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Viscoelastic and gas transport properties of a series of multiblock copolymer ionomers
Polymer 52 (2011) 3963e3969
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Viscoelastic and gas transport properties of a series of multiblock copolymer ionomers Yanfang Fan a, Mingqiang Zhang b, Robert B. Moore b, Hae-Seung Lee b, James E. McGrath b, Chris J. Cornelius a, * a b
Department of Chemical Engineering, University of Connecticut, Storrs, Connecticut 06269, USA Department of Macromolecular Science and Engineering, Virginia Polytechnic Institute and State University, Blacksburg, VA 24061, USA
a r t i c l e i n f o
a b s t r a c t
Article history: Received 2 April 2011 Received in revised form 17 June 2011 Accepted 25 June 2011 Available online 1 July 2011
A series of sulfonated poly(arylene ether sulfone) (BPSH-BPS) multiblock copolymer (MBC) ionomers were studied with respect to structure and physical properties. Ion exchange capacity (IECw) on a weight basis was constant for all MBC ionomers with controlled ionic and nonionic block lengths of 5k:5k, 10k:10k and 15k:15k. Viscoelastic properties were investigated as a function of block length, relaxation time, and temperature ranging from 200 C to 250 C. Relaxation time and glassy state modulus decreased with increasing ionic block length. Williams-Landel-Ferry (WLF) theory adequately modeled the viscoelastic changes in the BPSH-BPS MBC ionomers due to changes in time and temperature using a shift factor aT. Small Angle X-ray Scatter (SAXS) revealed an average inter-chain d-spacing that increased with increasing ionic block length (21.1, 31.4 and 36.4 nm). The larger ionic domain size is attributed to an increase in bulk ionic density within an ionic domain. Bondi’s group contribution method predicts a lower fractional free volume (FFV) with increasing ionic block length. The 15k BPSHBPS MBC ionomer had the lowest FFV of the series, which corresponds to the lowest hydrogen permeability of 2.46 Barrers at 30 C and 4 atm. Ó 2011 Elsevier Ltd. All rights reserved.
Keywords: Ionomer Viscoelasticity Transport
1. Introduction Today, fuel cells are considered to be a promising technology for alternative energy for stationary power and transportation [1,2]. Proton Exchange Membrane Fuel Cells (PEMFC) are being studied due to this diverse energy potential [3]. While PEMFC technology has the potential to be environmentally friendly, a critical property needing to be solved is long-term durability. Research studies have shown that several factors can reduce the lifetime of a PEMFC, including corrosion of the carbon-support, platinum-particle dissolution and sintering, and membrane thinning due to material relaxation [4]. The challenges of minimizing dynamic mechanical and material stresses leading to fuel cell failures in PEMFC systems is an ongoing area of investigation. These failures are due to inadequate membrane ionomer physical properties resulting in its low durability [3]. One mechanism associated with the physical degradation of membrane ionomers is operational cycling due to variations in temperature and humidity. These variations cause * Corresponding author. Tel.: þ1 860 486 3686; fax: þ1 860 486 2959. E-mail address: [email protected] (C.J. Cornelius). 0032-3861/$ e see front matter Ó 2011 Elsevier Ltd. All rights reserved. doi:10.1016/j.polymer.2011.06.050
changes in the hydration conditions of the ionomer, which contributes to the formation of micro cracks during volumetric expansion and contraction. Once a micro-crack forms, enhanced mechanical stresses accelerate crack propagation, which results in a significant decrease in the durability of a proton exchange membrane (PEM) fuel cell, and its failure as a membrane separator [5,6]. Generally, fuel cell durability is highly correlated to viscoelastic properties of PEMs, which are timeetemperature dependent physical properties. Many researchers have focused on the investigation of PEMs viscoelastic properties. Lai et al. proposed a linear viscoelastic stress model to account for time, temperature, and humidity effects in polyfluorosulfonic acid (PFSA) membranes [7]. The results showed that large variations in humidity create a faster hydration/dehydration rate resulting in high residual tensile stresses. Patankar et al. performed relaxation tests using DMA on Gore-Select 57 PEMs [8]. Thermal and hygral master curves were constructed using time, temperature, and humidity to create superposition plots based on the modulus. This model predicts the glassy plateau modulus (Eg) and stresses only for PSFA type membranes. Hydrocarbon based ionomers are one class of promising alternative membrane for fuel cells [9]. Research from the McGrath
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O
O HO3S
O S O
SO3H O A
F
F
O
O F
O
F
O S O
Bn
Fig. 1. Chemical structure of BPSH-BPS MBC ionomers.
group has focused upon the development of poly(arylene ether sulfone) MBC ionomers. The acronym for these materials is based on BP representing Biphenol, S for sulfone, H for the acidified form, and xx is the degree of disulfonation (BPSH-xx or BPS-xx) [9]. A series of BPSH-BPS (x:y) MBC ionomers with x and y representing the block length of the oligomer is shown in Fig. 1. When a volumetric based ion exchange capacity (IECv) is used to compare proton conductivities, these materials are comparable to Nafion. In this work, time dependent modulus responses of BPSH-BPS MBC ionomers is correlated with physical properties and molecular structure. Morphology, viscoelasticity, and gas transport are studied with respect to ion and non-ion containing block lengths of 5k:5k, 10k:10k, and 15k:15k. The effects of structure and block length with respect to changes in viscoelastic and hydrogen transport properties are discussed. 2. Experimental 2.1. Materials Synthesis of BPSH-BPS MBC ionomers has been documented in detail by Hae-Seung et al. [9]. BPSH-BPS was synthesized via a coupling reaction between phenoxide terminated BPSH100 oligomers and HFB end-capped BPS0 oligomers. As for the hydrophilic oligomer, BPSH100 was synthesized by varying the molecular weight via step growth polymerization. BPS0 of hydrophobic oligomer was first synthesized, and then end-capped with HFB via a nucleophilic aromatic substitution reaction. Nafion (1100 EW, sulfonic acid form, Hþ) was purchased from VWR LLC as a comparison with BPSH-BPS copolymers. Sulfonic acid (2 mol/L) was purchased and used as received from SigmaeAldrich.
form by boiling in 0.5 M H2SO4 for 2 h, rinsing in DI (deionized) water, and then boiling in DI water for 2 h to remove any excess acid. Wet samples were dried under vacuum at 60 C for 3 h.
2.3. Thermal analysis The glass transition temperature for Hþ form BPSH-BPS MBC ionomers was obtained using Dynamic Mechanical Analyzer Q800 (DMA, TA instruments). Membrane samples were analyzed in the tensile mode. Material tests were done using a frequency of 1 Hz with a heating rate of 2 C/min from 50 C to 250 C and a tensile strain 3t of 0.1%.
2.4. Short-term stress relaxation tests Frequency dependent relaxation tests for BPSH-BPS MBC ionomers were performed over a temperature range of 200e250 C using a Q800 DMA. A static force (stress) of 0.01 N was used in all experiments. This stress was experimentally determined to give a linear viscoelastic response. Materials were converted to the Hþ form prior to testing.
2.5. Gas transport properties Hydrogen permeability through dry films was measured using the salt form of a BPSH-BPS MBC ionomer. A gas permeation system was custom designed and built to determine gas transport properties using the time-lag method [10,30]. The gas purity used in this study was 99.99%. All experiments use a feed pressure of 4 atm and varied the temperature from 30 to 90 C.
2.2. Membrane preparation The salt form of BPSH-BPS MBC copolymers were dissolved in NMethylpyrrolidone (NMP) yielding transparent solutions (5e7 wt.%), which were solution cast on Teflon or glass substrates. Membrane films were created by controllably evaporating solvent from the polymer solution. Solvent removal was done in multiple steps involving drying for in a vacuum for 24 h at 60 C with an additional vacuum drying step at 110 C for 48 h. Membranes were solution cast into films having a nominal thickness ranging from 0.020 to 0.070 mm. Membrane films were converted to their Hþ
2.6. Density measurement A hydrostatic weighing method was used to measure the density of a BPSH-BPS MBC ionomer. Dry BPSH-BPS MBC ionomer films were weighed in the air a total of three times to yield an average sample mass w0. w is the films buoyant mass measured in hexane. The polymer density r can be obtained by the following relationship using hexane density rhexane and a small correction for the density of air rair (Eqn. (1)). The average error is 0.8%.
Fig. 2. Temperature dependence of BPSH-BPS MBC ionomers in Hþ form versus (a) tan d and (b) storage modulus E0 .
Y. Fan et al. / Polymer 52 (2011) 3963e3969
qMAX
5
40
15
d1
2qMAX 4
d 1 (nm)
Log I(q) [a.u.]
10
15k
2.5
IECV
32
3.0
2.0 24
2
10k
1.5
5k
16 5
0 0 .1
0.3
10
15
3
1.0
Molecular Weight × 10 (g/mol)
1
-1
IEC V (mol/L)
6
3965
q (nm )
Fig. 4. Dependence on block length versus d-spacing (square) and IECV (circle).
Fig. 3. SAXS spectra of BPSH-BPS MBC ionomers.
w0 þ rair *r w0 w hexane
(1)
2.7. Small angle x-ray scattering Small angle X-ray scattering (SAXS) using a Rigaku S-Max 300 3 pinhole SAXS system was used to evaluate structure. The radiation X-ray source is CuKa. The incident X-ray beam was attenuated to a wave length of 0.154 nm with a sample-to-detector distance of 16.0 cm. 3. Results and discussion 3.1. Physical properties A series of BPSH-BPS MBC ionomers were evaluated using DMA to determine if block length effects chain mobility and the glass transition temperature (Fig. 2). A single glass transition temperature (Tg) was observed for the MBC ionomers in the Hþ form that is attributed to the regularity of block length, fixed IEC, and similar phase separation properties. The sulfonated hydrophilic blocks have a Tg varying between 248 Cand 251 C, which is in the range of a random BPSH copolymer ionomer (230 Ce280 C) [11]. BPSHBPS MBC ionomers’ sulfonic acid groups were converted to the salt form to increase thermal stability and partially shield ionic charge interactions. When samples are converted to their salt forms, a lower Tg is observed around 180 C. This thermal transition is associated with the unsulfonated hydrophobic blocks. The observation of a lower Tg is due to a reduction in the concentration of ionic interactions among polymer chains that masks the Tg of the unsulfonated domain. This result is close to commercial polyphenylsulfone with a Tg of 190 C [12]. The primary difference between the two Tgs is attributed to differences in molecular weight Table 1 SAXS domain space comparison. MPC Ionomer
qmax nm1
2qmax nm1
d1 nm
d2 nm
BPSH-BPS (5k:5k) BPSH-BPS (10k:10k) BPSH-BPS (15k:15k)
0.297 0.200 0.173
0.587 0.406 0.356
21.2 31.4 36.4
21.4 31.0 35.3
T-T 0 ( o C) -40
-20
0
20
40
5k 10k 15k WLF equation
5
log(aT )
r¼
of the unsulfonated block length, and the degree and concentration of hydrogen bonding between ionic groups. Previous studies of these materials show that they have a tendency to form a phase-separated morphology. TEM and AFM phase images reported by Badami confirmed that the degree of phase separation is enhanced by an increase in block length [13]. Additional structural information of these materials was evaluated with through-plane measurements using SAXS. The scattering spectra of BPSH-BPS MBC ionomers are shown in Fig. 3 In this figure, I(q) is the scattered intensity reported in arbitrary units (a.u.) and q the scattering vector that is related to the scattering angle 2q [q ¼ (4p/l sin q), l is wave length of x-ray beam]. SAXS spectra of the BPSH-BPS MBC ionomers reveals a maximum scattering peak qmax at a lower scattering angle corresponding to a first order reflection. Assuming that qmax is due to interparticle interferences between ionic cluster domains, this peak can be used to estimate the inter-ionic distance d1. This distance is calculated by relating the maximum scattering vector qmax to Bragg’s law (d1 ¼ 2p/qmax) [14]. The nominal d-spacing between chains was 21.2 nm for 5k, 31.4 nm for 10k, and 36.4 nm for 15k (Table 1). This result suggests that ionic domain size is a function of
5
0
0
-5
-5
200
220
240
260
o
T ( C) Fig. 5. Shift factor aT versus temperature.
280
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a
b 10
10
o
9
10
10
o
E' (Pa)
E' (Pa)
220 C o 225 C o 230 C o 235 C o 240 C o 245 C o 250 C o 255 C o 260 C
10
9
10
8
220 C o 225 C o 230 C o 235 C o 240 C o 245 C o 250 C o 255 C o 260 C
10
1
f (HZ)
5
10
2
10
-1
10
-4
f× aT (HZ)
Fig. 6. a) Frequency versus Temperature of a 5k:5k BPSH-BPS MBC ionomer and b) reconstructed modulus using the shift factors aT.
block length (Fig. 4). A weak secondary order peak at 2qmax is observed for all three MBC ionomers. This scattering behavior suggests the presence of a well-ordered and periodic microstructure that may have a lamellar morphology since d2 is similar to d1. Improvements in ionomer phase separation may lead to better chemical stability and mechanical properties because of hydrophobic domains, while hydrophilic domains may enhance proton transport. These morphology considerations will be discussed with respect to viscoelastic and gas transport properties. The densities of BPSH-BPS MBC ionomers were 1.41, 1.43 and 1.48 g/cm3 that increased with block length. A special feature of these MBC ionomers is a fixed IECw based on mass. Using the polymer density, a volumetric IEC (IECv, mmol/cm3) was defined as the molar concentration of sulfonic acid groups per unit volume of dry membrane [15]. The IECv versus block length in Fig. 4 shows a direct relationship between increasing IECv and block length. This increase in IECv may be indicative of an increase in ionic density due to block length and domain size. Domain size expands with increasing local sulfonic acid group concentration. This results in a greater concentration of expanded chains observed by an increasing d-spacing measured by SAXS. Overall, the materials morphology and structure are tied to block length, which is one crucial factor influencing phase separation, ionic domain size, and ionic density. 3.2. Timeetemperature superposition (TTS) Polymer viscoelastic properties are dependent on time and temperature [16]. Due to this interdependence, changes in properties can be correlated using a timeetemperature superposition (TTS) [17]. This interrelationship of time and temperature has been
a 10
used to successfully characterize the physical properties of many amorphous polymers [18,19]. In this work, TTS is utilized to characterize time and temperature dependent ionomer properties. Mathematically, changes in the modulus (E) due to different temperatures (T1,T2) and time (t1,t2) can be equated to each other using the following relationship E(t1,T1) ¼ E(t2/aT,T2). A time dependent shift factor aT, compensates for viscoelastic changes in the time scale due to temperature. aT is proportional to the polymer chain’s temperature dependence, segmental mobility, frictional diffusion, and reptation [18]. The shift factor aT is used to construct master curves at a reference temperature To in order to reconstruct viscoelastic responses from time t1 to t2. Shift factors are determined based on the Williams, Landel and Ferry (WLF) relationship (Eqn. (2)) [18]. In the WLF relationship, empirical constants C1 and C2 are calculated based upon the reference temperature To and aT. 0 The coefficients C1 and C2 can be appropriately transformed to C1 0 and C2 with Tg being used as the reference temperature (Eqns. (3) and (4)).
Log aT ¼
(3)
C20 ¼ C2 þ Tg To
(4)
The WLF relationship was utilized to study time and temperature dependent physical properties of BPSH-BPS MBC ionomers. A common reference temperature To was chosen for investigating the
0
tan(Delta)
E'(Pa)
BPSH-BPS(5K:5K) BPSH-BPS(10K:10K) BPSH-BPS(15K:15K)
10
9
BPSH-BPS (5K:5K) BPSH-BPS(10K:10K) BPSH-BPS(15K:15K) 10
8
10
6
10
3
10
f×aT (HZ)
0
10
-3
(2)
C1 C2 C2 þ Tg To
C10 ¼
b 10
10
C 0 ðT To Þ C1 T Tg ¼ 0 1 C2 þ ðT To Þ C2 þ T Tg
10
-1
10
-2
10
-6
10
-3
10
0
f×aT (HZ)
10
3
10
6
Fig. 7. Master curves of BPSH-BPS MBC ionomers of different block length using a KWW fit for (a) Storage Modulus and (b) tan d.
Y. Fan et al. / Polymer 52 (2011) 3963e3969 Table 3 WLF constants and Tg.
5k
18
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MPC Ionomer
Tg C
C1
248 251 250
122 122 122
τ (s)
o
BPSH-BPS (5k:5k) BPSH-BPS (10k:10k) BPSH-BPS (15k:15k)
10k
12
"
5
10
fðtÞ ¼ exp
15
Molecular Weight × 10 3 (g/mol) Fig. 8. Characteristic time s versus block molecular weight.
dependence of block length and temperature based on aT. The viscoelastic time and temperature dependent properties for all BPSH-BPS MBC ionomers were fitted with the WLF equation using aT and the empirical constants C1 and C2 listed in Table 3 [19]. A plot of aT versus T and (T To) of different block length materials resulted in nearly identical values and slopes (Fig. 5). This result suggests that block length may not be a significant factor affecting the relationship between aT and temperature. Many researchers commonly use the Tg to compare shift factors between different material systems [16]. In this case, the Tg measured by DMTA was used as the reference temperature. The 0 0 corresponding values of C1 and C2 are listed in Table 3. It has been 0 0 suggested that C1 and C2 are material specific. Because BPSH-BPS MBC ionomers have the same structure, it is not surprising that 0 0 C1 and C2 did not vary significantly with polymer block length [20]. 0 0 Generally, C1 is often related to FFV at Tg and C2 is related to the differences between Tg and the thermodynamically defined transition temperature where free volume is zero [16]. The frequency versus temperature of BPSH-BPS MBC ionomers is shown in Fig. 6. A high storage modulus is present at low temperatures and high frequencies. The modulus drops rapidly at higher temperatures and lower frequencies. This corresponds to the material’s thermal transition through Tg. Master curves of BPSH-BPS MBC ionomers were constructed using the TTS principle by horizontally shifting the storage modulus curves with respect to a reference temperature To ¼ 230 C (Fig. 7) [8]. Kohlrausch-Williams-Watts (KWW) equation (Eqn. (5)) is generally used to describe the relaxation behavior of polymers in the glass transition region [16,21]. In the KWW equation, the stretched exponential function f(t) is equivalent to the time dependent changes in the glassy Eg and equilibrium Ee moduli. These changes are related to the effective relaxation time s, the coupling parameter n, and the distribution of relaxation times (1 n) that is related to the polydispersity in molecular weight [19]. In this work, the KWW equation was used to describe the
Table 2 Copolymer physical properties. MPC Ionomer
IECv FFV Density Error IECw mmol/g mmol/cm3 g/mL %
BPSH-BPS (5k:5k) BPSH-BPS (10k:10k) BPSH-BPS (15k:15k) Nafion
1.41 1.43 1.48 1.99
1.5 0.7 0.6 e
1.30 1.38 1.40 0.91
1.83 1.97 2.07 1.81
EA (H2) V0 L/mol kJ/mol
0.121 7.3 0.115 13.6 0.096 21.6 e e
0
To C
C1
o
C2 K
1067 1067 1067
230 230 230
120 120 120
1085 1088 1087
viscoelastic properties of BPSH-BPS MBC ionomers. Table 4 summarizes the fitting parameters obtained from the KWW equation.
15k
6
0
C2 K
20.1 19.4 20.6 23.0
1n # t
s
¼
EðtÞ Ee Eg Ee
(5)
Clearly, block length significantly influences relaxation times s and polymer chain relaxation (1 n). As seen from the storage modulus master curves in Fig. 7a, the rapid decline in modulus at longer times for BPSH-BPS MBC ionomers is linked to a decrease in the characteristic relaxation time s. The characteristic relaxation time s is plotted versus the molecular weight of a repeat unit in Fig. 8. s showed a linearly decreasing trend with increasing of block molecular weight. The parameter (1 n) is correlated to the width of the relaxation spectrum in the glass transition region and hence to the polydispersity (PDI). A more polydisperse system should have a lower (1 n) value due to a greater degree of polymer chain motion [22]. For a monodisperse polystyrene, (1 n) was found to be in the range of 0.57e0.60 [23]. BPSH-BPS MBC ionomers were synthesized by a two-step polymerization method, which generally creates more polydisperse polymers. Reasonable values of (1 n) in a range of 0.257e0.287 were obtained for these materials. The coupling parameter n is used to describe polymer chain motion or reptation, which is a measure of the strength of intermolecular interactions [23]. Increases in the value of n were directly proportional to block length. This result implies that longer block lengths have a more complicated relaxation process due to stronger intermolecular interactions as observed for the 15k BPSH-BPS MBC ionomer. The master curve of tan d (Fig. 7b) for BPSH-BPS MBC ionomers clearly illustrates the shifting of peaks to higher frequencies and narrowing of peak breadth as the block length increases. The equilibrium modulus, Ee of the MBC ionomers converged with each other at large time scales. However, the glassy storage modulus Eg decreased with increasing block length. This reveals that shorter block lengths are more favorable to enhancing mechanical properties with simpler viscoelastic material properties. 3.3. Gas permeability 3.3.1. Fractional free volume Volumetric properties are extremely important for a multitude of phenomenon, processes, and there exists many interesting relationships between a polymers free volume and its van der Waals volume (Vw) [24e27]. The van der Waal volume is defined as the molecular space occupied by a molecule, which is impenetrable
Table 4 KWW fitting parameters. MPC Ionomer BPSH-BPS (5k:5k) BPSH-BPS (10k:10k) BPSH-BPS (15k:15k)
1n 0.287 0.280 0.257
n 0.713 0.720 0.743
s s
Eg Pa
Ee Pa
17.8 11.7 4.2
3.19E9 1.70E9 1.43E9
3.70E8 2.95E8 2.95E8
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T(K) 360
320
H2
300
10k
Ea 20.6 kJ/mol
2.0
5k
3.2
Ea 20.1 kJ/mol Ea 19.4 kJ/mol
P H2 (Barrers)
2.5
LnP(Barrers)
340
1.5
2.8
15k 2.4
BPSH-BPS(5k:5k) BPSH-BPS(10k:10k) BPSH-BPS(15k:15k)
1.0
2. 6
2 .8
8
9
10
11
1/FFV
3. 0
3 .2
3.4
Fig. 10. H2 permeability versus 1/FFV at 30 C as a function of block length.
1000/T(K) Fig. 9. Hydrogen permeability versus temperature for BPSH-BPS MBC ionomers.
to other molecules with normal energies. This molecular volume can be estimated based upon the structural groups of the polymer and their van der Waals volumes. Bondi calculated the contributions of about 60 structural groups for their Vw. and developed group contribution relationships for predicting glassy state polymer properties [24]. In this work, the FFV is calculated by the chemical group contribution method proposed by Bondi, which is summarized by equations (6) and (7). In these equations, V is the specific volume of the polymer based on its density, Vo is the volume occupied by polymer chains based on Vw, and K is the total number of structural groups of the polymer repeat unit [24,25,28,29]. Table 2 summarizes the results of the FFV calculations of the BPSH-BPS MBC ionomers.
FFV ¼
V Vo V
Vo ¼ 1:3
K X
(6)
ðVw Þk
(7)
k¼1
The results from this model predict that longer ionic block lengths will have lower free volume. Consequently, longer ionic block lengths should have lower gas permeability and enhanced gas barrier properties. 3.3.2. Gas permeability Hydrogen permeability was measured from 30 C to 90 C for BPSH-BPS MPC ionomers and Nafion NE1035 in the dry state. Gas transport at these conditions is complex due to the nonequilibrium and temperature-time dependent chain motions [30]. However, an Arrhenius relationship between increasing temperature and gas permeability is observed for these materials shown in Fig. 9. Table 5 H2 permeability coefficient at 4 atm of BPSH-BPS MBC ionomers and Nafion. BPSH-BPS BPSH-BPS Nafion Temperature oC BPSH-BPS (5k:5k) Barrers (10k:10k) Barrers (15k:15k) Barrers Barrers 30 50 70 90
3.2 5.4 8.3 13.0
3.0 4.9 7.2 11.1
2.5 4.3 6.4 9.8
6.0 10.3 17.3 e
Block length noticeably influenced H2 permeability. The H2 permeability decreased from 13 Barrers to 10 Barrers as the block length increased from 5k to 15k at 90 C (Table 5). Over the entire temperature range, the 15k block length had the lowest H2 permeability. Moreover, the H2 permeability of all BPSH-BPS MBC ionomers was lower than Nafion. H2 transport through a polymer is strongly dependent on its free volume [31]. H2 permeability as a function of FFV is shown in Fig. 10. A linear relationship between 1/FFV and log(PH2) is observed for these materials. Longer 15k blocks were calculated to have a higher ionic density and lower gas permeability as shown in Fig. 10. H2 permeability is plotted as a function of temperature and block length (Fig. 9). Its apparent activation energy Ea is 20.1 kJ/mol (5k), 19.4 kJ/mol (10k), 20.6 kJ/mol (15k), and 23.0 kJ/mol (Nafion). While Nafion Ea (H2) is 15% higher than BPSH-MBC ionomers, variations may not be statistically significant. Gas transport is dependent on polymer chain dilation and proportional to its dspacing. The BPSH-MBC ionomer d-spacing is between 21.2 nm and 36.4 nm. These are two-orders of magnitude larger than the kinetic diameter of H2 at 0.289 nm. Therefore, H2 may be too small to be strongly influenced by these chain motions and dimensions. Future work will use larger probe molecules to understand polymer chain motions as it relates to polymer structure and composition.
4. Conclusions BPSH-BPS MBC ionomers with fixed IECw have comparable proton conductivities as Nafion. SAXS reveals that inter-domain spacing increases with block length. This may imply that a periodic lamellar morphology is present. The viscoelastic relaxation processes of these complex materials are adequately modeled using WLF and KWW theory. Block length is inversely proportional to FFV, characteristic chain relaxation time s, and the breadth of the relaxation process (1 n). The polymer chain’s molecular interaction parameter (n) increases with block length. Increasing block length decreases FFV, H2 permeability, and mechanical properties; but it increases ion conductivity and gas impermeability. The Ea for H2 permeation does not vary significantly between BPSH-BPS MBC ionomers (19.4e20.6 kJ/mol). This is attributed to chain dilation motions for gas transport being larger than the size of permeating H2, which creates insensitivity to thermally activated gas transport based on size. However, an onset of differences in Ea (H2) between Nafion and BPSH-BPS MBC ionomers is observed. Increasing block length improves ion domain density, conductivity, and lowers gas
Y. Fan et al. / Polymer 52 (2011) 3963e3969
permeability. Conversely, favorable physical properties are inversely related to the block size. This dichotomy illustrates the complexities of ionomer research, and the need for tailoring ordered materials with desirable properties. Acknowledgments This work was partially supported by the DuPont Company. The authors of this paper would like to thank Professor Dillard at Virginia Tech, and Professor Shaw and the IMS at UCONN for providing access to their DMA instruments. Finally, we acknowledge partial support by the National Science Foundation MRI program DMR0923107 that allowed us to do SAXS measurements in this study. References [1] Neburchilov V, Martin J, Wang H, Zhang J. Journal of Power Sources 2007; 169(2):221. [2] Inabaa M, Kinumotoa T, Kiriakea M, Umebayashia R, Tasakaa A, Ogumib Z. Electrochimica Acta 2006;51(26):5746. [3] Kocha SS, Yang JD, Yi JS. AIChE Journal 2006;52(5):1916. [4] Wilson MS, Garzon FH, Sickafus KE, Gottesfeld S. Journal of the Electrochemical Society 1993;140(10):2872. [5] Sethuraman VA, Weidner JW, Haug AT, Protsailo LV. The Electrochemical Society 2008;155(2):B119. [6] Zhang L, Mukerjee S. Journal of the Electrochemical Society 2006;153(6): A1062. [7] Lai YH, Mittelsteadt CK, Gittleman CS, Dillard DA. ASME Journal of Fuel Cell Science and Technology 2009;6:021002. [8] Patankar KA, Dillard DA, Case SW, Ellis MW, Lai Y-H, Budinski MK, et al. Mechanics of Time-Dependent Materials 2008;12(3):221. [9] Lee H-S, Roy A, Lane O, Dunn S, McGrath JE. Polymer 2008;49(3):715.
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[10] Recio R, Lozano AE, Pradanos P, Macros A, Tejerina F, Hernandez A. Journal of Applied Polymer Science 2008;107(2):1039. [11] Wang F, Hickner M, Kim YS, Zawodzinski TA, McGrath JE. Journal of Membrane Science 2002;197(1e2):231. [12] Rudnik E, Dobkowski Z. Journal of Thermal Analysis and Calorimetry 1995; 45(5):1153. [13] Badami AS, Lane O, Lee H-S, Roy A, McGrath JE. Journal of Membrane Science 2009;333(1e2):1. [14] Elabd YA, Walker CW, Beyer FL. Journal of Membrane Science 2004;231(1e2): 181. [15] Bae B, Miyatake K, Watanabe M. Macromolecules 2010;43(6):2684. [16] Ferry J. Viscoelastic properties of polymers. 3rd ed. New York: Wiley; 1980. [17] Shaw MT, MacKnight WJ. Introduction to polymer viscoelasticity. 3rd ed. New York, NY: Wiley Interscience; 2005. [18] Williams ML, Landel RF, Ferry JD. Journal of the American Chemical Society 1955;77(14):3701. [19] John DF, Grandine LD, Edwin RF. Journal of Applied Physics 1953;24(7):911. [20] Schwarzl FR, Zahradnik F. Rheologica Acta 1980;19(2):137. [21] Wimberger-Friedl R, Bruin JG. Rheologica Acta 1991;30(5):419. [22] Chen HY, Stepanov EV, Chum SP, Hiltner A, Baer E. Macromolecules 2000; 33(23):8870. [23] Santangelo PG, Ngai KL, Roland CM. Polymer 1998;39(3):681. [24] Bondi A. The Journal of Physical Chemistry 1964;68(3):441. [25] Krevelen DWV, Hoftyzer PJ. Properties of polymers, their estimation and correlation with chemical structure. 2nd ed. Amsterdam; New York: Elsevier; 1976. [26] Krevelen DWV, Nijenhuis KT. Properties of polymers: their correlation with chemical structure; their numerical estimation and prediction from additive group contributions. 4th ed. Amsterdam; Boston: Elsevier; 2009. [27] Park JY, Paul DR. Journal of Membrane Science 1997;125(1):23. [28] Sugden S. Journal of the Chemical Society; 1927:1786. [29] Nagel C, Günther-Schade K, Fritsch D, Strunskus T, Faupel F. Macromolecules 2002;35(6):2071. [30] Cornelius CJ. Physical and gas permeation properties of a series of novel hybrid inorganic-organic composites based on a synthesized fluorinated polyimide. Virginia Polytechnic Institute and State University; 2000. [31] Cromyn J. Polymer permeability. London; New York, NY: Elsevier Applied Science Publishers; 1985.
Contents lists available at ScienceDirect
Polymer journal homepage: www.elsevier.com/locate/polymer
Viscoelastic and gas transport properties of a series of multiblock copolymer ionomers Yanfang Fan a, Mingqiang Zhang b, Robert B. Moore b, Hae-Seung Lee b, James E. McGrath b, Chris J. Cornelius a, * a b
Department of Chemical Engineering, University of Connecticut, Storrs, Connecticut 06269, USA Department of Macromolecular Science and Engineering, Virginia Polytechnic Institute and State University, Blacksburg, VA 24061, USA
a r t i c l e i n f o
a b s t r a c t
Article history: Received 2 April 2011 Received in revised form 17 June 2011 Accepted 25 June 2011 Available online 1 July 2011
A series of sulfonated poly(arylene ether sulfone) (BPSH-BPS) multiblock copolymer (MBC) ionomers were studied with respect to structure and physical properties. Ion exchange capacity (IECw) on a weight basis was constant for all MBC ionomers with controlled ionic and nonionic block lengths of 5k:5k, 10k:10k and 15k:15k. Viscoelastic properties were investigated as a function of block length, relaxation time, and temperature ranging from 200 C to 250 C. Relaxation time and glassy state modulus decreased with increasing ionic block length. Williams-Landel-Ferry (WLF) theory adequately modeled the viscoelastic changes in the BPSH-BPS MBC ionomers due to changes in time and temperature using a shift factor aT. Small Angle X-ray Scatter (SAXS) revealed an average inter-chain d-spacing that increased with increasing ionic block length (21.1, 31.4 and 36.4 nm). The larger ionic domain size is attributed to an increase in bulk ionic density within an ionic domain. Bondi’s group contribution method predicts a lower fractional free volume (FFV) with increasing ionic block length. The 15k BPSHBPS MBC ionomer had the lowest FFV of the series, which corresponds to the lowest hydrogen permeability of 2.46 Barrers at 30 C and 4 atm. Ó 2011 Elsevier Ltd. All rights reserved.
Keywords: Ionomer Viscoelasticity Transport
1. Introduction Today, fuel cells are considered to be a promising technology for alternative energy for stationary power and transportation [1,2]. Proton Exchange Membrane Fuel Cells (PEMFC) are being studied due to this diverse energy potential [3]. While PEMFC technology has the potential to be environmentally friendly, a critical property needing to be solved is long-term durability. Research studies have shown that several factors can reduce the lifetime of a PEMFC, including corrosion of the carbon-support, platinum-particle dissolution and sintering, and membrane thinning due to material relaxation [4]. The challenges of minimizing dynamic mechanical and material stresses leading to fuel cell failures in PEMFC systems is an ongoing area of investigation. These failures are due to inadequate membrane ionomer physical properties resulting in its low durability [3]. One mechanism associated with the physical degradation of membrane ionomers is operational cycling due to variations in temperature and humidity. These variations cause * Corresponding author. Tel.: þ1 860 486 3686; fax: þ1 860 486 2959. E-mail address: [email protected] (C.J. Cornelius). 0032-3861/$ e see front matter Ó 2011 Elsevier Ltd. All rights reserved. doi:10.1016/j.polymer.2011.06.050
changes in the hydration conditions of the ionomer, which contributes to the formation of micro cracks during volumetric expansion and contraction. Once a micro-crack forms, enhanced mechanical stresses accelerate crack propagation, which results in a significant decrease in the durability of a proton exchange membrane (PEM) fuel cell, and its failure as a membrane separator [5,6]. Generally, fuel cell durability is highly correlated to viscoelastic properties of PEMs, which are timeetemperature dependent physical properties. Many researchers have focused on the investigation of PEMs viscoelastic properties. Lai et al. proposed a linear viscoelastic stress model to account for time, temperature, and humidity effects in polyfluorosulfonic acid (PFSA) membranes [7]. The results showed that large variations in humidity create a faster hydration/dehydration rate resulting in high residual tensile stresses. Patankar et al. performed relaxation tests using DMA on Gore-Select 57 PEMs [8]. Thermal and hygral master curves were constructed using time, temperature, and humidity to create superposition plots based on the modulus. This model predicts the glassy plateau modulus (Eg) and stresses only for PSFA type membranes. Hydrocarbon based ionomers are one class of promising alternative membrane for fuel cells [9]. Research from the McGrath
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O
O HO3S
O S O
SO3H O A
F
F
O
O F
O
F
O S O
Bn
Fig. 1. Chemical structure of BPSH-BPS MBC ionomers.
group has focused upon the development of poly(arylene ether sulfone) MBC ionomers. The acronym for these materials is based on BP representing Biphenol, S for sulfone, H for the acidified form, and xx is the degree of disulfonation (BPSH-xx or BPS-xx) [9]. A series of BPSH-BPS (x:y) MBC ionomers with x and y representing the block length of the oligomer is shown in Fig. 1. When a volumetric based ion exchange capacity (IECv) is used to compare proton conductivities, these materials are comparable to Nafion. In this work, time dependent modulus responses of BPSH-BPS MBC ionomers is correlated with physical properties and molecular structure. Morphology, viscoelasticity, and gas transport are studied with respect to ion and non-ion containing block lengths of 5k:5k, 10k:10k, and 15k:15k. The effects of structure and block length with respect to changes in viscoelastic and hydrogen transport properties are discussed. 2. Experimental 2.1. Materials Synthesis of BPSH-BPS MBC ionomers has been documented in detail by Hae-Seung et al. [9]. BPSH-BPS was synthesized via a coupling reaction between phenoxide terminated BPSH100 oligomers and HFB end-capped BPS0 oligomers. As for the hydrophilic oligomer, BPSH100 was synthesized by varying the molecular weight via step growth polymerization. BPS0 of hydrophobic oligomer was first synthesized, and then end-capped with HFB via a nucleophilic aromatic substitution reaction. Nafion (1100 EW, sulfonic acid form, Hþ) was purchased from VWR LLC as a comparison with BPSH-BPS copolymers. Sulfonic acid (2 mol/L) was purchased and used as received from SigmaeAldrich.
form by boiling in 0.5 M H2SO4 for 2 h, rinsing in DI (deionized) water, and then boiling in DI water for 2 h to remove any excess acid. Wet samples were dried under vacuum at 60 C for 3 h.
2.3. Thermal analysis The glass transition temperature for Hþ form BPSH-BPS MBC ionomers was obtained using Dynamic Mechanical Analyzer Q800 (DMA, TA instruments). Membrane samples were analyzed in the tensile mode. Material tests were done using a frequency of 1 Hz with a heating rate of 2 C/min from 50 C to 250 C and a tensile strain 3t of 0.1%.
2.4. Short-term stress relaxation tests Frequency dependent relaxation tests for BPSH-BPS MBC ionomers were performed over a temperature range of 200e250 C using a Q800 DMA. A static force (stress) of 0.01 N was used in all experiments. This stress was experimentally determined to give a linear viscoelastic response. Materials were converted to the Hþ form prior to testing.
2.5. Gas transport properties Hydrogen permeability through dry films was measured using the salt form of a BPSH-BPS MBC ionomer. A gas permeation system was custom designed and built to determine gas transport properties using the time-lag method [10,30]. The gas purity used in this study was 99.99%. All experiments use a feed pressure of 4 atm and varied the temperature from 30 to 90 C.
2.2. Membrane preparation The salt form of BPSH-BPS MBC copolymers were dissolved in NMethylpyrrolidone (NMP) yielding transparent solutions (5e7 wt.%), which were solution cast on Teflon or glass substrates. Membrane films were created by controllably evaporating solvent from the polymer solution. Solvent removal was done in multiple steps involving drying for in a vacuum for 24 h at 60 C with an additional vacuum drying step at 110 C for 48 h. Membranes were solution cast into films having a nominal thickness ranging from 0.020 to 0.070 mm. Membrane films were converted to their Hþ
2.6. Density measurement A hydrostatic weighing method was used to measure the density of a BPSH-BPS MBC ionomer. Dry BPSH-BPS MBC ionomer films were weighed in the air a total of three times to yield an average sample mass w0. w is the films buoyant mass measured in hexane. The polymer density r can be obtained by the following relationship using hexane density rhexane and a small correction for the density of air rair (Eqn. (1)). The average error is 0.8%.
Fig. 2. Temperature dependence of BPSH-BPS MBC ionomers in Hþ form versus (a) tan d and (b) storage modulus E0 .
Y. Fan et al. / Polymer 52 (2011) 3963e3969
qMAX
5
40
15
d1
2qMAX 4
d 1 (nm)
Log I(q) [a.u.]
10
15k
2.5
IECV
32
3.0
2.0 24
2
10k
1.5
5k
16 5
0 0 .1
0.3
10
15
3
1.0
Molecular Weight × 10 (g/mol)
1
-1
IEC V (mol/L)
6
3965
q (nm )
Fig. 4. Dependence on block length versus d-spacing (square) and IECV (circle).
Fig. 3. SAXS spectra of BPSH-BPS MBC ionomers.
w0 þ rair *r w0 w hexane
(1)
2.7. Small angle x-ray scattering Small angle X-ray scattering (SAXS) using a Rigaku S-Max 300 3 pinhole SAXS system was used to evaluate structure. The radiation X-ray source is CuKa. The incident X-ray beam was attenuated to a wave length of 0.154 nm with a sample-to-detector distance of 16.0 cm. 3. Results and discussion 3.1. Physical properties A series of BPSH-BPS MBC ionomers were evaluated using DMA to determine if block length effects chain mobility and the glass transition temperature (Fig. 2). A single glass transition temperature (Tg) was observed for the MBC ionomers in the Hþ form that is attributed to the regularity of block length, fixed IEC, and similar phase separation properties. The sulfonated hydrophilic blocks have a Tg varying between 248 Cand 251 C, which is in the range of a random BPSH copolymer ionomer (230 Ce280 C) [11]. BPSHBPS MBC ionomers’ sulfonic acid groups were converted to the salt form to increase thermal stability and partially shield ionic charge interactions. When samples are converted to their salt forms, a lower Tg is observed around 180 C. This thermal transition is associated with the unsulfonated hydrophobic blocks. The observation of a lower Tg is due to a reduction in the concentration of ionic interactions among polymer chains that masks the Tg of the unsulfonated domain. This result is close to commercial polyphenylsulfone with a Tg of 190 C [12]. The primary difference between the two Tgs is attributed to differences in molecular weight Table 1 SAXS domain space comparison. MPC Ionomer
qmax nm1
2qmax nm1
d1 nm
d2 nm
BPSH-BPS (5k:5k) BPSH-BPS (10k:10k) BPSH-BPS (15k:15k)
0.297 0.200 0.173
0.587 0.406 0.356
21.2 31.4 36.4
21.4 31.0 35.3
T-T 0 ( o C) -40
-20
0
20
40
5k 10k 15k WLF equation
5
log(aT )
r¼
of the unsulfonated block length, and the degree and concentration of hydrogen bonding between ionic groups. Previous studies of these materials show that they have a tendency to form a phase-separated morphology. TEM and AFM phase images reported by Badami confirmed that the degree of phase separation is enhanced by an increase in block length [13]. Additional structural information of these materials was evaluated with through-plane measurements using SAXS. The scattering spectra of BPSH-BPS MBC ionomers are shown in Fig. 3 In this figure, I(q) is the scattered intensity reported in arbitrary units (a.u.) and q the scattering vector that is related to the scattering angle 2q [q ¼ (4p/l sin q), l is wave length of x-ray beam]. SAXS spectra of the BPSH-BPS MBC ionomers reveals a maximum scattering peak qmax at a lower scattering angle corresponding to a first order reflection. Assuming that qmax is due to interparticle interferences between ionic cluster domains, this peak can be used to estimate the inter-ionic distance d1. This distance is calculated by relating the maximum scattering vector qmax to Bragg’s law (d1 ¼ 2p/qmax) [14]. The nominal d-spacing between chains was 21.2 nm for 5k, 31.4 nm for 10k, and 36.4 nm for 15k (Table 1). This result suggests that ionic domain size is a function of
5
0
0
-5
-5
200
220
240
260
o
T ( C) Fig. 5. Shift factor aT versus temperature.
280
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Y. Fan et al. / Polymer 52 (2011) 3963e3969
a
b 10
10
o
9
10
10
o
E' (Pa)
E' (Pa)
220 C o 225 C o 230 C o 235 C o 240 C o 245 C o 250 C o 255 C o 260 C
10
9
10
8
220 C o 225 C o 230 C o 235 C o 240 C o 245 C o 250 C o 255 C o 260 C
10
1
f (HZ)
5
10
2
10
-1
10
-4
f× aT (HZ)
Fig. 6. a) Frequency versus Temperature of a 5k:5k BPSH-BPS MBC ionomer and b) reconstructed modulus using the shift factors aT.
block length (Fig. 4). A weak secondary order peak at 2qmax is observed for all three MBC ionomers. This scattering behavior suggests the presence of a well-ordered and periodic microstructure that may have a lamellar morphology since d2 is similar to d1. Improvements in ionomer phase separation may lead to better chemical stability and mechanical properties because of hydrophobic domains, while hydrophilic domains may enhance proton transport. These morphology considerations will be discussed with respect to viscoelastic and gas transport properties. The densities of BPSH-BPS MBC ionomers were 1.41, 1.43 and 1.48 g/cm3 that increased with block length. A special feature of these MBC ionomers is a fixed IECw based on mass. Using the polymer density, a volumetric IEC (IECv, mmol/cm3) was defined as the molar concentration of sulfonic acid groups per unit volume of dry membrane [15]. The IECv versus block length in Fig. 4 shows a direct relationship between increasing IECv and block length. This increase in IECv may be indicative of an increase in ionic density due to block length and domain size. Domain size expands with increasing local sulfonic acid group concentration. This results in a greater concentration of expanded chains observed by an increasing d-spacing measured by SAXS. Overall, the materials morphology and structure are tied to block length, which is one crucial factor influencing phase separation, ionic domain size, and ionic density. 3.2. Timeetemperature superposition (TTS) Polymer viscoelastic properties are dependent on time and temperature [16]. Due to this interdependence, changes in properties can be correlated using a timeetemperature superposition (TTS) [17]. This interrelationship of time and temperature has been
a 10
used to successfully characterize the physical properties of many amorphous polymers [18,19]. In this work, TTS is utilized to characterize time and temperature dependent ionomer properties. Mathematically, changes in the modulus (E) due to different temperatures (T1,T2) and time (t1,t2) can be equated to each other using the following relationship E(t1,T1) ¼ E(t2/aT,T2). A time dependent shift factor aT, compensates for viscoelastic changes in the time scale due to temperature. aT is proportional to the polymer chain’s temperature dependence, segmental mobility, frictional diffusion, and reptation [18]. The shift factor aT is used to construct master curves at a reference temperature To in order to reconstruct viscoelastic responses from time t1 to t2. Shift factors are determined based on the Williams, Landel and Ferry (WLF) relationship (Eqn. (2)) [18]. In the WLF relationship, empirical constants C1 and C2 are calculated based upon the reference temperature To and aT. 0 The coefficients C1 and C2 can be appropriately transformed to C1 0 and C2 with Tg being used as the reference temperature (Eqns. (3) and (4)).
Log aT ¼
(3)
C20 ¼ C2 þ Tg To
(4)
The WLF relationship was utilized to study time and temperature dependent physical properties of BPSH-BPS MBC ionomers. A common reference temperature To was chosen for investigating the
0
tan(Delta)
E'(Pa)
BPSH-BPS(5K:5K) BPSH-BPS(10K:10K) BPSH-BPS(15K:15K)
10
9
BPSH-BPS (5K:5K) BPSH-BPS(10K:10K) BPSH-BPS(15K:15K) 10
8
10
6
10
3
10
f×aT (HZ)
0
10
-3
(2)
C1 C2 C2 þ Tg To
C10 ¼
b 10
10
C 0 ðT To Þ C1 T Tg ¼ 0 1 C2 þ ðT To Þ C2 þ T Tg
10
-1
10
-2
10
-6
10
-3
10
0
f×aT (HZ)
10
3
10
6
Fig. 7. Master curves of BPSH-BPS MBC ionomers of different block length using a KWW fit for (a) Storage Modulus and (b) tan d.
Y. Fan et al. / Polymer 52 (2011) 3963e3969 Table 3 WLF constants and Tg.
5k
18
3967
MPC Ionomer
Tg C
C1
248 251 250
122 122 122
τ (s)
o
BPSH-BPS (5k:5k) BPSH-BPS (10k:10k) BPSH-BPS (15k:15k)
10k
12
"
5
10
fðtÞ ¼ exp
15
Molecular Weight × 10 3 (g/mol) Fig. 8. Characteristic time s versus block molecular weight.
dependence of block length and temperature based on aT. The viscoelastic time and temperature dependent properties for all BPSH-BPS MBC ionomers were fitted with the WLF equation using aT and the empirical constants C1 and C2 listed in Table 3 [19]. A plot of aT versus T and (T To) of different block length materials resulted in nearly identical values and slopes (Fig. 5). This result suggests that block length may not be a significant factor affecting the relationship between aT and temperature. Many researchers commonly use the Tg to compare shift factors between different material systems [16]. In this case, the Tg measured by DMTA was used as the reference temperature. The 0 0 corresponding values of C1 and C2 are listed in Table 3. It has been 0 0 suggested that C1 and C2 are material specific. Because BPSH-BPS MBC ionomers have the same structure, it is not surprising that 0 0 C1 and C2 did not vary significantly with polymer block length [20]. 0 0 Generally, C1 is often related to FFV at Tg and C2 is related to the differences between Tg and the thermodynamically defined transition temperature where free volume is zero [16]. The frequency versus temperature of BPSH-BPS MBC ionomers is shown in Fig. 6. A high storage modulus is present at low temperatures and high frequencies. The modulus drops rapidly at higher temperatures and lower frequencies. This corresponds to the material’s thermal transition through Tg. Master curves of BPSH-BPS MBC ionomers were constructed using the TTS principle by horizontally shifting the storage modulus curves with respect to a reference temperature To ¼ 230 C (Fig. 7) [8]. Kohlrausch-Williams-Watts (KWW) equation (Eqn. (5)) is generally used to describe the relaxation behavior of polymers in the glass transition region [16,21]. In the KWW equation, the stretched exponential function f(t) is equivalent to the time dependent changes in the glassy Eg and equilibrium Ee moduli. These changes are related to the effective relaxation time s, the coupling parameter n, and the distribution of relaxation times (1 n) that is related to the polydispersity in molecular weight [19]. In this work, the KWW equation was used to describe the
Table 2 Copolymer physical properties. MPC Ionomer
IECv FFV Density Error IECw mmol/g mmol/cm3 g/mL %
BPSH-BPS (5k:5k) BPSH-BPS (10k:10k) BPSH-BPS (15k:15k) Nafion
1.41 1.43 1.48 1.99
1.5 0.7 0.6 e
1.30 1.38 1.40 0.91
1.83 1.97 2.07 1.81
EA (H2) V0 L/mol kJ/mol
0.121 7.3 0.115 13.6 0.096 21.6 e e
0
To C
C1
o
C2 K
1067 1067 1067
230 230 230
120 120 120
1085 1088 1087
viscoelastic properties of BPSH-BPS MBC ionomers. Table 4 summarizes the fitting parameters obtained from the KWW equation.
15k
6
0
C2 K
20.1 19.4 20.6 23.0
1n # t
s
¼
EðtÞ Ee Eg Ee
(5)
Clearly, block length significantly influences relaxation times s and polymer chain relaxation (1 n). As seen from the storage modulus master curves in Fig. 7a, the rapid decline in modulus at longer times for BPSH-BPS MBC ionomers is linked to a decrease in the characteristic relaxation time s. The characteristic relaxation time s is plotted versus the molecular weight of a repeat unit in Fig. 8. s showed a linearly decreasing trend with increasing of block molecular weight. The parameter (1 n) is correlated to the width of the relaxation spectrum in the glass transition region and hence to the polydispersity (PDI). A more polydisperse system should have a lower (1 n) value due to a greater degree of polymer chain motion [22]. For a monodisperse polystyrene, (1 n) was found to be in the range of 0.57e0.60 [23]. BPSH-BPS MBC ionomers were synthesized by a two-step polymerization method, which generally creates more polydisperse polymers. Reasonable values of (1 n) in a range of 0.257e0.287 were obtained for these materials. The coupling parameter n is used to describe polymer chain motion or reptation, which is a measure of the strength of intermolecular interactions [23]. Increases in the value of n were directly proportional to block length. This result implies that longer block lengths have a more complicated relaxation process due to stronger intermolecular interactions as observed for the 15k BPSH-BPS MBC ionomer. The master curve of tan d (Fig. 7b) for BPSH-BPS MBC ionomers clearly illustrates the shifting of peaks to higher frequencies and narrowing of peak breadth as the block length increases. The equilibrium modulus, Ee of the MBC ionomers converged with each other at large time scales. However, the glassy storage modulus Eg decreased with increasing block length. This reveals that shorter block lengths are more favorable to enhancing mechanical properties with simpler viscoelastic material properties. 3.3. Gas permeability 3.3.1. Fractional free volume Volumetric properties are extremely important for a multitude of phenomenon, processes, and there exists many interesting relationships between a polymers free volume and its van der Waals volume (Vw) [24e27]. The van der Waal volume is defined as the molecular space occupied by a molecule, which is impenetrable
Table 4 KWW fitting parameters. MPC Ionomer BPSH-BPS (5k:5k) BPSH-BPS (10k:10k) BPSH-BPS (15k:15k)
1n 0.287 0.280 0.257
n 0.713 0.720 0.743
s s
Eg Pa
Ee Pa
17.8 11.7 4.2
3.19E9 1.70E9 1.43E9
3.70E8 2.95E8 2.95E8
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Y. Fan et al. / Polymer 52 (2011) 3963e3969
T(K) 360
320
H2
300
10k
Ea 20.6 kJ/mol
2.0
5k
3.2
Ea 20.1 kJ/mol Ea 19.4 kJ/mol
P H2 (Barrers)
2.5
LnP(Barrers)
340
1.5
2.8
15k 2.4
BPSH-BPS(5k:5k) BPSH-BPS(10k:10k) BPSH-BPS(15k:15k)
1.0
2. 6
2 .8
8
9
10
11
1/FFV
3. 0
3 .2
3.4
Fig. 10. H2 permeability versus 1/FFV at 30 C as a function of block length.
1000/T(K) Fig. 9. Hydrogen permeability versus temperature for BPSH-BPS MBC ionomers.
to other molecules with normal energies. This molecular volume can be estimated based upon the structural groups of the polymer and their van der Waals volumes. Bondi calculated the contributions of about 60 structural groups for their Vw. and developed group contribution relationships for predicting glassy state polymer properties [24]. In this work, the FFV is calculated by the chemical group contribution method proposed by Bondi, which is summarized by equations (6) and (7). In these equations, V is the specific volume of the polymer based on its density, Vo is the volume occupied by polymer chains based on Vw, and K is the total number of structural groups of the polymer repeat unit [24,25,28,29]. Table 2 summarizes the results of the FFV calculations of the BPSH-BPS MBC ionomers.
FFV ¼
V Vo V
Vo ¼ 1:3
K X
(6)
ðVw Þk
(7)
k¼1
The results from this model predict that longer ionic block lengths will have lower free volume. Consequently, longer ionic block lengths should have lower gas permeability and enhanced gas barrier properties. 3.3.2. Gas permeability Hydrogen permeability was measured from 30 C to 90 C for BPSH-BPS MPC ionomers and Nafion NE1035 in the dry state. Gas transport at these conditions is complex due to the nonequilibrium and temperature-time dependent chain motions [30]. However, an Arrhenius relationship between increasing temperature and gas permeability is observed for these materials shown in Fig. 9. Table 5 H2 permeability coefficient at 4 atm of BPSH-BPS MBC ionomers and Nafion. BPSH-BPS BPSH-BPS Nafion Temperature oC BPSH-BPS (5k:5k) Barrers (10k:10k) Barrers (15k:15k) Barrers Barrers 30 50 70 90
3.2 5.4 8.3 13.0
3.0 4.9 7.2 11.1
2.5 4.3 6.4 9.8
6.0 10.3 17.3 e
Block length noticeably influenced H2 permeability. The H2 permeability decreased from 13 Barrers to 10 Barrers as the block length increased from 5k to 15k at 90 C (Table 5). Over the entire temperature range, the 15k block length had the lowest H2 permeability. Moreover, the H2 permeability of all BPSH-BPS MBC ionomers was lower than Nafion. H2 transport through a polymer is strongly dependent on its free volume [31]. H2 permeability as a function of FFV is shown in Fig. 10. A linear relationship between 1/FFV and log(PH2) is observed for these materials. Longer 15k blocks were calculated to have a higher ionic density and lower gas permeability as shown in Fig. 10. H2 permeability is plotted as a function of temperature and block length (Fig. 9). Its apparent activation energy Ea is 20.1 kJ/mol (5k), 19.4 kJ/mol (10k), 20.6 kJ/mol (15k), and 23.0 kJ/mol (Nafion). While Nafion Ea (H2) is 15% higher than BPSH-MBC ionomers, variations may not be statistically significant. Gas transport is dependent on polymer chain dilation and proportional to its dspacing. The BPSH-MBC ionomer d-spacing is between 21.2 nm and 36.4 nm. These are two-orders of magnitude larger than the kinetic diameter of H2 at 0.289 nm. Therefore, H2 may be too small to be strongly influenced by these chain motions and dimensions. Future work will use larger probe molecules to understand polymer chain motions as it relates to polymer structure and composition.
4. Conclusions BPSH-BPS MBC ionomers with fixed IECw have comparable proton conductivities as Nafion. SAXS reveals that inter-domain spacing increases with block length. This may imply that a periodic lamellar morphology is present. The viscoelastic relaxation processes of these complex materials are adequately modeled using WLF and KWW theory. Block length is inversely proportional to FFV, characteristic chain relaxation time s, and the breadth of the relaxation process (1 n). The polymer chain’s molecular interaction parameter (n) increases with block length. Increasing block length decreases FFV, H2 permeability, and mechanical properties; but it increases ion conductivity and gas impermeability. The Ea for H2 permeation does not vary significantly between BPSH-BPS MBC ionomers (19.4e20.6 kJ/mol). This is attributed to chain dilation motions for gas transport being larger than the size of permeating H2, which creates insensitivity to thermally activated gas transport based on size. However, an onset of differences in Ea (H2) between Nafion and BPSH-BPS MBC ionomers is observed. Increasing block length improves ion domain density, conductivity, and lowers gas
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permeability. Conversely, favorable physical properties are inversely related to the block size. This dichotomy illustrates the complexities of ionomer research, and the need for tailoring ordered materials with desirable properties. Acknowledgments This work was partially supported by the DuPont Company. The authors of this paper would like to thank Professor Dillard at Virginia Tech, and Professor Shaw and the IMS at UCONN for providing access to their DMA instruments. Finally, we acknowledge partial support by the National Science Foundation MRI program DMR0923107 that allowed us to do SAXS measurements in this study. References [1] Neburchilov V, Martin J, Wang H, Zhang J. Journal of Power Sources 2007; 169(2):221. [2] Inabaa M, Kinumotoa T, Kiriakea M, Umebayashia R, Tasakaa A, Ogumib Z. Electrochimica Acta 2006;51(26):5746. [3] Kocha SS, Yang JD, Yi JS. AIChE Journal 2006;52(5):1916. [4] Wilson MS, Garzon FH, Sickafus KE, Gottesfeld S. Journal of the Electrochemical Society 1993;140(10):2872. [5] Sethuraman VA, Weidner JW, Haug AT, Protsailo LV. The Electrochemical Society 2008;155(2):B119. [6] Zhang L, Mukerjee S. Journal of the Electrochemical Society 2006;153(6): A1062. [7] Lai YH, Mittelsteadt CK, Gittleman CS, Dillard DA. ASME Journal of Fuel Cell Science and Technology 2009;6:021002. [8] Patankar KA, Dillard DA, Case SW, Ellis MW, Lai Y-H, Budinski MK, et al. Mechanics of Time-Dependent Materials 2008;12(3):221. [9] Lee H-S, Roy A, Lane O, Dunn S, McGrath JE. Polymer 2008;49(3):715.
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