PABS-SWCNTs, 66 mg were dispersed in 10 mL of deionized water by energetic agitation of CNTs/H2O using an ultrasonic tip (GEX600, 700W, 20%, 120 kHz) for 10 min. This was followed by 4 h of sonication in an ultrasonic bath (Branson 1510 MT). Each of the CNT dispersions was then mixed with a prepared TPU/ H2O emulsion using the same ultrasonic tip (5 min) to give composite dispersions with final concentrations of CNTs of 5.2 mg/mL and of TPU of 10.6 mg/mL. These stock solutions were used for all further work. Each of the stock composite solutions was diluted with TPU/H2O to give a series of solutions of varying nanotube mass fraction. Each of these dilutions was sonicated for 5 min under the ultrasonic tip, followed by further sonication for 4 h in the ultrasonic bath. To prepare ODA-SWCNT composites, the supplied TPU was dried at 60 °C for 3 days in Petri dishes. The dried TPU was

cut into small pieces. A 50 mg/mL solution of this dried TPU was prepared in tetrahydrafuran (THF) at room temperature by magnetic stirring. The THF was purchased from Sigma Aldrich. The rest of the ODA-SWCNT composite solutions preparation procedure was the same as that which was carried out for waterbased systems. The only difference was that ultrasonication was undertaken in an ice beaker to avoid evaporation of the highly volatile THF. Films were prepared from both composite and polymer-only solutions by dropcasting into 40 × 40 mm Teflon trays. The water-based solutions were dried at 60 °C in an oven for 12 h, while the THF-based solutions were dried at room temperature for 12 h followed by further heating at 60 °C for 2 h. Thermogravimetry showed virtually no solvent remained after drying. Free standing films were obtained. These films were cooled in ambient condition to stabilize them. Films were cut into strips using a die cutter with a fixed spacing of 2.25 mm. The film thicknesses were measured using a micrometer screw and were typically ∼0.08 mm. Thermogravimetric analysis (TGA) was carried out to assess the solvent entrapment in the sample and also to determine the onset of polymer degradations using a Perkin-Elmer Pyris 1 TGA. The temperature range was from 25 - 900 °C with scan of 10 °C/minute. TGA analysis of the TPU revealed the presence of trapped solvent and polymer degradation above 230 °C (Supporting Info). For this reason the maximum temperature was kept below 230 °C for all of the other thermal studies (DSC and DMTA). Differential scanning calorimetry (DSC) was carried out with a Perkin-Elmer Diamond DSC with a heating scan rate of 20 °C/minute. Dynamic thermal mechanical tests were carried out using a Perkin-Elmer Diamond DMA in the temperature range from -60 to +170 °C with a frequency of 1 Hz. Tensile tests were carried out using a Zwick Roell with 100 N load cell at a strain rate of 100 mm/min. Acknowledgment. The authors acknowledge Science Foundation Ireland through the principle investigators scheme for supporting this work. References and Notes (1) Che, J. W.; Cagin, T.; Goddard, W. A. Nanotechnology 2000, 11, 65. (2) Berber, S.; Kwon, Y.-K.; Toma´nek, D. Phys. ReV. Lett. 2000, 84, 4613. (3) Wei, B. Q.; Vajtai, R.; Ajayan, P. M. Appl. Phys. Lett. 2001, 79, 1172. (4) Kim, B. M.; Fuhrer, A. M. S. J. Phys.: Condensed Matter 2004, 16, R553. (5) Biercuk, M. J.; Llaguno, M. C.; Radosavljevic, M.; Hyun, J. K.; Johnson, A. T. Appl. Phys. Lett. 2002, 80, 2767. (6) Xie, S. S.; Li, W. Z.; Pan, Z. W.; Chang, B. H.; Sun, L. F. J. Phys. Chem. Solids 2000, 61, 1153. (7) Coleman, J. N.; Khan, U.; Blau, W. J.; Gun’ko, Y. K. Carbon 2006, 44, 1624. (8) Iijima, S. Nature 1991, 354, 56. (9) Coleman, J. N.; Cadek, M.; Blake, R.; Nicolosi, V.; Ryan, K. P.; Belton, C.; Fonseca, A.; Nagy, J. B.; Gun’ko, Y. K.; Blau, W. J. AdV. Funct. Mater. 2004, 14, 791. (10) Velasco-Santos, C.; Martinez-Hernandez, A. L.; Fisher, F.; Ruoff, R.; Castano, V. M. J. Phys. D: Appl. Phys. 2003, 36, 1423. (11) Wagner, H. D.; Lourie, O.; Feldman, Y.; Tenne, R. Appl. Phys. Lett. 1998, 72, 188.

11408

J. Phys. Chem. C, Vol. 114, No. 26, 2010

(12) Kilbride, B. E.; Coleman, J. N.; Fraysse, J.; Fournet, P.; Cadek, M.; Drury, A.; Hutzler, S.; Roth, S.; Blau, W. J. Journal Of Applied Physics 2002, 92, 4024. (13) Giordani, S.; Bergin, S. D.; Nicolosi, V.; Lebedkin, S.; Kappes, M. M.; Blau, W. J.; Coleman, J. N. J. Phys. Chem. B 2006, 110, 15708. (14) Giordani, S.; Bergin, S.; Nicolosi, V.; Blau, W. J.; Coleman, J. N. Phys. Status Solidi B 2006, 243, 3058. (15) Zhao, B.; Hu, H.; Yu, A.; Perea, D.; Haddon, R. C. J. Am. Chem. Soc. 2005, 127, 8197. (16) Zhao, B.; Hu, H.; Haddon, R. C. AdV. Funct. Mater. 2004, 14, 71. (17) Niyogi, S.; Hamon, M. A.; Hu, H.; Zhao, B.; Bhowmik, P.; Sen, R.; Itkis, M. E.; Haddon, R. C. Acc. Chem. Res. 2002, 35, 1105. (18) Amiran, J.; Nicolosi, V.; Bergin, S. D.; Khan, U.; Lyons, P. E.; Coleman, J. N. J. Phys. Chem. C 2008, 112, 3519. (19) Coleman, J. N.; Cadek, M.; Blake, R.; Nicolosi, V.; Ryan, K. P.; Belton, C.; Fonseca, A.; Nagy, J. B.; Gun’ko, Y. K.; Blau, W. J. AdV. Funct. Mater. 2004, 14, 791. (20) Dalton, A. B.; Collins, S.; Munoz, E.; Razal, J. M.; Ebron, V. H.; Ferraris, J. P.; Coleman, J. N.; Kim, B. G.; Baughman, R. H. Nature 2003, 423, 703. (21) Vigolo, B.; Penicaud, A.; Coulon, C.; Sauder, C.; Pailler, R.; Journet, C.; Bernier, P.; Poulin, P. Science 2000, 290, 1331. (22) Qian, D.; Dickey, E. C.; Andrews, R.; Rantell, T. Appl. Phys. Lett. 2000, 76, 2868. (23) Blighe, F. M.; Werner, B. J.; Coleman, J. N. Nanotechnology 2008, 19, 415709. (24) Koerner, H.; Price, G.; Pearce, N. A.; Alexander, M.; Vaia, R. A. Nat. Mater. 2004, 3, 115. (25) Meng, Q.; Hu, J.; Yeung, L. Smart Mater. Struct. 2007, 830. (26) Kwon, J.; Kim, H. J. Polym. Sci. Part A: Polym. Chem. 2005, 43, 3973. (27) Xia, H.; Song, M. Soft Matter 2005, 1, 386. (28) Chen, W.; Tao, X.; Liu, Y. Composite Sci. Technol. 2006, 66, 3029. (29) Xia, H.; Song, M. J. Mater. Chem. 2006, 16, 1843. (30) Yoo, H. J.; Jung, Y. C.; Sahoo, N. G.; Cho, J. W. J. Macromol. Sci., Part B 2006, 45, 441. (31) Korley, L. T. J.; Pate, B. D.; Thomas, E. L.; Hammond, P. T. Polymer 2006, 47, 3073. (32) Liff, S. M.; Kumar, N.; McKinley, G. H. Nat. Mater. 2007, 6, 76. (33) Crawford, D. M.; Bass, R. G.; Haas, T. W. Thermochim. Acta 1998, 323, 53. (34) Kim, H.-D.; Lee, T.-J.; Huh, J.-H.; Lee, D.-J. J. Appl. Polym. Sci. 1999, 73, 345.

Khan et al. (35) Paik, I. H.; Goo, N. S.; Jung, Y. C.; Cho, J. W. Smart Mater. Struct. 2006, 15, 1476. (36) Bliznyuk, V.; Singamaneni, S.; Kattumenu, R.; Atashbar, M. Appl. Phys. Lett. 2006, 88, 164101. (37) Luo, Y.; Wang, C.; Li, Z. Synth. Met. 2007, 157, 390. (38) Jeong, H. M.; Lee, J. B.; Lee, S. Y.; Kim, B. K. J. Mater. Sci. 2000, 35, 279. (39) Koerner, H.; Liu, W.; Alexander, M.; Mirau, P.; Dowty, H.; Vaia, R. A. Polymer 2005, 46, 4405. (40) Koerner, H.; Kelley, J. J.; Vaia, R. A. Macromolecules 2008, 41, 4709. (41) Cadek, M.; Coleman, J. N.; Ryan, K. P.; Nicolosi, V.; Bister, G.; Fonseca, A.; Nagy, J. B.; Szostak, K.; Beguin, F.; Blau, W. J. Nano Lett. 2004, 4, 353. (42) Bhattacharyya, A. R.; Sreekumar, T. V.; Liu, T.; Kumar, S.; Ericson, L. M.; Hauge, R. H.; Smalley, R. E. Polymer 2003, 44, 2373. (43) Sandler, J.; Broza, G.; Nolte, M.; Schulte, K.; Lam, Y. M.; Shaffer, M. S. P. J. Macromol. Sci.-Phys. 2003, B42, 479. (44) Xie, X. L.; Aloys, K.; Zhou, X. P.; Zeng, F. D. J. Thermochem. Anal. Calorim. 2003, 74, 317. (45) Jeong, H. M.; Song, J. H.; Lee, S. Y.; Kim, B. K. J. Mater. Sci. 2001, 36, 5457. (46) Hatakeyama, T.; Zhenhai, L. Handbook of thermal analysis; John Wiley & Sons Ltd.: New York, 1998. (47) Sperling, L. H. Introduction to Physical Polymer Science, 3rd ed.; Wiley: New York, 2001. (48) Khan, U.; Ryan, K. P.; Blau, W. J.; Coleman, J. N. Composite Sci. Technol. 2007, 67, 3158. (49) Buffa, F.; Abraham, G. A.; Grady, B. P.; Resasco, D. J. Polym. Sci. Part B: Polym. Phys. 2007, 45, 490. (50) Ward, I. M. Mechanical properties of solid polymers; Wiley: New York, 1971. (51) Finnigan, B.; Jack, K.; Campbell, K.; Halley, P.; Truss, R.; Casey, P.; Cookson, D.; King, S.; Martin, D. Macromolecules 2005, 38, 7386. (52) Yeh, F.; Hsiao, B. S.; Sauer, B. B.; Michel, S.; Siesler, H. W. Macromolecules 2003, 36, 1940. (53) Martin, D. J.; Meijs, G. F.; Gunatillake, P. A.; Yozghatlian, S. P.; Renwick, G. M. J. Appl. Polym. Sci. 1999, 71, 937. (54) Taylor, F. E.; Laity, P. R.; Wong, S. S.; Norris, K.; Khunkamchoo, P.; Cable, M.; Andrews, G.; Johnson, A. F.; Cameron, R. E. Microsc. Microanal. 2006, 12, 151. (55) http://www.hydrosize.com/. (56) http://www.carbonsolution.com.

JP102938Q