Article pubs.acs.org/Macromolecules
Influence of Water on the Structure and Properties of PDMSContaining Multiblock Polyurethanes Kimberly A. Chaffin,*,† Adam J. Buckalew,† James L. Schley,† Xiangji Chen,† Matthew Jolly,† Julie A. Alkatout,† Jennifer P. Miller,† Darrel F. Untereker,† Marc A. Hillmyer,*,‡ and Frank S. Bates*,§ †
Science and Technology, Medtronic Incorporated, 710 Medtronic Parkway, Minneapolis, Minnesota 55432, United States Department of Chemistry and §Department of Chemical Engineering and Materials Science, University of Minnesota, Minneapolis, Minnesota 55455, United States
‡
S Supporting Information *
ABSTRACT: Segmented polyurethane multiblock polymers containing polydimethylsiloxane and polyether soft segments form tough and easily processed thermoplastic elastomers. Two commercially available examples, Elast-Eon E2A (denoted as E2A) and PurSil 35 (denoted as P35), were evaluated for molecular and mechanical stability after immersion in buffered water for up to 52 weeks at temperatures ranging from 37 to 85 °C. Dynamic mechanical spectroscopy experiments, performed in tension and shear, were used to characterize the linear viscoelastic properties of compression-molded (dry) specimens. Smallangle X-ray scattering measurements indicated a disorganized microphase-separated morphology for all test conditions. Upon aging in phosphate buffered saline, samples of E2A and P35 were analyzed by size exclusion chromatography (SEC) and tensile testing as a function of time and temperature. The absolute molar mass of each material prior to aging in water was determined by SEC using a multiangle light scattering detector. Aging at 85 °C and 52 weeks lead to a 67% and 50% reduction in molar mass from the original values for E2A and P35, respectively. We attribute the reduction in molar mass to hydrolysis of the polymer backbone and have evaluated the data using a pseudo-zero-order kinetics analysis. The temperature dependence of the extracted rate data is consistent with an activated (i.e., Arrhenius) process, and thus all the molar mass reduction data can be reduced to a single master curve. Concomitant with the reduction in molar mass, E2A and P35 transformed with aging from strain-hardening to strain-softening materials, characterized by substantially reduced tensile strength (stress at failure) and ultimate elongation (strain at failure) relative to the original properties.
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INTRODUCTION
responsible for producing mechanical integrity in these hybrid materials.10−12 Although various diisocyantes are available for polyurethane syntheses, the two classic examples of toluene diisocyante and methylene diphenyl diisocyanate dominate the landscape.10 Thus, the characteristics of practical segmented polyurethanes are primarily controlled by the choice of the dihydroxy components. In particular, monomeric diols that participate in so-called hard segment formation and oligomeric or polymeric diols that make up the soft segments largely determine ultimate properties. Dihydroxy-terminated polydimethylsiloxane (PDMS diol) is a particularly intriguing soft segment component as it offers an unusual degree of control over morphology and viscoelastic behavior due to its extreme nonpolar character and ultralow glass transition temperature.13,14 Furthermore, polyurethanes that contain both PDMS segments and various polyether soft segments, like the materials we have studied here, have been reported to exhibit enhanced biostability relative to conventional thermoplastic polyether urethane (PEU) elastomers composed of pure polyether soft segments such as the well-known Elasthane
Advances in synthetic polymer chemistry coupled with improved structural characterization techniques and predictive theory have transformed the design of synthetic polymers for commodity and value-added commercial applications. For example, the organization of discrete sequences of chemically distinct repeat units into multiblock architectures provides scientists and engineers with nearly unlimited opportunities for creating products that simultaneously exhibit enhanced end-use properties and facile processing characteristics.1 Segmented polyurethanes, prepared by condensation polymerization of diisocyanate and dihydroxy compounds (monomers and oligomers or polymers), represent an important class of such multiblock polymers.2 These materials find numerous uses as flexible foams for insulation,3 bedding and automotive purposes, protective coatings, as adhesives,4,5 and in various biomedical devices.5,6 The ultimate properties of segmented polyurethanes, including mechanical response, use temperatures, and environmental stability are controlled by the choice of ingredients, drawn from an ever-expanding pallet of commercially available starting materials.7,8 Thermodynamic incompatibility of chemically different blocks generally leads to nanoscale segregation. Often referred to as microphase separation,9 this organization at the molecular scale is primarily © 2012 American Chemical Society
Received: September 19, 2012 Revised: November 5, 2012 Published: November 13, 2012 9110
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80A (produced by DSM Biomedical).15,16 Biostability of these new materials is usually associated with the well-established oxidative stability of the PDMS backbone.17 In this work, we examine the chemical and mechanical stability of two commercially available PDMS-based PEU compoundsElast-Eon E2A (denoted as E2A) (manufactured by AorTech, International) and PurSil 35 (denoted as P35) (manufactured by DSM Biomedical)following exposure to deoxygenated water for up to one year at various temperatures. These thermoplastic elastomers are produced by combining four principal ingredients: 4,4′-methylenediphenyl diisocyanate (MDI), 1,4-butanediol (BDO), α,ω-bis(6hydroxyethoxypropyl)polydimethylsiloxane (PDMS-diol), and either poly(hexamethylene oxide) (PHMO) in E2A18,19 or poly(tetramethylene oxide) (PTMO) in P35.15 The structures of these building blocks are illustrated in Figure 1. (The P35
irregularly within individual chains. These constraints, combined with relatively large dispersities in chain length resulting from the statistical nature of the condensation polymerization, produce a state of molecular packing frustration that results in a disorganized morphology. This scenario, microphase separation without long-range order, is in contrast with the more familiar morphologies exhibiting particular ordered state symmetries (e.g., lamellae, gyroid, cylinders, and spheres) that characterize more well-defined block polymers (e.g., discrete AB diblock copolymers and ABC triblock terpolymers) produced by controlled polymerization techniques.23 The absorption of water by polyurethanes in general can influence their properties and stability.24 For example, Elasthane 80A, a PTMO-based PEU elastomer, absorbs about 2 wt % water under physiological conditions (see below). As we demonstrate, the addition of hydrophobic PDMS to a PEU formulation reduces the overall water absorption as expected. However, the associated nanoscale morphology described above forces intimate contact between water and all the constituent blocks, including the hydrophobic PDMS. This situation is distinct from that of pure PDMS-based elastomers. To probe the influence of this nanoscale morphology on the long-term stability of PDMS containing PEU elastomers, we have investigated the molar mass and tensile properties of E2A and P35 with equilibrium water contents following exposure to deoxygenated buffered saline solutions at temperatures between 37 and 85 °C for up to one year. Through these studies, we determined that the observed reduction in the molar mass of both E2A and P35 exhibits an Arrhenius rate dependence and is accompanied by significant reduction in the tensile strength and extension at failure (both measured at 37 °C).
Figure 1. Structures of the components used to synthesize E2A and P35. The polyether in E2A is PHMO, and the polyether in P35 is PTMO. Also depicted is a schematic of one possible sequence of blocks in these multiblock polyurethanes.
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recipe reportedly includes a monohydroxy PDMS oligomer that is known as a surface modifying end group (SME), which terminates a small percentage of the polyurethane chain ends.20) Both PDMS-PEU materials are produced commercially in a bulk polymerization, leading to multiblock polymers that contain soft (PDMS and PHMO or PTMO) and hard (MDI-BDO) segments coupled together by carbamate (urethane) linkages. Because the condensation polymerization process involves three different diols (ignoring the monofunctional SME in the case of P35), the product is a complex mixture of macromolecules, each chain containing a statistical array of coupled segments. For example, E2A has six possible MDI-linked triads (PDMS-MDI-PDMS, PDMS-MDI-PHMO, PDMS-MDI-BDO, PHMO-MDI-PHMO, PHMO-MDI-BDO, and BDO-MDIBDO) that in turn can combine to form numerous higher order block sequences depending on the reaction stoichiometry, sequence of monomer addition, and relative reactivities. Here we define soft segments as either individual PDMS or PH(T)MO chains (or MDI coupled PDMS and/or PH(T)MO chains) flanked by MDI-BDO units. The hard segments in these polyurethanes consist of -(MDI-BDO)m- sequences of various lengths. In addition to the classic hard segment/soft segment self-organization, the incompatibility between PDMS and PH(T)MO results in microphase separation of these components as well.21,22 The morphology and physical properties of the resulting elastomers are directly influenced by the ensemble of chain architectures produced.18 Most notably, microphase separation must accommodate distributions of soft and hard segment lengths strung together
EXPERIMENTAL SECTION
Materials. Elast-Eon E2A (also referred to as Elast-Eon 2A and denoted here as E2A) (Aortech, International) and Elasthane 80A and PurSil 35 (denoted P35) (DSM Biomedical) were received from the manufacturers. E2A and P35 are both soft durometer materials with Shore hardness values of 90A and 80A, respectively. Elasthane 80A (Shore 80A hardness), an analogue of P35 absent the PDMS blocks, was used as a reference material. The chemical composition of the polyurethanes was determined by nuclear magnetic resonance (NMR) spectroscopy. Polymers were dissolved in THF-d8 at a concentration of ∼20 mg/mL, and spectra were acquired using a JEOL Eclipse 400 MHz NMR spectrometer. Polyether and PDMS block lengths (i = 6.7 for PHMO, i = 13.6 for PTMO, and j = 9.2, Figure 1) and the composition of the polymers were estimated based on integration values in the corresponding 1H NMR spectra. The results, listed in Table 1 in the form of polymerization feedstock ingredients, are consistent with published values for these materials.15,18,21 As-received polymer pellets were dried and compression-molded into 0.7 mm thick sheets at 220 °C. E2A formed transparent films while the P35 sheets displayed an inhomogeneous cloudy appearance (see Supporting Information). ASTM D1708 microtensile bars and circular disks were stamped from the pressed films and used for
Table 1. Weight Percentages of the Components in E2A and P35
a
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PHMO
E2Aa P35b
12
PTMO
PDMS-diol
MDI
BDO
24
49 36
33 34
6 6
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Table 2. Molar Mass Characterization of E2A and P35 SEC-MALLSc sample a
E2A P35b a
−1
SEC (PS standards)c −1
−1
−1
Mn (kg mol )
Mw (kg mol )
Đ
dn/dc (mL g )
Mn (kg mol )
Mw (kg mol−1)
Đ
48 45
101 88
2.2 2.0
0.1020 0.1075
103 119
200 214
1.9 1.9
Elast-Eon E2A. bPurSil35. cUncertainty in Mw and Mn are estimated to be 5−10%
mechanical and rheological testing, respectively. Molding at 220 °C was shown to eliminate allophanates (bonding of terminal isocyanate and carbamate groups) in the as received materials as discussed in the Supporting Information. Polymer Aging in PBS. Tensile bars of E2A and P35 were immersed in phosphate buffered saline (PBS) solution obtained from Sigma-Aldrich (P5368). PBS is a salt-based water solution with osmolarity and ion concentrations that match those found in the human body (isotonic). House nitrogen was continuously bubbled through the solutions to minimize the presence of oxygen; periodic monitoring with an electronic probe (VWR sympHony SP90M5 handheld meter with dissolved oxygen probe) showed less than 4% of the saturation concentration of oxygen in PBS at ambient conditions (i.e., air at room temperature and pressure). The pH was monitored during the course of the experiments with a pH meter and remained relatively constant at pH = 7.4 ± 0.5 over the course of the aging experiments. Samples were aged in PBS solution at 37, 55, 70, and 85 °C (±2 °C) for up to one year. Weight gain was measured gravimetrically over several days following immersion of E2A, P35, and Elasthane 80A in deionized water at 37 °C and in PBS at 37, 70, and 85 °C. Samples were removed from solution periodically, blotted dry, weighed, then dried under reduced pressure at ambient temperature for three days, and weighed again. In all instances the percent water adsorbed reached a constant value within 4 h of immersion, and the results for deionized water and PBS were indistinguishable. Longer-term (up to 52 weeks) weight gain experiments were performed in PBS at 37 °C. Molar Mass Determination. The absolute weight-average molar mass Mw and dispersity Đ = Mw/Mn of compression-molded E2A and P35 were determined by multiangle laser light scattering−size exclusion chromatography (MALLS-SEC) operated with THF as the mobile phase at 1 mL min−1 and 25 °C. The instrument was outfitted with an Agilent 1260 Infinity series injection system, three Waters Styragel columns (HR6, HR4, and HR1), a Wyatt DAWN Heleos II multiangle laser light scattering, MALLS, detector with 18 measurable angles from 10° to 160°, and a laser wavelength of 663.3 nm. Eluted samples were analyzed by a Wyatt OPTILAB T-rEX refractive index (RI) detector employing a laser wavelength of 658 nm. Data from the MALLS and RI detectors were used in combination with the Wyatt Astra 6 software to determine the refractive index increment dn/dc and absolute values of Mw and Mn. (Here we note that compositional heterogeneity between different chains (i.e., variations in dn/dc) produces additional light scattering intensity, resulting in an overestimation of Mw.25 Accordingly, the values listed in Table 2 represent upper limits.) The relative molar mass of solution aged specimens were evaluated using either an Agilent HP1100 or 1200 HPLC system fitted with two PLgel 10um MIXED-B columns and a UV detector (270 nm wavelength), operated with dimethylformamide (containing 0.05 M LiBr) as the mobile phase at a flow rate of 1 mL min−1 at 53 °C. The instrument was calibrated with 11 low dispersity polystyrene (PS) standards ranging from 600 to 600 000 g mol−1 obtained from Polymer Laboratories. Small-Angle X-ray Scattering. Small-angle X-ray scattering (SAXS) data were acquired using the 2 m instrument located in the College of Science and Engineering Characterization Facility at the University of Minnesota. Measurements were taken at a 58 cm sampleto-detector distance using Cu Kα radiation (λ = 1.542 Å). The scattering patterns were collected on a two-dimensional multiwire detector and corrected for background and detector response characteristics. Azimuthal averages produced one-dimensional plots
of intensity (arbitrary units) versus the scattering wave vector q = 4πλ−1 sin(θ/2), where θ represents the scattering angle. Fresh and aged compression-molded samples were assessed both in the dry state and following exposure to deionized water for three days. Samples immersed in water were loaded into cells containing Kapton windows separated by an O-ring spacer. Sample temperature was controlled to within ±1 °C during data acquisition. Dynamic Mechanical Spectroscopy. A TA Instruments Q800 DMA was employed for measurements in the tensile mode, operated at 1 Hz and 0.1% strain amplitude. Strain sweeps established the limit of linear response, and all data were acquired subject to this limitation. Dynamic elastic (E′) and loss (E″) moduli were recorded while heating the sample from −150 to 180 °C at 3 °C min−1. An ARES strain-controlled rheometer fitted with 25 mm diameter parallel plates was used to determine the dynamic shear moduli, G′ and G″, as a function of temperature (60−220 °C) and frequency (0.1 ≤ ω ≤ 100 rad s−1) in the linear viscoelastic limit. Samples were loaded onto the fixtures at room temperature and heated to 220 °C (any bubbles present were removed by gentle shearing), and the shear moduli were measured as a function of frequency. Subsequently, the samples were cooled at 3 °C min−1 from 220 to 80 °C for E2A or 60 °C for P35 while being probed at a frequency of 1 rad s−1. The material was then heated back to 220 °C at 3 °C/min while continuously sampling G′ and G″ at 1 rad s−1. A second frequency sweep was then acquired at 220 °C. Tensile Testing. The tensile properties of E2A and P35 were measured as a function of PBS exposure time and temperature according to ASTM D1708 (Rev. B). Between 10 and 15 samples were evaluated at each test condition. Experiments labeled as “initial” were prehydrated in water for three days prior to testing. All samples were removed from the test solutions and maintained in a deionized water bath for up to 1 h at 37 °C prior to testing at 37 °C. Samples were pulled to failure on a MTS Sintech 1/D tensile tester at 5.0 in min−1. The load values were recorded as a function of displacement and converted to engineering stress and strain based on the initial crosssectional area and the gage length of the tensile bars. The total duration of each tensile test did not exceed 2 min, and thus dehydration of the test samples was not a factor.
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RESULTS AND ANALYSIS Molar Mass Determination. The condensation polymerization utilized in the preparation of PDMS containing PEUs is expected to result in a most probable distribution of chain lengths (Đ = 2), and previous publications dealing with related PDMS-based PEUs confirm such behavior based on SEC measurements.26,27 We were unable to find reports with absolute number-average (Mn) or weight-average (Mw) molar mass measurements; these quantities are routinely listed with respect to polystyrene (PS) calibration standards.13−15,18,26,27 (The SEC determination of Đ is more reliable due to the power-law (i.e., Mark−Houwink) dependence of chain hydrodynamic volume with respect to molar mass).25 Yet ultimate mechanical properties such as tensile strength and strain at failure are critically dependent on the absolute chain lengths, reflected in Mn and Mw. Therefore, we determined Mw for E2A and P35 using light scattering measurements conducted during simultaneous SEC analysis (Figure 2). Table 2 lists the absolute and PS-calibrated Mn, Mw and Đ values for both materials. 9112
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consistent with the gradual formation of inhomogeneities larger in size than the nanoscale morphology (see below). Small-Angle X-ray Scattering. Nanoscale morphology was probed using small-angle X-ray scattering experiments performed on freshly molded dry specimens at 25 °C and wet specimens between 25 and 85 °C following saturation with water for three days (see Figure 3). Single broad peaks corresponding to average segregated domain spacings of d = 7 nm for E2A and 6 nm for P35 independent of temperature are shown in Figure 4 (d = 2π/q*, where q* is the peak scattering
Figure 2. SEC traces of (a) E2A and (b) P35. Absolute weight-average molar masses were determined while recording these chromatographic traces using a separate MALLS detector.
Within experimental uncertainty the PS-based Mn result for E2A and dispersity agree with those cited previously.13 The results obtained for P35 are quite similar to E2A. However, the absolute molar mass for E2A (Mw = 101 kg/mol) and P35 (Mw = 88 kg/mol) are roughly half the PS-calibrated values. The values of molar mass dispersity for both samples approached 2, consistent with a condensation polymerization process. Water Uptake. Immersion of molded E2A and P35 in phosphate buffered saline (PBS, see Experimental Section for details) leads to absorption of water to a nearly constant weight within several hours in both samples. Figure 3 illustrates the weight percentage of water in these materials as a function of exposure time at 37 °C and the steady-state fraction of water as a function of temperature. Modification of PEU with PDMS reduces the hygroscopic character of these materials, nearly in proportion to the concentration of hydrophobic PDMS blocks (see Table 1); i.e., E2A and P35 take up slightly more than onethird and one-half, respectively, as much water as the PEU Elasthane 80A. The amount of water absorbed is weakly dependent on temperature, increasing by 10−15% between 37 and 85 °C in both E2A and P35. E2A and P35 both lost transparency with long-time exposure to PBS; for a given time specimens developed greater opacity at higher temperatures (see Supporting Information). Drying opaque samples (E2A and P35 after 52 weeks at 85 °C) under reduced pressure at room temperature returned clarity to the original as-molded levels. Also, there was no noticeable impact on the associated small-angle X-ray scattering patterns
Figure 4. SAXS patterns from E2A (upper) and P35 (lower) obtained from dry and wet (following three days of soaking in distilled water) specimens at various temperatures.
wave vector). These data have been normalized to a common high-q intensity level to account for variations in the overall scattering intensity attributable to variable X-ray beam transmission. These domain spacings are consistent with those reported in previous publications dealing with similar PDMS-based PEU.14,26 Aging these materials in PBS at elevated temperature had a minor impact on the morphology as indicated by the SAXS patterns shown in Figure 5 for specimens of E2A and P35 following 32 weeks of soaking in PBS at 70 °C. Both materials exhibit a 15% increase in d-spacing over the initial wet value but without noticeable changes in the breadth of the SAXS peaks. Subsequent drying at reduced pressure and room temperature to a constant weight had no effect on the slightly expanded domain dimension, which is also shown in Figure 5. Drying and
Figure 3. Water absorption determined gravimetrically as a function of (a) temperature following three day exposures and (b) time at 37 °C. 9113
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Figure 5. SAXS results obtained from (a) E2A and (b) P35 in the initial wet state and following aging in PBS at 70 °C. Also shown are data recorded after drying a specimen aged at 70 °C followed by reprocessing (molding in the dry state at 220 °C) and rehydration. All data were acquired at 25 °C.
Figure 6. Linear dynamic elastic (E′) and loss (E″) moduli measured in tension at a frequency of 1 Hz for (a) E2A and (b) P35 while heating at 3 °C min−1. Glass transition temperatures are indicated by the arrows.
220 °C remolding of both elastomers after 32 weeks in PBS at 70 °C leads to slight (ca. 6%) decreases in d as compared to the initial wet samples. Modest broaden of the peak was also observed. These SAXS results suggest that aging in water at elevated temperatures does not significantly influence the disorganized morphology present following processing at 220 °C. Dynamic Mechanical Spectroscopy. The linear dynamic elastic and loss moduli were determined in tension (E′ and E″) and shear (G′ and G″) as a function of frequency and temperature for dry specimens of E2A and P35 that were molded at 220 °C. Figure 6 shows the E′ and E″ results obtained at a frequency of ω = 1 Hz while heating the specimens at 3 °C min−1. Both materials display peaks in E″ attributable to the glass transitions of domains formed (i) predominately by PDMS around −100 °C for both samples and (ii) at −20 °C for the PHMO containing E2A and at 0 °C for the PTMO containing P35, in agreement with literature values.28 Above about 170 °C the materials become too soft for accurate measurements in the tensile geometry. Complementary parallel plate experiments for determination of G′ and G″ were conducted between 60 and 220 °C to characterize the rheological behavior of the materials up to and including the standard temperature at which the material is processed. Specimens were loaded between the rheometer plates and quickly (ca. 10 min) heated to 220 °C under a purge of dry nitrogen followed by acquisition of frequency spectra as shown in Figure 7a,b. At this temperature E2A displays terminal liquid-like behavior (G′ ∼ ω2 and G″ ∼ ω) while P35 exhibits
nonterminal G′(ω) behavior at the lowest measurement frequencies. The low frequency response of G′ for P35 is consistent with some degree of micro- or macrophase separation.29 This assessment may explain anecdotal evidence of difficulties encountered when processing P35 and is consistent with the cloudy appearance of the molded samples (see Supporting Information). Cooling from 220 °C at 3 °C min−1 (following the isothermal frequency scan) was accompanied by the acquisition of isochronal (ω = 1 rad s−1) G′(T) data (Figure 7c,d). The dynamic elastic modulus increases monotonically with both materials down to about 100 °C, and then G′ rises sharply between 100 and 90 °C to a level consistent with the results in Figure 6 (i.e., E′ ≈ 3G′). This increase in G′ upon cooling is attributed to demixing of the constituent blocks. Upon reheating from 60 °C (at 3 °C min−1) the G′ data show no evidence of the 100 °C transitions seen during cooling. For E2A, the rate of decline in elasticity with temperature increases at roughly 140 °C, displays a slight inflection around 180 °C, and then drops precipitously at 200 °C. Similar behavior was observed for P35; however, the G′ data plateau to some extent above 200 °C. A second frequency scan reveals that the E2A material has returned precisely to the same liquid-like rheological state that it exhibited at the start of the cooling and heating cycle and P35 returns to a similar nonterminal scaling behavior with a modest offset in the magnitude of G′(ω) and G″(ω) as seen in Figure 7b. In P35 the variation in G′(ω) and G″(ω) at 220 °C may reflect additional macrophase separation as the polymer was noticeably more opaque when 9114
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Figure 7. Dynamic elastic (G′) and loss (G″) moduli measured in shear for E2A (a, c) and P35 (b, d). Frequency spectra were obtained after quickly heating to 220 °C (filled symbols, a and b) and then again following isochronal (ω = 1 rad s−1) measurements shown in panels c and d taken while cooling the specimens at 3 °C min−1 to 60 °C and reheating at 3 °C min−1 to 220 °C (open symbols, a and b). E2A displays a reproducible liquidlike terminal behavior at 220 °C, while P35 exhibits a nonterminal response at this temperature.
measured ultimate properties, and we suspect this behavior derives from macroscopic heterogeneities that are also reflected in sample cloudiness (see Supporting Information) and nonterminal flow (Figure 7b) during processing. Both dry materials are relatively tough elastomers (Figure 8a,d) that strain-harden and fail with similar σT (∼20−25 MPa) and εb (∼500−750%) values. E2A is stiffer (see also Figures 6 and 7) and displays a smaller σy than P35. Absorption of water reduces the initial (nonaged) values of σy and εb in E2A (Figure 8a) but has little impact on P35 (Figure 8d). Most obvious is the dramatic loss of strength and toughness (i.e., area under the stress strain response) that occurs with increased time and temperature of exposure to PBS. After 52 weeks at 85 °C, both materials no longer strain-harden and σy and εb are reduced to less than half their initial values. Loss in mechanical integrity is more severe with P35 than E2A. Drying and reprocessing (molding) aged materials does not lead to recovery of the initial wet mechanical properties as illustrated in Figure 8b,e (reprocessed curves) for 32 weeks at 70 °C. Molar Mass Changes. We monitored the molar mass distribution of the PDMS-based PEU elastomers at all stages of aging by PS-calibrated SEC measurements. Figure 9 shows a representative set of results obtained with E2A over a one year period at 85 °C. Increased time in PBS solution leads to a shift in the SEC curves to higher elution volumes consistent with a lowering of the molar mass of the samples. (These SEC traces show no evidence of allophanate formation,30 which would result in a high molar mass shoulder; see Supporting Information.) Figure 10 summarizes how the PS-relative values
removed from the instrument at the end of the measurement cycle. These rheological results are consistent with microphase separation of soft (PDMS and PH(T)O rich) and hard (MDIBDO rich) blocks leading to a disorganized three-domain nanoscale morphology. We associate the pronounced hysteresis in G′(T) between about 100 and 180 °C with mixing (upon heating) and demixing (upon cooling) of the constituent blocks, reflecting the aforementioned packing frustration created by the lack of regularity in molecular architecture imposed by condensation polymerization of the principal building blocks. Nevertheless, desirable elastomeric properties are established upon cooling below 100 °C and persist upon heating to about 140 °C. We also can conclude that these materials, particularly E2A, exhibit thermal stability under dry conditions for periods of at least 30 min at 220 °C, enabling facile melt processing. Tensile Properties. Tensile testing was performed at 37 °C on ASTM D1708 dog-bone-shaped specimens of E2A and P35 following exposure to PBS for up to one year at 37, 55, 70, and 85 °C. Between 10 and 15 individual engineering stress (σ) versus engineering strain (ε) experiments were performed at each combination of time and temperature, and representative results are presented in Figure 8. All the experiments on E2A yielded stress−strain curves closely clustered around average responses as illustrated by Figure 8a−c; the variability in yield stress (σy), tensile strength (σT), and strain at break (εb) is typical of thermoplastic elastomers.10 As seen in Figure 8d,e, P35 produced significantly broader distributions in the 9115
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Figure 8. Representative stress−strain results measured at 37 °C for E2A (a, b, c) and P35 (d, e, f) following aging in PBS at representative temperatures and times. (A complete collection of results is provided in the Supporting Information.) Data shown in red were obtained from initial dry specimens, and the results shown in blue illustrate the tensile behavior following drying, reprocessing (molding at 220 °C), and rehydration of specimens aged for 32 weeks at 70 °C. The data in black were acquired for wet samples.
the changes in molar mass based on conventional chemical reaction rate theory and consider how this behavior is connected to loss of mechanical integrity. Molar Mass Reduction. The molar mass reduction observed for both E2A and P35 upon aging in PBS at various temperatures depicted in Figure 10 can reasonably be interpreted as due to hydrolytic cleavage of the segmented polyurethane backbone as water (and associated salts in the buffer) is the only reagent present in the aging studies, and the materials take up between 0.5% and 1% water by mass (Figure 3) under all aging conditions. Consistent with random chain cleavage the originally broad molar mass dispersity remained relatively constant over the course of the aging studies.31 Every hydrolysis event in a given macromolecule will lead to two chains with a number-average average molar mass equal to onehalf of the original value Mn(0). A second hydrolysis event will lead to a one-third reduction and so on. The rate of this molar mass reduction can be related to the hydrolysis reaction kinetics. A simple reaction rate analysis suggests that the rate of backbone bond (BB) loss (i.e., hydrolysis) can be given by eq 1, where [BB] and [H2O] are the effective concentrations of hydrolytically sensitive backbone bonds and water in the sample during aging, respectively.
Figure 9. SEC results for E2A aged in PBS at 85 °C. The chromatographic response shifts to larger elution times as a function of aging time, indicative of a reduction in the molar mass. Solid symbols identify the elution times of polystyrene (PS) calibration standards.
of Mn change with time as a function of aging temperature for both materials.
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DISCUSSION The results presented in the previous section demonstrate that exposure of E2A and P35 to PBS at all temperatures between 37 and 85 °C for extended periods of time results in a steady reduction in molecular mass with concomitant degradation of the ultimate tensile properties. In this section we consolidate
−
d[BB] = k[BB][H 2O] dt
(1)
As the maximum molar mass reduction we observed in Figure 10 is approximately one-third of the original value, the concentration of both water and backbone bonds is 9116
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Figure 10. Mn determined by SEC (PS calibrated) versus aging time at various temperatures for (a) E2A and (b) P35.
Figure 11. Mn(0)/Mn(t), obtained from Figure 10, plotted versus time for (a) E2A and (b) P35, where Mn(0) is the initial molar mass. The slope of the indicated lines is proportional to k′ for each temperature as defined in eq 4.
temperatures. The linear dependence of Mn(0)/Mn(t) on time is evident and consistent with this simple model. Values of k′ [EG]0−1 at the various aging temperatures can be extracted from the slopes of the lines fitted to the data in Figure 11. The fundamental second order rate constant k is embedded in k′, and an estimate of k would require knowledge of [EG]0, [BB]0, and [H2O]0. While the overall average value of [H2O]0 can be calculated from the data in Figure 3, we expect a nonuniform local distribution of water in the sample; thus, the effective value of [H2O]0 is unclear. Also, we are not certain as to the specific backbone bond that is most susceptible to hydrolysis in either elastomer; therefore, [BB]0 is also unknown. Thus, we are reluctant to estimate actual values of the second-order rate constant for hydrolysis k from this analysis. Nonetheless, we can evaluate the temperature dependence of k using an Arrhenius analysis. Assuming that the hydrolysis is an activated process (i.e., the simple second-order chemical reaction represented by eq 1), the temperature dependence of k′ should be given by eqs 5 and 6, where Ea is the apparent activation energy, R is the gas constant, T is the absolute temperature, and A is the Arrhenius prefactor.
insignificantly changed from their original values even at the longest reaction times. Thus, eq 1 can be simplified to the pseudo-zero-order reaction rate shown in eq 2, where k′ = k[BB]0[H2O]0. −
d[BB] 1 d[EG] = = k′ dt 2 dt
(2)
Every backbone bond cleavage generates two additional end groups (EG) in the sample with the rate of EG formation being twice the rate of backbone cleavage, thus the factor of 1/2 in eq 2. The concentration of end groups ([EG]) in a macromolecular sample is inversely proportional to the numberaverage molar mass; thus, the reduction in molar mass for a given polymer sample can be related to the end-group concentration in that sample according to eq 3.
[EG]t M n(0) = M n (t ) [EG]0
(3)
Integration of eq 2 and substitution into eq 3 give a prediction for the dependence of molar mass reduction as a function of time for aging of these polyurethanes under hydrolysis conditions as eq 4.
M n(0) 2k′ = t+1 M n (t ) [EG]0
k′ = k[BB]0 [H 2O]0 = A e−Ea / RT [BB]0 [H 2O]0
(5)
(4)
The molar mass data are plotted in Figure 11 as a function of time according to eq 4 for both E2A and P35 at various
ln k′ = ln(A[BB]0 [H 2O]0 ) − 9117
Ea RT
(6)
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The values of k′ are embedded in the slopes (= 2k′[EG]0−1, eq 4) of the lines in Figure 11. Since [EG]0 = 2ρMn(0)−1 for linear polymers (ρ is the polymer density), the measured slope is equal to k′Mn(0)ρ−1. Accordingly, the slopes taken from Figure 11 are plotted as (ln k′ + C) versus T−1 in Figure 12, where C =
Figure 13. Master curve for the hydrolysis of (a) E2A and (b) P35 based on application of the shift factor bT (eq 7) with a reference temperature of 37 °C. Figure 12. Arrhenius plot for the results obtained from Figure 11 and based on eq 6 for (a) E2A and (b) P35.
one covalent bond, i.e., splitting XPX into XP, has virtually no effect on the morphology but dramatically influences the ultimate tensile properties. XP strain-softens beyond about 200% extension and fails at approximately one-third the tensile strength of XPX. These characteristics guide our interpretation of the changes in mechanical behavior with the PDMS-based PEUs. E2A and P35 contain a more complex mix of molecular structures than XPX as described in the Introduction. Nevertheless, an underlying multiblock molecular architecture coupled with a disorganized nanoscale morphology explains the initial stress−strain response of these materials and the similarity to XPX. By analogy, the observed correlation between reduced molar mass and the deterioration in mechanical properties in PDMS-based PEUs is not surprising. The overall molar mass influences the tensile strength and strain at break of most polymeric materials.34,35 (Of course, these properties also are affected by a host of other factors, e.g., whether the material is glassy, semicrystalline, or rubbery, the entanglement molecular weight (Me), extent of cross-linking, state of microphase separation, etc.) In the limit of many individual hard and soft block sequences per chain (i.e., high overall molar mass) or where the molar mass is much greater than the entanglement value Me, sacrificing one backbone connection (i.e., statistically halving Mn) would not necessarily produce the results found in Figure 8. However, E2A (Figure 8a) and P35 are not high molar mass compounds (compared to most commercially available elastomers and plastics) with respect to the ultimate mechanical properties, even in the initial pristine form (Table 2). Halving the average chain length of either material brings the number-average molar mass below levels normally associated with high-strength elastomers.36
ln(Mn(0)ρ−1). The above analysis assumes that the initial concentrations of [EG]0, [BB]0, and [H2O]0 are not strong functions of temperature. This is reasonable as the density of polymers are not greatly dependent on temperature; thus, [BB]0 and [EG]0 should be constant, and we determined that the average water contents in these materials are not strongly temperature dependent (Figure 3). The approximately linear dependence in Figure 12 is consistent with an activated chemical process, and the values of Ea for the two polymers extracted from this analysis (∼30 kJ mol−1) are within experimental error of one another.32 With the activation barriers from the above analysis, the reduction in molar mass data shown in Figure 11 can be shifted and plotted on a master curve at a particular reference temperature (see Supporting Information). The time values (t) for the Mn(0)/Mn(t) data in Figure 11 at 55, 70, and 85 °C were shifted to a new reference time (tref) by an acceleration factor (bT) and plotted on a 37 °C master curve (Tref = 37 °C or 310 K) in Figure 13 using the relationship shown in eq 7. tref = bT t
where bT = e(Ea / R)(1/ Tref + 1/ T )
(7)
The values of Ea taken for the E2A and P35 were 36 and 24 kJ mol−1, respectively (Figure 12). This analysis allows a prediction for the long-time molar mass reduction for both E2A and P35 at 37 °C using data acquired at higher temperatures. Mechanical Properties. Molecular architecture plays a pivotal role in the control of morphology and mechanical performance in multiblock polymer materials. A recent study of CEC-P tetrablock and CEC-P-CEC heptablock terpolymers (also referred to as XP and XPX, respectively) underscores this point.33 These model thermoplastic elastomers, formed by sequencing glassy poly(cyclohexylethylene) (C), semicrystalline poly(ethylene) (E), and rubbery poly(ethylene-altpropylene) (P), are strikingly similar to PDMS-based PEUs. They form a disorganized three-domain microphase-separated morphology containing hard (C) and soft (E and P) domains mechanically coupled by covalent bonds. The tensile stress− strain response of XPX mimics the initial strain-hardening behavior of E2A (Figure 8a) and P35 (Figure 8d). Breaking just
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PERSPECTIVE While we do not know precisely which backbone bonds break when PDMS-based PEUs are exposed to water, we can reasonably correlate the observed reduction in ultimate properties with the reduction in molar mass. Thus, determination of the specific bonds being broken in these elastomers under hydrolytic aging will allow for the design of more hydrolytically resistant materials and presumably more mechanically stable materials. Searching for specific end groups 9118
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reaction rate at low temperature and vice versa. In the studies described above, we determined Ea for the hydrolysis reaction of these polyurethanes. From Figure 13 the long-term degradation of polymer molar mass at 37 °C can be predicted based on the high-temperature rate data acquired over relatively short times. This analysis suggests that the maximum reduction in molar mass (and tensile properties) we observed at 85 °C and 52 weeks would be achieved at 37 °C in approximately three years for P35 and approximately six years for E2A. The associated mechanical performance in actual device applications will of course depend on how the material is loaded in a given design and the specific in vivo environment. These results clearly show that the maximum sustainable load decreases with exposure time to water (Figure 8).
that result from the hydrolytic cleavage is one approach. However, end-group concentrations are small for these samples, thus rendering their detection challenging. Studying the hydrolysis of appropriately designed model compounds that are representative for the linkages in these PDMS-based PEU ealstomers is an alternative strategy that may shed light on this problem. Nonetheless, from our studies we have determined an apparent activation barrier for the hydrolytic process of ∼30 kJ mol−1, and the Arrhenius approach we have taken is consistent with other work in the literature on the hydrolytic stability of polyurethane elastomers.37,38 This activation energy can give us some clue as to the nature of the hydrolysis event. However, given the assumptions in the analysis, the unknown influence of the degradation products on the rate of subsequent hydrolysis events, and the statistical error associated with our approach, we are careful to not over interpret the activation energy. The bonds that are most susceptible to hydrolysis in the polyurethane are the Si−O of the PDMS segments, the C−O of the polyether segments, and a specific polyurethane connection (N(H)−C(O)−O). We suspect that the urethane bonds adjacent to either a PH(T)MO chain or a PDMS chain are the most likely candidates as the hydrolytic stability and degradation in mechanical properties of simple PTMO/BDO/ MDI polyurethanes has been studied and hydrolysis of the PTMO/MDI carbamate has been implicated.39 Comparison of the activation energy we determined to the activation energies for carbamate hydrolysis would be one way to support this conjecture. However, from our initial search of the literature, the activation energy we determined for the hydrolysis of these PDMS-based PEUs appears to be anomalously low compared to other small molecule carbamates.40,41 While incorporation of PDMS into these materials reduces the water uptake and thus the effective concentration of water in the samples, the high degree of incompatibility between the PDMS segments and the polar carbamate linkage may induce stretching of the PDMS blocks near the hard segment interface. Cleaving the carbamate linkage adjacent to the PDMS chain would result in a release of this strain, possibly a driving force for hydrolysis and a potential reason for the relatively low activation energy calculated. Of course, we recognize that these two commercial polyurethanes are used in implantable biomedical devices that must endure long-term exposure to physiological conditions. In such aqueous and oxidative environments, degradation of molar mass in such elastomers should be of principal concern when evaluating and predicting long-term performance characteristics. Our work here suggests that the mechanical integrity of these PDMS-based PEUs may be compromised in such applications (e.g., as insulation on pacemaker/defibrillator leads) as a result of simple hydrolysis of the segmented polyurethane macromolecules. Thus, while the oxidative stability of these PDMS-containing elastomeric polyurethanes is enhanced and the level of water uptake is reduced, our results suggest that the hydrolytic cleavage of these polymers is not eliminated and raises cause for concern in long-term implantable biomedical devices. Predicting long-term materials performance from accelerated testing can be fraught with difficulties. This is especially true when evaluating the failure behavior as defects can dominate the ultimate properties. However, for simple chemical reactions that follow Arrhenius rate dependences, the correlation between reaction rate and temperature is well established for myriad reactions.42,43 With the apparent activation energy (Ea), the reaction rate at high temperature can be predicted from the
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SUMMARY Two commercial PDMS-based PEU thermoplastic elastomers, Elast-Eon E2A (denoted as E2A) and PurSil 35 (denoted as P35), were comprehensively investigated for changes in molar mass and ultimate mechanical properties in tension following exposure to a deoxygenated phosphate buffered saline solution for up to one year at temperatures between 37 and 85 °C. Our results demonstrate a degradation in molar mass with increased aging time and temperature. The molar mass reduction observed follows an Arrhenius relationship, presumably due to hydrolytic cleavage of the polymer chains. P35 absorbs nearly twice the amount of water as E2A, and the molar mass degrades at almost twice the rate. Reduction in molar mass is accompanied by substantial degradation of the tensile strength and toughness of both materials. Additional research is necessary to firmly establish the specific reactions responsible for these results.
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ASSOCIATED CONTENT
S Supporting Information *
Allophanate analysis, sample appearance upon exposure to water, complementary electron microscopy experiments, complete tensile testing data, and Arrhenius acceleration factor analysis. This material is available free of charge via the Internet at http://pubs.acs.org.
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AUTHOR INFORMATION
Corresponding Author
*E-mail: kim.chaffi
[email protected] (K.A.C.); hillmyer@ umn.edu (M.A.H.);
[email protected] (F.S.B.). Notes
The authors declare the following competing financial interest(s): Frank S. Bates and Marc A. Hillmyer are paid consultants for Medtronic Incorporated.
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ACKNOWLEDGMENTS Portions of this work were carried out at the Characterization Facility in the College of Science and Engineering at the University of Minnesota, which receives partial support from the National Science Foundation through the NNIN program and the Materials Research Science and Engineering Center (NSF-MRSEC) at the University of Minnesota (NSF-DMR0212302). We acknowledge Justin Kennemur for acquiring the SEC data presented in Figure 2. The authors thank Timothy P. Lodge for his critical reading of this manuscript and helpful comments. The authors also thank SuPing Lyu for helpful discussions concerning experimental design and data analysis. 9119
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