Polymers in Membrane Electrode Assemblies

10.36

Polymers in Membrane Electrode Assemblies

DS Kim, C Welch, RP Hjelm, and YS Kim, Los Alamos National Laboratory, Los Alamos, NM, USA MD Guiver, National Research Council, Ottawa, ON, Canada; Hanyang University, Seoul, Republic of Korea © 2012 Elsevier B.V. All rights reserved.

10.36.1 10.36.2 10.36.2.1 10.36.2.1.1 10.36.2.1.2 10.36.2.1.3 10.36.2.2 10.36.2.2.1 10.36.2.2.2 10.36.2.3 10.36.2.4 10.36.3 10.36.3.1 10.36.3.2 10.36.3.2.1 10.36.3.2.2 10.36.3.2.3 10.36.3.2.4 10.36.3.3 10.36.3.3.1 10.36.3.3.2 10.36.3.3.3 10.36.4 References

Introduction Polymer Electrolyte Membranes Water Uptake Chemistry effect Polymer architecture effect Effect of polymer processing Proton Conductivity Under fully hydrated conditions Under partial humidified conditions Methanol Permeability Summary Polymer Electrolyte Ionomers in the Electrode Historical Background Structural Effect Ionomer composition Ion exchange capacity Hydrophobicity Dispersing solvent Membrane–Electrode Interface Interfacial resistance between PEM and electrode Origin of interfacial failure Methods to improve interfacial compatibility Summary

10.36.1 Introduction Modern polymer electrolyte membrane fuel cells (PEMFCs) to date have incorporated a membrane electrode assembly (MEA) that consists of (1) a polymer electrolyte membrane (PEM), (2) anode and cathode catalyst layers, and (3) the gas diffusion layers (GDLs). Figure 1 depicts the five-layered MEA structure and the component dimensions typically used for PEMFCs. The PEM is located in the center of the MEA, is sandwiched between two thin anode and cathode layers, and allows ions to pass from the anode to the cathode while the electrons induce a current through an external circuit to the cathode. The electrodes contact with the GDLs, which provide a pathway for fuel and oxidant and an electrically conductive pathway for current collection. The GDL is typically made from carbon paper or cloth having a hydropho­ bic microporous coating layer, which provides effective wicking of liquid water from the electrode to the diffusion media. Typically, polytetrafluoroethylene (PTFE) is used as a hydropho­ bic component to enhance water removal capability. The role of the PEM is to provide ionic conductivity, to prevent the flow of electrons, to act as a barrier to the reactants, and to maintain chemical and mechanical stabilities. Anode and cathode catalyst layers, where the fuel oxidation and oxygen reduction reactions (ORRs) take place, respectively, consist of electrocatalyst, catalyst-supporting material, and polymer electrolyte ionomer. The ionomer contacts with electrocatalyst and catalyst-supporting materials to help the electrochemical reaction by providing ion and reactant pathway and sustaining mechanical integrity. Polymer Science: A Comprehensive Reference, Volume 10

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Typical fuel cell MEAs are prepared either by catalyst-coated membrane (CCM) or by gas diffusion electrode (GDE) meth­ ods. The CCM method includes the processes to apply the catalyst layer to the PEM by direct coating or decal transfer and lamination of GDLs to the CCM. The GDE method involves a process of application of catalyst to the GDLs and lamination of the GDEs to the PEM. A variety of MEA fabrica­ tion techniques such as modified thin-film electrodes,1 vacuum deposition,2,3 electrodeposition,4,5 dual ion-beam-assisted deposition,6 electroless deposition,7 catalyst impregnation,8 and modified GDE9,10 have been developed based on CCM and GDE configurations. Polymeric materials are utilized as PEMs and electrode ionomers, which are important elements in the MEA structure and greatly impact fuel cell performance. In this chapter, PEMs and electrode ionomers utilized in MEA structures are described, with a focus on the utilization of these materials in practical PEMFC devices.

10.36.2 Polymer Electrolyte Membranes PEMs play two basic roles in MEA: (1) ion conductor and (2) separator. As an ion conductor, PEMs require high ionic conductivity. Generally, the ionic resistance of the cell is much greater than the electronic resistance; cell performance can be greatly improved by increasing ion conductivity. As a separator, PEMs must not conduct electrons, and must transport fuel and

doi:10.1016/B978-0-444-53349-4.00287-9

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Polymers in Energy Applications | Polymers in Membrane Electrode Assemblies

Anode

PEM

Cathode

4. failure of the MEA active area edge, where the wet and dry regions meet and mechanical stress is concentrated.

GDL Carbon (20–50nm)

Water uptake of PEMs is usually reported by measuring the weight difference between dry and wet PEM samples: WU ¼

Ionomer Electrocatalyst (2–10nm) 50–400μm

1–20μm 20–200μm

Figure 1 Schematic diagram of membrane electrode assembly.

oxidant in a minimal quantity. Electron conduction results in short-circuiting; fuel and oxidant permeation through PEMs increase electrode overpotentials. The separator requires mechanically robust film-forming properties to provide mechanical integrity of MEA. Although PEMs require a number of combined properties, the relative importance of each prop­ erty is not equal. For example, the electron-insulating property is not as important as ionic conductivity, since most PEMs are good electron insulators by nature. Common key properties of PEMs include ion conductivity, reactant permeability, and mechanical properties. Selection of a PEM for specific fuel cell applications is achieved from optimizing those key properties by changing polymer chemistry, architecture, and processing conditions. In this section, general structure–property relation­ ships of sulfonated PEMs are explained in order to discuss how those key properties are optimized and why specific families of PEMs are preferred in certain fuel cell applications. Water uptake of PEMs is first discussed in terms of polymer chemistry and architecture, and then proton conductivity and methanol permeability are discussed in the following part.

10.36.2.1

Water Uptake

Sulfonated PEMs absorb water because sulfonic acid groups have affinity to water. Water that is absorbed in PEMs is directly linked to proton conductivity, mechanical properties, and reac­ tant permeability. Additional water beyond an optimal amount impacts MEA processibility and fuel cell durability. Several problems related to highly swollen PEMs have been observed in MEA processing and fuel cell operations: 1. membrane wrinkling and/or membrane dissolution during CCM MEA fabrication procedure due to the nonuniform contact of catalyst ink solution on membranes; 2. interfacial delamination between membrane and electrode (will be discussed in Section 10.36.3.3); 3. increased creep and pin hole formation during fuel cell operations; and

WWET − WDRY  100 WDRY

½1

where WWET and WDRY are the polymer weights in wet and dry states, respectively. In general, PEMs with high water uptake ( > 100 wt.%) have limited stability during MEA fabrication or fuel cell test­ ing. Water uptake of PEMs can be controlled by sulfonic acid concentration, chemical structure, and polymer architecture of PEMs. Water uptake is proportional to the sulfonic acid con­ centrations, which is well understood. The effect of polymer chemical structure and architecture of a wide range of PEMs, on the other hand, is relatively less understood. In this section, PEM chemistry and architecture effects on water uptake are discussed in order to better understand the structure–property relationship. Water uptake of various PEMs are compared at a fixed sulfonic acid concentration that is expressed by weight-based ion exchange capacity, IECW (meq. g− 1), which is the inverse value of equivalent weight (EW, the number of grams of dry membrane per mole of sulfonic acid groups).

10.36.2.1.1

Chemistry effect

Polymer structure has a profound impact on the water uptake of PEMs. In order to analyze this effect in a systematic way, PEMs are arbitrarily categorized based on their molecular com­ position into four classes: (1) perfluorinated sulfonic acid (PFSA), (2) fluorinated, (3) hydrocarbon (HC)-based, and (4) functional PEMs. PFSAs denote a polymer family that has perfluorovinyl ether groups terminated with sulfonated groups onto a PTFE backbone. Nafion is an examples of a PFSA PEMs. Fluorinated PEMs include polymer blends with PFSA and par­ tially fluorinated HC. HC-based PEMs include nonfluorinated polymers having no specific functional groups that interact with sulfonic acid or water molecules. Sulfonated polyaro­ matics derived from polyethers, polysulfones, polyketones, or polystyrene and sulfonated heterocycles such as polyimides are classified into this category. Functional PEMs include polymers or polymer blends containing basic functionality such as imi­ dazole or amine, and polar groups such as nitrile or phosphine oxide. The functional groups in the PEMs provide additional interactions through acid–base, dipolar, or hydrogen-bonding interactions besides the interactions between sulfonic acid group and water. Figure 2 compares water uptake for fluori­ nated PEMs as a function of the degree fluorination at an IECW of 1.4 meq.g−1. Wholly aromatic nonfluorinated poly(arylene ether ketone) showed the lowest water uptake (17 wt.%), while highly fluorinated decafluorobiphenyl and hexafluoro (6F)-based poly(arylene ether) showed the highest water uptake among those compared. Poly(arylene ether)s with inter­ mediate levels of fluorine content showed higher water uptake than wholly aromatic polymers, but lower water uptake than highly fluorinated ones. Water uptake of HC-based PEMs is further influenced by the backbone stiffness and specific inter­ actions. It was observed that highly rigid polyimide has a noticeably lower water uptake than relatively flexible polysul­ fone at a similar sulfonic acid concentration and the water

Polymers in Energy Applications | Polymers in Membrane Electrode Assemblies

693

Figure 2 Effect of fluorine content on water uptake for poly(arylene ether) polymers: (a) wholly aromatic poly(arylene ether ketone),11, 12 (b) 6F-based poly(arylene ether ketone),13 (c) decafluorobiphenyl-based poly(arylene ether),14 (d) 6F and decafluorobiphenyl-based poly(arylene ether) random copolymers.15

uptake of poly(arylene ether) with phenyl phosphine oxide (PPO) group decreased with increasing amounts of PPO func­ tional groups. In this case, PPO has a dipole–dipole interaction with sulfonic acid groups, essentially replacing the interaction with water molecules. A reduction of water uptake has also been observed in PEMs having ionic, acid–base, and hydrogen-bonding interactions. Increasing water uptake with fluorine content is also observed in radiation-grafted poly(styrene sulfonic) acids (PSSAs). Another example of the effect of fluorine content on water uptake is depicted in Figure 3 which shows the water uptake of non-cross-linked radiation-grafted PSSAs as a func­ tion of polymer backbone fluorination. As the polymer backbone transitions from polyethylene, poly(vinylidene fluoride) (PVDF), PVDF-co-hexafluoropropylene (HFP) and finally to fluorinated ethylene propylene (FEP), water uptake increased from 57 to 98 wt.%, even though the IECW slightly decreased from 1.94 to 1.80 meq. g−1. Both Figures 2 and 3

clearly show that the water uptake increases with fluorination at a given IECW, which is somewhat counterintuitive, since adding fluorine typically makes materials more hydrophobic. There are two major reasons that help to explain this beha­ vior. The first reason is related to the density difference between fluorine and hydrogen atoms.20 The atomic mass of fluorine is 19 times greater than that of hydrogen, while the atomic radius of fluorine is only 25% greater than that of hydrogen. This means that when hydrogen atoms are replaced with fluorine atoms in the polymer, the actual volume concentration of sul­ fonic acid groups is much higher than that estimated from weight-based sulfonic acid concentration that is expressed in IECW. The second reason is related to the water domain struc­ ture.21 Fluorinated PEMs have larger water domains than nonfluorinated PEMs. The water domain structure of sulfonated PEMs is determined by local sulfonic acid concentration and the local environment of sulfonic acid groups. Sulfonated PEMs have phase-separated structure primarily due to the greater

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Polymers in Energy Applications | Polymers in Membrane Electrode Assemblies

Figure 3 Effect of fluorine content on water uptake for radiation-grafted polymers containing fluorine groups.16–19

extent of difference in chemical structure between hydrophobic (nonsulfonated) and hydrophilic (sulfonated) segments. While the hydrophilicity of sulfonic acid group containing segments (i.e., solubility parameter) are similar for both fluorinated and nonfluorinated systems, the hydrophobicity of the fluorinated PEMs segments is much greater than the nonfluorinated analogs. This greater contrast between hydrophilic and hydrophobic nat­ ures of PEMs contributes the phase contrast of each domain structure. Consequently, fluorinated PEMs have greater local concentrations of sulfonic acid groups, which enhances water absorption. Fluorinated PEMs also have higher acidity due to the electron-withdrawing nature of fluorine, while the aryl sulfonic acid groups in nonfluorinated systems is somewhat weaker.

10.36.2.1.2

Polymer architecture effect

Polymer architecture also has a significant impact on water uptake of PEMs. Figure 4 compares water uptake of wholly aromatic poly (arylene ether)s having a different polymer architecture, but at similar IECW values of 1.8 meq. g−1. Cross-linked poly(arylene ether) showed the lowest water uptake of 17 wt.%. The reduced water uptake of the cross-linked PEMs is attributed to the volume restriction caused by chemical bonding. A similar volume restriction effect is also found in pore-filling composites,23 and PTFE-reinforced composites.24 Poly(arylene ether) with homopolymer-like architecture exhibited the second lowest

water uptake of 28 wt.%. The water uptake increased with ran­ dom copolymer (WU = 75 wt.%) and multiblock copolymers (WU = 100–140 wt.%). Within the class of multiblock copoly­ mers, water uptake increased with block length. The difference in water uptake between homopolymer-like structure, random, and multiblock copolymers is due to the intrinsic distribution of sulfonic acid and structural difference. The homopolymer-like structure provides a more uniform distribution of sulfonic acid groups, which reduces the local sulfonic acid concentration, while multiblock copolymers have self-assembled morphology where the sulfonic acid groups aggregate in the hydrophilic domains. Grafting polymer architecture, on the other hand, appears to have much less impact on water uptake. Figure 4 shows the effect of side chain length of poly(arylene ether) copolymers at an IECW of 1.2 meq. g−1. While a slight increase of water uptake was observed in both cases, it is arguable that water uptake increase for the longer graft chain is due to a higher local concentration of sulfonic acid group. Irrelevancy of graft chain architecture to water uptake make significant difference in water uptake between multiblock and grafted copolymers. Although graft chains may not have such a direct impact on water uptake, they could change the crystalline structure of PEMs and indirectly impact the water uptake. Figure 5 com­ pares water uptake of short side chain (SSC) and long side chain (LSC) PFSA at an IECW of 1.27 meq. g−1. It is noted that

Polymers in Energy Applications | Polymers in Membrane Electrode Assemblies

695

Figure 4 Effect of side chain length on water uptake for poly(arylene ether)s.22

Figure 5 Effect of polymer architecture (short side chain vs. long side chain) on water uptake of (a) short side chain PFSA, (b) long side chain PFSA, and (c) PTFE-g-PSSA.26–28

SSC PFSA has significantly lower water uptake (WU = 20– 33 wt.%) than LSC PFSA (WU = 47 wt.%). LSC PFSA is a totally amorphous structure and the disruption of crystalline structure

is due to sulfonic acid side chain, while SSC PFSA maintains a crystalline structure, which acts as barriers to solvent swelling.25,29

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10.36.2.1.3

10.36.2.2

δ

Effect of polymer processing

Water uptake of polymer electrolytes is also significantly affected by membrane processing. A number of studies have reported obtaining different properties for PFSAs, depending on their processing conditions. Preboiled Nafion PEMs have significantly increased water uptake compared to membranes dried at ele­ vated temperature and rehydrated at room temperature.30–33 Water uptake of Nafion is also affected by heat treatment. Nafion without heat treatment is referred to as the ‘E-form’ (expanded form), and membranes which were heat treated at 80 and 105 °C are referred to as the ‘N-form’ (normal form) and ‘S-form’(shrunken form) (normal form) and “S-form” (shrun­ ken form), respectively. After Nafion was dried at elevated temperature, it was observed to have decreased water uptake compared to that dried at lower temperature.34,35 It was also reported that fully hydrated HC-based PEMs treated by water at higher temperatures showed greater water absorption than PEMs treated at a lower temperature (30 °C).36,37 This is because at a certain temperature, termed the hydrogel temperature or Thg, the phase-separated domain structure is significantly disrupted and becomes a hydrogel, where water uptake increases considerably.

Proton Conductivity

Proton conductivity is one of the key properties in fuel cells and has been investigated extensively. There are two general mechan­ isms responsible for ion conduction: the vehicular mechanism, which relies on the physical transport of a vehicle to move ions, and the Grotthuss mechanism, which involves the ion being handed-off from one hydrogen-bonding site to another. The Grotthuss mechanism depends on the rate of polymer reorgani­ zation and solvation, while the vehicular mechanism depends on the rate of physical diffusion of the vehicle. Water plays a significant role in both mechanisms as a solvating agent for polymer electrolytes and as an ion carrier. Ion conduction via the Grotthuss mechanism is known to be much more efficient than proton conduction via the vehicular mechanism, since ion hopping is 1 order of magnitude faster than water diffusion. In polymer electrolytes, proton mobility via the Grotthuss mechan­ ism is very fast under fully hydrated conditions while significant suppression of intermolecular proton transfer occurs under less-humidified conditions and thus proton mobility via vehicu­ lar mechanism becomes more important.38 In this section, PEM chemistry and architecture effects on proton conductivity under fully and partially hydrated conditions are explained using length scale parameters to compare a wide range of PEMs and demon­ strate fuel cell performance using selected PEMs.

O

S O

O H

O

H O

H H

S

O

O

O H

Figure 6 Schematic diagram of proton conduction in hydrated sulfo­ nated PEMs.

estimate of equivalent volume (cm3 per ionomer or the mol equivalent of acid groups) based on the summation of molar volume subunits (Table 1) rather than true volume measure­ ments, as shown in eqn [3]: Table 1 Groups

Molar volume increments of selected groups39 Va(298) (cm3 mol)−1

5.5

Groups

Va(298) (cm3 mol)−1

32.5

9.12

16.37

18.72

23.7

73.3

19.85

5.28

21.87 84.16

65.5

30.7 20.0

69

49.0 100.5

112

17.3

10.36.2.2.1

Under fully hydrated conditions

A simplified model for proton conduction is illustrated in Figure 6 that considers two contributing factors, interanionic distance (δ) and number of participating ion-carrying media (hydration number, λ). In this model, proton conductivity can be expressed by λ ConductivityðσÞ α ½2 δ where λ is hydration number (number of water molecules per sulfonic acid group) and δ is the distance between sulfonic acid groups in hydrated polymer. The interanionic distance in dry polymers can be approxi­ mated by Molar Volume per Charge (MVC). MVC is an

8.5

6.88

8.0 34.08

20.22

25.76

8.32 40.5

26.56

Polymers in Energy Applications | Polymers in Membrane Electrode Assemblies MVC ¼

X

ni vi

½3

i

where vi is the volumetric contribution of the ith structural group that appears ni times per charge. The distance between sulfonic acid groups in hydrated poly­ mers can be approximated by eqn [4]: MVCWET ¼ MVC þ VH2 O  λ

½4 −1

where VH2O is the molar volume of water, 18 cm mol . The MVCWET parameter assumes all absorbed water resides in hydrophilic domains (which is reasonable) and neglects polymer morphological features (which is not valid in general but the morphological effect can be considered minimized at high water content where ionic clusters are well developed). Then conductivity is proportional to Percent Conducting Volume (PCV) that is defined as 3

PCV ¼

VH2 O  λ ðMVC þ VH2 O  λÞ

½5

The numerator of the eqn [5] is an estimate of the volume of the aqueous (hydrophilic) phase per acid site, and the denomi­ nator is an estimate of the volume of the hydrated membrane (both hydrophilic and hydrophobic phases) per acid site. PCV is essentially a ratio of the volume of the hydrophilic phase to the hydrated membrane analogous to the conducting volume of the membrane. Figure 7 plots the conductivity of a wide variety of polymers electrolytes as a function of PCV under fully hydrated

Relative conductivity

10

1

0.1 PFSA Partially fluorinated HC-based (random) Functional HC-based (cross-linked) Regression

0.01

0.001 0.0

0.1

0.2

0.3

0.4

0.5

0.6

0.7

Percent conducting volume (PCV) Figure 7 Relative proton conductivity vs. percent conducting volume (PCV) of various sulfonated polymers under fully hydrated conditions at ambient temperature.20 Reprinted from Ann. Rev. Chem. Biomol. Eng., 2010, Y. S. Kim et al., Moving beyond mass-based parameters for con­ ductivity analysis of sulfonated polymers, 123–148, 2010, with permission from Annual Reviews.

Table 2

697

conditions. A general correlation exists between proton con­ ductivity in fully hydrated PEMs and PCV, in spite of the data scattering. Proton conductivities of arbitrarily categorized PEMs under fully hydrated conditions do not show significant differ­ ences at a given PCV, and correlate with a master curve (shown by the black dash line). This indicates that there is an apparent diminished role of molecular structural/morphology effects on proton conductivity for full hydrated PEMs using the PCV parameter, probably due to the fact that incorporation of sig­ nificant amounts of water (λ  > l0) in fully hydrated PEMs diminishes the importance of structural characteristics by the formation of more continuous, less tortuous hydrophilic domains. Consequently, the conductivity of all different kinds of polymers under fully hydration can be approximated by a single PCV parameter. Proton conductivity increases 1 order of magnitude as PCV increases from 0 to 0.35 and then increases more slowly. The marginal increase of proton conductivity beyond a PCV value of 0.35 is due to the volume dilution of sulfonic acid groups with water, such that proton conduction occurs less effectively. In this higher PCV range, PEMs exhibit excessive water uptake and their mechanical properties (mod­ ulus and strength) deteriorate accordingly. As a result, most well-performing PEMs for H2/air fuel cells in a fully hydrated environment have PCV values of 0.35 where conductivity is relatively high (0.05–0.1 S cm−1) while mechanical properties are reasonably good. Table 2 shows properties of selected PEMs at PCV values of around 0.33. The PEMs selected have reason­ ably low water uptake, which helps to remove issues related to excessive water uptake during fuel cell operations. It is seen that all PEMs in Table 2 have similar conductivity of 80 mS cm−1, significantly different MVC and λ values. The functional PEM has relatively low MVC and λ, while PFSA PEM has relatively high MVC and λ. In other words, a similar conductivity can be obtained with higher sulfonic acid concen­ tration and lower water uptake or conversely, with lower sulfonic acid concentration and higher water uptake. The water absorption capacity of each PEM is consistent with the results in the water uptake section, which showed that func­ tional or cross-linked PEMs have relatively lower water uptake than HC-based PEMs, which in turn have lower water uptake than fluorinated PEMs. Figure 8 compares H2/air fuel cell performance using the selected PEMs. All PEMs have a similar thickness (50 µm) and fabricated in a same manner and tested under same operating conditions. It is noted that H2/air fuel cell performance and high frequency resistance (HFR) using these membranes are comparable with each other, as expected.

Comparison of PEM properties14 IEC

MVC

PEM

(meq. g−1)

(cm3 equiv. mol.−1)

λ

Poly(arylene ether nitrile) homopolymer Poly(arylene ether) cross-linked polymer Poly(arylene ether sulfone) random copolymer PFSA Nafion copolymer

2.7 1.66 1.54 0.91

277 435 486 524

6.2 10.4 14.4 15.0

WU

σ

PCV

(wt.%)

(mS cm−1)

0.29 0.30 0.35 0.34

43 30 40 26

81 74 72 91

698

Polymers in Energy Applications | Polymers in Membrane Electrode Assemblies

1000 conductivity (mS cm–1)

Cell potential (V)

1.0

0.8

0.6

a) SPAEN 1.0 b) CSFQH80-BP c) BPSH-35 d) Nafion 212

HFR (Ω cm2)

0.4

0.2

0.2

0.4

0.6

0.8

Current density (A

Perfluorinated (grafted/side chain) Aromatic (block) Partially fluorinated aromatic (random) Aromatic (random) Aromatic (homo/alternating)

0.1

0.1

0.2

0.3

0.4

0.5

Percent conducting volume (PCV)

1.0

1.2

cm−2)

Under partial humidified conditions

In H2/air fuel cells, low relative humidity (RH) operations are desirable. At low RH operation, parasitic power losses from gas humidification is reduced, electrode ‘flooding’ at high current densities (liquid water begins to collect in the electrode and GDL and limit access of gases to the catalyst) is prevented and operating temperature can be increased. At high operating temperature, reaction kinetics of electrocatalyst are enhanced, water and thermal management are simplified, useful waste heat can be recovered, and lower quality reformed hydrogen may be used as the fuel. Figure 9 shows the conductivity versus PCV of various PEMs as the amount of absorbed water decreased. Note the much more significant data scattering at low PCV values compared with Figure 7 (e.g., conductivity varies from 0.1 to 10 mS cm−1 at a PCV of 0.1) and the scattering becomes less at high PCV values, which is consistent with the above argu­ ment that PCV can be used as a single parameter for Table 3

1

0.0

Figure 8 H2/air fuel cell performance under fully humidified conditions at 80 °C using sulfonated poly(arylene ether nitrile) homopolymer (SPAEN 1.0), cross-linked partially sulfonated poly(arylene ether) random copo­ lymer (CSFQH80-BP),14 sulfonated poly(arylene ether sulfone) random copolymer (BPSH-35),13 and PFSA (Nafion 212).

10.36.2.2.2

10

0.01

0.0 0.0

100

Figure 9 Proton conductivity vs. percent conducting volume of various sulfonated polymers under partially hydrated conditions at ambient temperature.20 Graph is taken from Ann. Rev. Chem. Biomol. Eng., 2010, Y. S. Kim et al., Moving beyond mass-based parameters for conductivity analysis of sulfonated polymers, 123–148, 2010, with permission from Annual Reviews. Individual data were taken from 12 independent references in the literatures.

conductivity prediction at fully hydrated conditions. The large data scattering at low PCV (i.e., low RH) are attributable to proton hopping becoming very sluggish and the contribution the vehicular conduction mechanism via diffusion of water molecules plays a bigger role under low RH conditions. Closer scrutiny of the graph indicates a clear tendency that reflects the morphological features of each categorized polymers. PFSAs showed the highest of conductivity and aromatic (multiblock) and partially fluorinated (random) showed the second highest conductivity, while aromatic ran­ dom copolymers and aromatic homopolymer-like structure have the least conductivity. Table 3 shows the properties of various PEMs at 50% RH. There are a few important observations from Table 3, regarding polymer structural effects. 1. When comparing PEMs having similar PCV, fluorinated polymers have higher λ and MVC values. For example, λ and MVC of Nafion 112 are 3.5 and 524, respectively, which are higher than of those of HC-based BPSH-35 at a PCV of 0.11. This suggests that fluorinated PEMs are bene­ ficial for low RH fuel cell operations.

Properties of HC-based, partially fluorinated, and PFSA PEMs at 50% RH

PEM

Category

Architecture

IECW (meq. g−1)

MVC (cm3 equiv. mol.−1)

λ

PCV

WU (wt.%)

σ (mS cm−1)

Ph-PEEKDK BPSH-35 BPSH-40 BPSH-PI (15k-15k) BPSH-BPS (3k-3k) BPSH-BPS (5k-5k) BPSH-BPS (10k-10k) 6F-40 Nafion 112 Nafion 212 Dow 700

HC based HC based HC based HC based HC based HC based HC based Fluorinated PFSA PFSA PFSA

Homo Random Random Multiblock Multiblock Multiblock Multiblock Random Side chain Side chain Side chain

1.60 1.54 1.72 1.55 1.33 1.39 1.28 1.30 0.91 0.95-1.0 1.42

488 486 430 430 414 414 414 522 524 477 335

3.2 3.3 4.2 3.4 2.8 3.0 3.8 3.4 3.5 3.7 3.3

0.11 0.11 0.15 0.12 0.11 0.12 0.14 0.11 0.11 0.12 0.15

9.2 9.1 11.7 9.4 6.7 7.2 9.1 8.0 5.7 6.1 8.5

0.2 0.9 1.3 2.3 5 9 14 1.8 10 12 33

Polymers in Energy Applications | Polymers in Membrane Electrode Assemblies

Figure 10 shows the H2/air fuel cell performance using selected PEMs from Table 3 at two different gas inlet humidites (70% and 40% RH). When both the inlet gases were maintained at 70% RH, the random and multiblock copolymers performed similarly, while much lower cell performance was observed using the homopolymer-like structure. This is probably attri­ butable to the homopolymer-like structure having a much lower conductivity, even at moderately reduced humidity. This is in contrast with the similar cell performance being obtained for the homopolymer-like structure and other PEMs when the cells are operated under fully hydrated conditions. When anode and cathode inlet humidity levels are reduced to 40% RH, the performance of the random copolymer suffered

significantly, while the fuel cell performance of multiblock copolymer and PFSA PEMs decreased to a lesser degree. This demonstrates the importance of PEM structures under reduced RH fuel cell operations.

10.36.2.3

Unlike PEMs used in H2/air fuel cells, where a single key property (i.e., proton conductivity) has been the primary emphasis, PEMs for direct methanol fuel cells (DMFCs) have been investigated with a strong emphasis on methanol perme­ ability. In DMFCs, high methanol crossover through membrane is an issue. High methanol crossover results in not only wasting fuels but also decreasing fuel cell performance/ durability by increasing cathode overpotential, decreasing membrane conductivity, and increasing solvent swelling level of membrane. Methanol permeability measured on free-standing films is derived from Fick’s diffusion law j ¼ −D

j¼

DH Δc l

1.0

0.4 0.3

0.4 Ph-PEEKDK 0.2 BPSH-35 BPSH-PI (15k–15k) 0.1 Nafion 212

0.2

0.0

0.0 0.0

0.2

0.4

0.6

0.8

Current density (A cm−2)

1.0

1.2

Nafion 212

0.8 Cell potential (V)

Cell potential (V)

0.6

½7

0.7 BPSH-35 BPSH-PI (15k–15k) 0.6

40% RH

0.6 0.5

½6

where H is the partition coefficient, l the thickness of the membrane, and Δc is the concentration difference between the solution in contact with the membrane. Methanol permeability (DH) is defined as the product of the diffusion coefficient and partition coefficient and takes into account both solubility and diffusivity. Although, in operating fuel cells, there are combinations of other forces that can influ­ ence transport such as electro-osmotic drag or hydraulic permeation due to pressure difference, the driving force by the concentration difference is largely counted for screening of DMFC PEMs. Methanol permeability of PEMs is also influenced by

70% RH

0.8

dc1 dz

where j is the flux, D the diffusion coefficient, c1 the concentra­ tion of methanol, and z the position within the membrane. This equation can be integrated to yield

0.7 1.0

Methanol Permeability

0.5 0.4

0.6

0.3 0.4 0.2 0.2

Power density (W cm−2)

2. Increasing IEC improves the low RH conductivity as PCV value increases; For example, as the IEC of BPSH increases from 1.54 to 1.72 meq. g−1, PCV increases from 0.11 to 0.15 since MVC value decreases from 486 to 430 and λ value increases from 3.3 to 4.2. Another example is the PFSA PEMs (i.e., Nafion 112, 212, and Dow 700). This suggests that increasing IEC is beneficial for low RH fuel cell operations. 3. As block length increases in multiblock copolymers, λ (and thus PCV) increases. Consequently, higher conductivity is expected. For example, BPSH-BPS copolymers showed these trends; increasing the block length from 3 to 10 k led to improved conductivity from 5 to 14 mS cm−1. This suggests that increasing block length is beneficial for low RH fuel cell operations. The λ value alone (or even with MVC) cannot predict the conductivity. For example, the conductivity of BPSH-35 is greater by more than 4 times than that of Ph-PEEKDK at 50% RH, in spite of the similar λ and MVC values. Another example is a comparison of BPSH-40 random copolymer and BPSH-PI multiblock copolymer. This suggests that mor­ phological features and state of water can be important factors that are not reflected in the λ numerical values.

699

0.1 0.0

0.0 0.0

0.1

0.2

0.3

0.4

0.5

0.6

Current density (A cm−2)

Figure 10 H2/air fuel cell performance at 100 °C and under 70% and 40% RH conditions using sulfonated poly(arylene ether ketone) homopolymer (Ph-PEEKDK), sulfonated poly(arylene ether) random copolymer (BPSH-35), sulfonated poly(arylene ether sulfone)-polyimide multiblock copolymer (BPSH-PI), and PFSA (Nafion 212).40 Reprinted with permission from Chem. Mater., 20, M. L. Einsla et al., Toward improved conductivity of sulfonated aromatic proton exchange membranes at low relative humidity, 5636–5642, 2008, Copyright 2008 American Chemical Society.

700

Polymers in Energy Applications | Polymers in Membrane Electrode Assemblies

chemical structure and morphology, although there are limited systematic approaches to correlate methanol permeability and with chemical/morphological features of sulfonated polymers. Figure 11 shows the relative methanol permeability (PM) of various PEMs as a function of PCV. The normalized methanol permeability to benchmark Nafion® (EW = 1100) was used in order to reduce the systematic errors derived from different test conditions (temperature and methanol concentration) and meth­ odology. Methanol permeability increases as the PCV value increases without a definite slope transition, which implies that dilution by the absorbed water at a PCV value of 0.35 does not have much influence on methanol permeability. This is because proton transport occurs largely via the proton hopping mechan­ ism under fully hydrated conditions, where both the concentration of proton exchange sites (sulfonic acid group) and that of the proton carrier (water molecules) are equally important. Thus, dilution of sulfonic acid groups by water reduces the transport rate, while methanol transport occurs solely via molecular diffusion, where the methanol diffusion predomi­ nantly relies on the amount of water absorbed in the PEMs. Figure 11 shows a clear trend of structural effects on methanol permeability unlike the conductivity plot: the methanol perme­ ability increases in the order of PFSA > partially fluorinated > HCbased > functional PEMs. Blends of PFSA with methanol barrier polymers or additives (classified here as partially fluorinated PEMs) exhibit lower methanol permeability than PFSAs at a given PCV, suggesting that using PFSA blends have the potential to improve DMFC performance over unmodified PFSAs, in spite of the rather limited methanol barrier efficiency compared to HC-based or functional PEMs. HC-based PEMs show lower methanol permeability than PFSA and partially fluorinated PEMs. Functional PEMs having additional specific interactions showed the least methanol permeability. One exception is cross-linked PEMs. Although most cross-linked PEMs displayed here do not have additional specific interactions, the methanol permeability is substantially lower than those of functional PEMs. Since methanol permeability in addition to proton conduc­ tivity impacts fuel cell performance, selectivity, which is defined as the ratio of proton conductivity to methanol

permeability, has been suggested as a single key property of PEMs for DMFCs to compare one polymer versus another.42 σ Selectivity ðsÞ ¼ ½8 DH It is expected the membrane having high selectivity performed better in DMFC because PEMs having higher selectivity allows less methanol crossover at a given conductivity. The greatest limitation of selectivity as a general gauge of DMFC PEMs is the requirement of a minimum conductivity necessary for effective operation of the DMFC, regardless of how low the methanol permeability is.43 This minimum conductivity is related to the fact that the membrane itself has a minimum attainable thick­ ness to achieve low cell resistance. Selectivity also fails to account for issues of mechanical/chemical robustness or the ability to be fabricated into high-performance MEAs. Therefore, selectivity serves only as a screening tool for potential DMFC PEMs. Figure 12 shows the relative selectivity (Φ) of various PEMs as a function of PCV. The relative selectivity plot shows a general trend that relative selectivity decreased as PCV increased due to the fact that the slope of methanol permeability is greater than those of proton conductivity. As a result, only a few highly selective PEMs are available at high PCV range (e.g., only 2 out of 48 PEMs have selectivity of > 4 at PCV at 0.4 or greater). This suggests that PEMs having high PCV values may be less attractive for DMFC applications not only due to their excessive water uptake but also due to their low selectivity. PEMs having low PCV ( < 0.3), on the other hand, have limited to use due to their low conductivity in spite of their high selectivity values. For example, a PEM having PCV of 0.1 and a relative conductivity of 0.05 would need to be 20 times thinner than Nafion PEM in order to incur similar ohmic loss. In practical terms, this is insufficient sustain mechanical integ­ rity. The effects of molecular compositions and cross-linking are well reflected with the selectivity plot, in spite of the rela­ tively large data scatterings for HC-based and functional PEMs. This large data scattering for HC-based and functional PEMs is likely due to the significant structural differences existing within these categories.

12 1

0.1 PFSA Partially fluorinated HC-based (random) Functional HC-based (cross-linked)

0.01

Relative selectivity

Relative methanol permeability

PFSA Partially fluorinated HC-based (random) Functional HC-based (cross-linked)

15

10

9 6 3 0

0.001 0.0

0.1 0.2 0.3 0.4 0.5 0.6 Percent conducting volume (PCV)

0.7

Figure 11 Relative methanol permeability PM vs. PCV of various sulfo­ nated polymers under fully hydrated conditions as a function of polymer category.41 Reprinted from Kim, Y.S.; et al J. Membr. Sci. 2010, 374, 49–58, with permission from Elsevier.

0.0

0.1

0.2 0.3 0.4 0.5 0.6 Percent conducting volume (PCV)

0.7

Figure 12 Relative selectivity vs. PCV of various sulfonated polymers under fully hydrated conditions as a function of polymer category.41 Reprinted from Kim, Y.S.; et al J. Membr. Sci. 2010, 374, 49–58, with permission from Elsevier.

Polymers in Energy Applications | Polymers in Membrane Electrode Assemblies

701

Table 4 Average values of PCV, MVC, λ, relative proton conductivity (σ), relative methanol permeability (PM), and relative selectivity (Φ) for PEMs (PCV = 0.35)41

PEMs

No. data

PCV

MVC (cm3 equiv. mol.−1)

λ (H2O/SO3H)

σ

PM

Φ

PFSA Partially fluorinated HC based Functional Cross-linked

3 11 12 14 12

0.35 (0.03) 0.35 (0.02) 0.35 (0.01) 0.35 (0.03) 0.35 (0.02)

550 (40) 531 (73) 459 (85) 407 (83) 397 (45)

16.5 (1.8) 15.9 (2.0) 13.7 (2.3) 12.5 (3.2) 11.6 (2.4)

0.85 (0.18) 0.87 (0.23) 0.86 (0.24) 0.83 (0.33) 0.98 (0.18)

0.88 (0.11) 0.73 (0.16) 0.35 (0.12) 0.24 (0.12) 0.21 (0.06)

0.9 (0.1) 1.2 (0.2) 2.5 (1.0) 5.1 (3.5) 6.3 (2.5)

Numbers in the parenthesis denote the standard deviation.

Table 4 shows the averaged values of MVC and λ values of the categorized PEMs at the PCV of 0.35 for discussion about the molecular design aspect. A decreasing trend of methanol permeability with molecular composition (i.e., PFSA > partially fluorinated > HC-based > functional) is consistent with the trend observed in Figure 11 for PEMs in the whole range of PCV. The MVC and λ decreased as molecular composition changed from PFSA to HC-based and to functional PEMs. Considering that the conductivity of PEMs are similar at a given PCV, this indicates that conductivity of PFSA is obtained with relatively low sulfonic acid volume concentration (i.e., high MVC) and high water content (i.e., high λ), while similar conductivity was obtained with functional PEM where the sulfonic acid volume concentration is high (i.e., low MVC) and water content is low (i.e., low λ). Although one should not assume that a single λ value can be used for quantitative estimation of methanol permeability because differences in dielectric constants, partitions, and states between water and methanol molecules were not accounted for in λ, the differ­ ences in methanol permeability seem to correlated well with the water holding capability of each categorized PEM system, depending on molecular composition. Table 5 shows the properties of functional PEMs at PCV of around 0.3. While all functional PEMs exhibit effectively

reduced methanol permeability by decreasing MVC and λ, the PEMs with high dipole interactions such as phosphine oxide or amide have a more profound effect than those with low specific interactions such as nitrile or carboxylic acid groups. For exam­ ple, the PEM with 2 vol.% phosphine oxide group has MVC and λ of 440 (cm3 equiv.−1 mol.) and 9.6 (H2O/SO3H), while the PEM with 8 vol.% nitrile group has a greater MVC and λ of 493 (cm3 equiv.−1 mol.) and 12.2 (nH2O/SO3H), respectively. High dipole interactions may also impact the proton conductivity as the proton conductivity of these functional PEMs is noticeably lower than those of other functional PEMs. This is not readily apparent in Figure 11, but it appears as slightly lower conduc­ tivity with greater data scattering for functional PEMs. This result suggests that strong dipole interactions, particularly those invol­ ving interaction of the functional group with sulfonic acid groups, may not be desirable for DMFC PEMs. Table 6 compares the MVC, λ, and PCV values for HC-based cross-linked, homo, and random (co)polymers at PCV of 0.25. It is evident that there is a clear trend that MVC and λ increase in the order of cross-linked copolymer, homopolymer-like struc­ ture, and random copolymer (MVC increased from 328 to 638 and λ increased from 5.6 to 12.0). Relative methanol perme­ ability increased accordingly from 0.16 to 0.19. This result suggests that lower methanol permeability is expected with

Table 5 PCV, MVC, λ, relative proton conductivity (σ), relative methanol permeability (PM), and relative selectivity (Φ) for HC-based and functional PEMs at PCV of 0.3041 Functional group (backbone)

Vol. %

PCV

MVC (cm3 equiv. mol.−1)

λ (H2O/SO3H)

σ

PM

Φ

Carboxylic acid (ketone) Nitrile (phenylene) Amide (ketone) Phosphine oxide (sulfone) Imidazole (sulfone)

5 8 4 2 14

0.31 0.31 0.32 0.28 0.28

518 493 408 440 441

12.6 12.2 10.5 9.6 9.1

0.36 0.84 0.19 0.26 0.43

0.20 0.24 0.18 0.16 0.10

1.8 3.5 1.0 1.7 4.3

Table 6 PCV, MVC, λ, proton conductivity (σ), normalized methanol permeability (PM), and relative selectivity (Φ) for HC-based homopolymers, random copolymers, multiblock copolymers, and cross-linked copolymer at PCV of 0.2541 Polymer architecture

PCV

MVC (cm3 equiv. mol.−1)

λ (H2O/SO3H)

σ

PM

Φ

Cross-linked copolymer Homopolymer Random copolymer

0.25 (0.06) 0.25 (0.05) 0.26 (0.03)

328 (67) 462 (35) 638 (75)

5.6 (1.7) 8.7 (1.5) 12.0 (1.9)

0.66 (0.2) 0.39 (0.1) 0.46 (0.2)

0.16 (0.04) 0.18 (0.05) 0.19 (0.08)

6.3 (3.8) 4.7 (0.6) 2.6 (1.0)

702

Polymers in Energy Applications | Polymers in Membrane Electrode Assemblies

cross-linked copolymers and homopolymer-like structures compared to non-cross-linked random copolymers and these would possibly have lower methanol permeability than multiblock copolymers. Unlike some functional PEMs, conductivity reduction due to the cross-linking was not observed, which leads to excellent selectivity for cross-linked PEMs. Although the availability of PEM data having different polymer architec­ ture is somewhat limited to generate statistical data, this result indicates the importance of polymer architecture in addition to molecular composition of sulfonated polymers on methanol barrier properties. These observations and trends described above provide guidance and the following suggestions for the design of advanced PEMs for DMFC applications: 1. The absence of fluorine in PEMs can significantly improve PEM properties for DMFC. Although replacing hydrogen with fluorine increases volume sulfonic acid concentration (and thus conductivity), the fluorination cannot adequately compensate high methanol permeability derived from lar­ ger water domain structures. 2. Methanol barrier properties of HC-based PEMs can be sig­ nificantly improved by changing polymer molecular structure or introducing hydrophobic substituent. 3. Incorporation of functional groups may also improve PEM properties. Functional groups having strong dipole interac­ tions effectively reduce the MVC and water uptake. However, these often reduce proton conductivity due to their strong interaction with sulfonic acid groups, producing relatively moderate increases in selectivity. Another issue with functional groups having strong dipole interactions is the composition range for incorporation is very narrow.44 In this case, it requires very precise control of incorporation of the functional group. On the other hand, functional groups with relatively weak dipolar interactions such as nitrile

group can provide a relatively broader compositional win­ dow and uniformity. 4. Changing polymer architecture to cross-linked or homopolymer-like structure from random may improve methanol barrier properties. Cross-linking is more aggressive way to improve methanol-blocking properties, although sig­ nificant reduction of MVC and λ are typically expected. Like the PEMs with strong specific interactions, the compositional window for cross-linking is narrow and thus, low levels of cross-linking may be desirable to obtain good PEM properties. Figure 13 demonstrates the DMFC performance of MEA using sulfonated poly(arylene ether nitrile) (m-SPAEEN-60, relative selectivity = 2.9), sulfonated poly(arylene ether sulfone) (BPSH-35, relative selectivity = 1.3) and PFSA (Nafion, relative selectivity = 1). Performance comparison of PEMs under DMFC operating conditions is complex since the influence of methanol crossover and cell resistance varies throughout current density range. Furthermore, operating conditions such as methanol-feed concentration, temperature, and so on have an influence on the methanol crossover and cell resistance different degrees. Typically, optimum performance for a particular PEM can be achieved by adjusting membrane thickness (and thus changing the relative values of methanol crossover and cell resistance) at certain operating conditions. When comparing several PEMs, it is not readily apparent form the polarization curves as to whether each thickness is optimized or not. In order to have a more meaningful comparison across several different PEMs, therefore, membrane thickness was selected to have similar methanol crossover (all cells have similar methanol crossover current of 50 mA cm −2 at 0.5 M methanol feed) and the DMFC performance was compared at two different methanol-feed con­ centrations. In this case, the performance difference is largely attributed to changes in ohmic losses primarily due to the PEM resistance. The cell using sulfonated polynitrile showed

0.8 0.7 0.6

Cell potential (V)

0.9 m-SPAEEN-60 (53 μm) BPSH-35 (74 μm) Nafion (250 μm)

0.7 0.6 0.5

0.4

0.4

0.3

0.3

0.2

0.2

0.1 0.0 0

100

200

300

Current density (mA cm−2)

400

m-SPAEEN-60 (53 μm) BPSH-35 (74 μm) Nafion (250 μm)

0.8

0.5

HFR (Ω cm2)

HFR (Ω cm2)

Cell potential (V)

0.9

0.1 0.0 0

100

200

300

400

500

600

Current density (mA cm−2)

Figure 13 DMFC performance of MEA using sulfonated poly(arylene ether nitrile) (m-SPAEEN-60, relative selectivity = 2.9), sulfonated poly(arylene ether sulfone) (BPSH-35, relative selectivity = 1.3) and PFSA (Nafion, relative selectivity = 1) at 0.5 and 2 M methanol feed concentration (cell temperature: 80 °C).45 Reprinted with permission from J. Electrochem. Soc., 2008, 155, B21. Copyright 2008, The Electrochemical Society.

Polymers in Energy Applications | Polymers in Membrane Electrode Assemblies

the best DMFC performance and the cell using PFSA Nafion showed the worst performance at both methanol-feed concen­ trations. The results in Figure 13 are consistent with the PEM property analysis.

703

Carbon

Pt

10.36.2.4

Summary

The primary goal of this section is to provide a general relation­ ship between structures and properties of PEMs for their use in H2/air fuel cells under fully hydrated and partially hydrated conditions, and in DMFCs. PEMs are classified under four categories by their chemistry: polymers having functional groups, HC-based, partially fluorinated, and PFSA PEMs. PEMs are further categorized by polymer architectures, that is, cross-linked, homopolymer, random copolymers, grafted/side chain copolymers, and multiblock copolymers. When water uptake is compared at a given IEC, clear trends emerged. For the chemistry effect: water uptake increased from polymers having functional groups to HC-based and to fluorinated. For the polymer architecture effect: water uptake increased from cross-linked to homopolymer-like structure to random and grafted/side chain copolymer and to multiblock copolymer. From a combination of these chemical structure and architec­ ture effects, high proton conductivity can be obtained from each polymer under fully hydrated conditions where the ratio of water volume to sulfonic acid concentration in the hydrated state (PCV) has a value of about 0.35. In general, polymers having functional groups have higher IECW (1.8–2.2 meq. g−1), polyaromatics have lower IECW (1.5–1.8 meq. g−1), and fluorinated polymers have the lowest IECW values (0.8–1.5 meq. g−1) to attain the required PCV value. Although similar conductivity under fully humidified conditions could be obtained regardless of polymer chemistry and polymer architecture for most sulfonated polymers by adjusting IECW, there are significant differences in local sulfo­ nic acid concentration or size of water domains due to morphological differences, which impacts conductivity at low RH and methanol crossover. For H2/air fuel cell applications under low RH operations, polymers having high local sulfonic acid concentration are beneficial since higher conductivity can be obtained at low RH; while for DMFC applications, polymers having uniform SO3H concentration are beneficial since lower methanol permeability can be obtained.

10.36.3 Polymer Electrolyte Ionomers in the Electrode Controlling the structure of the electrode is critical to obtaining desirable catalyst utilization, proton/water/reactant gas transport and mechanical durability. Figure 14 shows a TEM image of a fuel cell electrode consisting of platinum electrocatalyst, carbon catalyst support, and ionomer binder. Reactants such as hydro­ gen or oxygen supplied from the GDL are transported through the pores in the electrode to the interface with the electrode ionomer. Then the reactants permeate through the thin ionomer layer (ideally a few tens of nm scale) to reach the catalyst– ionomer interface. Platinum (Pt) catalyst is typically finely dispersed onto carbon-supporting materials to maximize the electrochemically available surface area. In the cathode catalyst layer, oxygen is electrochemically reduced to H2O.

Binder

10 nm Figure 14 TEM image of PEMFC electrode structure.46

The electrode ionomers must have a high proton conduc­ tivity to minimize loss due to high ion resistances in the electrodes. In addition, the ionomers must provide high fuel and oxidant permeability for the electrochemical reaction, which is in a stark contrast to PEMs that require minimal fuel and oxidant permeability. Good water transport properties of the ionomers are desirable for PEM hydration and for fast oxidant permeability. Good oxidative and hydrolytic stability remain as essential properties for electrode stability. The iono­ mers also should provide mechanical robustness and good adhesion to the PEM in order to maintain the electrode struc­ ture under fuel cell operating conditions. Chemical inertness and contaminant-free properties to the electrocatalyst are also required for the ionomer. Electronic conductivity of electrode ionomers is preferable but not essential, since the electrocata­ lyst and its supporting materials provide electron conductivity within the electrode.

10.36.3.1

Historical Background

Incorporation of a polymer electrolyte as an electrode ionomer was developed in the mid-1980s and substantially improved fuel cell performance. Schematic electrode structures of early efforts are illustrated in Figure 15. Until the mid-1980s, fuel cell electrodes contained a high loading of platinum catalyst (4 mg cm−2) (Figure 15(a)). Platinum catalysts were usually mixed with PTFE and hot pressed onto the membrane to make the MEA.47 In the late 1980s, Raistrick developed porous GDEs containing a low Pt loading (0.35 mg cm−2) by incorporating Nafion in the electrode (Figure 15(b)).48 Further optimization of the GDEs in terms of the amount of Nafion impregnated, hot-pressing conditions, operating temperature and pressure, and reactant humidification was achieved by Ticianelli et al.,49,50 (Figure 15(c)). Figure 15(d) shows that the GDEs (0.35 mg cm−2 Pt loading) prepared from the above procedure and compared with Prototech electrode (0.35 mg cm−2 Pt loading) and the GE-HS-UTC electrode with high Pt load­ ing (4 mg cm−2 Pt loading). The GDE electrode with Nafion-impregnated electrodes exhibit much improved perfor­ mance compared to the Prototech electrode having the same Pt loading. The fuel cell performance with Nafion-impregnated electrodes showed the same level of performance as an

704

Polymers in Energy Applications | Polymers in Membrane Electrode Assemblies

(d) 1.0 (a)

(b)

Carbon

Teflon

0.9 Platinum

0.8 A

0.7

Cell potential (V)

lonomer

(c)

0.6

B

C

0.5 0.4

–2

0.3

A. Nafion impregnated cell (0.35 mg Pt cm ) B. PTFE impregnated GE/HS-UTC cell (4 mg Pt cm–2)

0.2

C. As-received Prototech cell (0.35 mg Pt cm–2)

0.1 0.0 0

40

80 120 160 Current density (mA cm–2)

200

Figure 15 Cross-sectional schematics of the membrane–electrode interface for three constructions. (a) Unsupported Pt catalyst bonded to the membrane. (b) Carbon cloth electrode incorporating Pt–C catalyst layer attached onto membrane (catalyst layer thickness 50–100 µm) and (c) Thin catalyst film electrode in which a thin Pt/C-Nafion catalyst layer is attached to the membrane (catalyst layer thickness 3–5 µm).44 (d) Cell potential/current density plots for H2/O2 fuel cells at 50 °C, 1 atm pressure: A and C cells with Nafion-impregnated and as-received Prototech electrodes, respectively (0.35 mg-Pt cm−2, respectively; B cell with GE-HS-UTC membrane electrode assembly (4 mg-Pt cm−2).50 With kind permission from Springer Science, J. Appl. Electrochem., 22, 1–7 (1992); Reprinted with permission from J. Electrochem. Soc., 1988, 135, 2209. Copyright 1988, The Electrochemical Society.

electrode structure (Pt loading: 0.15 mg cm−2) showed good long-term performance ( > 2000 h of operating with overall power losses of 10%).51,52 Figure 16 shows the (a) catalyst utilization and (b) long-term fuel cell performance of electro­ des prepared by the thin-film electrode technique. Figure 16(a) shows that the differences in the specific activities for each type of electrode are dramatic and, clearly, the use of Nafion elec­ trode ionomer and the type of construction significantly impact the utilization of platinum electrocatalysts. Figure 16(b) shows the current density change during extended fuel cell operation for the Na+ and TBA+ forms of Nafion-bonded electrodes. Again significant differences in fuel cell life are noted, depend­ ing on the type of electrode ionomer used. During the period

electrode with high Pt loading. Hot pressing the electrodes to the membrane (rather than cold mechanical compression) is thought to greatly reduce activation and ohmic overpotential losses. In the early 1990s, Wilson developed the thin-film electrode technique that further improved replaced the porous GDE.44 The thin-film electrode was attached to the membrane by either direct application of Nafion-containing catalyst ink to the membrane or decal transfer using a Teflon blank and subse­ quent hot pressing at 125–145 °C. In the mid-1990s, Wilson replaced the proton form of Nafion with the tetrabuthylammo­ nium (TBA+) form, which allowed hot pressing at higher temperatures, 200–210 °C. The MEAs made with the new

(a)

(b)

Cell voltage (V)

0.8 0.6

1 0.4 2

3 0.2 0.0

0

4

8

Specific acitivity (A mg–2 Pt)

12

Current density (mA cm–2) at 0.5 V

1600 1. thin film electrode (0.15 mg/Ptcm–2) 2. carbon cloth electrode (0.35 mg Pt cm–2) 3. unsuported Pt catalyst (4 mg Ptcm–2)

1.0

1400 1200 1000 800 600 Na form TBA form

400 200 0

0

500

1000

1500

2000

2500

Hours

Figure 16 Effect of electrode ionomer on (a) catalyst specific activity: comparison between thin film electrode, carbon cloth electrode and unsupported Pt electrode.44 With kind permission from Springer Science, J. Appl. Electrochem., Thin-Film Catalyst layers for polymer electrolyte fuel-cell electrodes, 22, 1992, 1–7, M. S. Wilson et al., Fig. 6 and (b) electrode durability: comparison between Na+ and TBA+ forms of Nafion.51 Reprinted from Electrochimica Acta, 40, M. S. Wilson et al., Low platinum loading electrodes for polymer electrolyte fuel cells fabricated using thermoplastic ionomers, 355, Copyright (1995), with permission from Elsevier.

Polymers in Energy Applications | Polymers in Membrane Electrode Assemblies

from the early 1980s to the mid-1990s, the framework for the modern electrode structure had been established. During the period from mid-1990s to late 2000s, systematic approaches to investigate the electrode structure have been made with PFSA and HC-based electrode ionomers as the PEMFC technology was implemented in practical energy systems. Nevertheless, studies of the structural effects of the electrode ionomers are still ongoing. In the following sections, the structural effect of electrode ionomers on PEMFC performance is discussed based on current knowledge.

10.36.3.2 10.36.3.2.1

loading of HC-based S-PEEK binder is (partially) explained by the density of S-PEEK, which is typically 60% of that of Nafion. However, further reduction of ionomer loading is often required for S-PEEK (as seen in Figure 17(b)) since HC-bonded electrodes are further limited by slow oxygen transport. As the ionomer loading is increased to the optimum value, the fuel cell performance improves due to increases in the ion conductivity of the fuel cell electrode, which is proportional to the volume fraction of ionomer in the electrode.55,56 On the other hand, the subsequent decrease in fuel cell performance as the ionomer composition further increases above the ‘opti­ mum loading’ indicates that another factor such as electronic resistance, oxygen permeability, or other electrode structural change comes into play. Increasing electronic resistance in the electrode may not be the major factor for the decreased fuel cell performance, as it has been observed that the electronic resis­ tances for cells having higher electrode ionomer loading were still small.53 The major factor impacting fuel cell performance in electrodes with high ionomer loading can be elucidated by fitting the polarization curve in the low overpotential and linear regions using a Tafel analysis:

Structural Effect Ionomer composition

Since a fuel cell electrode consists of multiple components (electrocatalyst, catalyst support, and electrode ionomer), the composition of each component is critical to fuel cell perfor­ mance and provides useful information about key properties of electrodes. The optimum composition of electrode ionomer is typically discussed with the fuel cell polarization curve. Figure 17 shows fuel cell polarization curves as a function of electrode ionomer composition. PFSA Nafion (Figure 17(a)) and HC-based sulfonated poly(arylene ether ketone) (S-PEEK) (Figure 17(b)) are used to illustrate the effect of polymer chemistry. The performance of PFSA-based electrodes is much superior to that of HC-based electrodes. The fuel cell perfor­ mance improves as the electrode ionomer loading reaches the ‘optimum loading’ and then starts to decrease. There are ‘opti­ mum loadings’ of 30 and 11 wt.% for Nafion and S-PEEK electrode ionomers, respectively, for carbon-supported plati­ num catalyst. A lower optimum binder loading 12 wt.% of Nafion was found for platinum black catalyst. In terms of volume, the optimum loading for each case corresponds to 40–50 vol.% ionomer. The higher density of platinum black catalyst as compared to Pt/C explains the lower Nafion loading required for platinum black catalyst, while the lower optimum

Eo ¼ Eeq − b log io

0.7 0.6

1.0 0.9

0.5 0.4 0.3

0.7

0.5 0.4 0.3 0.2

0.1

0.1

0.0 0.2

0.4 0.6 Current density (A cm–2)

0.8

1.0

4.9 wt.% S-PEEK 11.0 wt.% S-PEEK 19.4 wt.% S-PEEK 29.9 wt.% S-PEEK 30 wt.% Nafion

0.6

0.2

0.0

Cathode loading

0.8

10 wt.% 20 wt.% 30 wt.% 40 wt.% 50 wt.%

½10

where Eeq is the theoretical equilibrium potential (1.2 V vs. RHE), and i0 is the exchange current. Fitting the experimental data to the eqn [10] suggests that the ORR mechanism was not influenced by the ionomer load­ ing and the differences of exchange current density are insignificant. On the other hand, the ohmic resistance (obtained after subtracting the HFR) showed similar trends as

Cathode nafion loading

0.8

½9

where b is the Tafel slope, R is an approximately ohmic resis­ tance responsible for the linear region of the polarization curve, and E0 the constant given by

(b)

1.0 0.9

Cell potential (V)

E ¼ E0 − b log i − Ri

Cell potential (V)

(a)

705

0.0 0.00

0.02

0.04 0.06 0.08 Current density (A cm–2)

0.10

0.12

Figure 17 Effect of ionomer composition on fuel cell performance; (a) MEAs using PFSA Nafion (IEC = 0.91 meq. g−1) cathode binder measured at H2/O2 (ambient temperature and pressure),53 Reprinted with permission from J. Electrochem. Soc., 150, C745 (2003). Copyright 2003, The Electrochemical Society. (b) MEAs using HC-based S-PEEK cathode (IEC = 1.88 meq. g−1) binder measured at H2/O2 (35 °C, and ambient pressure).54 Reprinted with permission from J. Electrochem. Soc, 152, A752 (2005). Copyright 2005, The Electrochemical Society.

706

Polymers in Energy Applications | Polymers in Membrane Electrode Assemblies

the polarization curves,53 which suggests that the oxygen trans­ port resistance is likely the main factor limiting electrode performance at high electrode ionomer loading.

10.36.3.2.2

Ion exchange capacity

The IECW of the electrode ionomer impacts proton conductiv­ ity. An electrode ionomer with higher IECW provides more proton conduction. Although direct measurement of proton conductivity in electrodes has been measured by impedance techniques, it is generally complex and thus available only for a few electrode systems. The higher ionic conductivity of electro­ des with higher IECW ionomers has been witnessed by catalyst utilization measurements. Catalyst utilization is calculated as the electrochemically active surface area (ECSA) of the plati­ num catalyst divided by the total surface area of the platinum. The ECSA is typically estimated from the charge under the H-adsorption or desorption in cyclic voltammetry. Figure 18 shows the catalyst utilization as a function of IECW of the electrode ionomer. Catalyst utilization increases with IECW for both PFSA Nafion and HC-based S-PEEK electrode iono­ mers. This is attributed to the higher concentration of sulfonic acid groups and hence, better proton conductivity at the electrode. The IECW of the electrode ionomer also impacts on oxygen permeability. Oxygen permeability of the electrode ionomer is typically measured by chronoamperometry at microelectrodes with respect to temperature, gas pressure, and humidity.58–60 Both the hydrophobic and hydrophilic phases strongly impact the electrochemical O2 transport process, The hydrophilic phase is predominantly involved in the O2 diffusion pathway: hence, increased connectivity of hydrophilic domains and more water uptake enable a higher diffusion coefficient of oxygen in the electrode ionomer. The solubility of oxygen is determined mainly by the fraction of the hydrophobic compo­ nent in the membrane; a lower degree of hydrophobicity and less water uptake tend to enhance O2 solubility. Therefore, the O2 permeability, defined as the product of the O2 solubility

(a)

and diffusion coefficients, is affected by both the hydrophilic and hydrophobic components of polymer electrolytes. Table 7 compares the O2 solubility, diffusion coefficient, and permeability of polymer electrolytes measured by solid-state microelectrode cell.61–63 O2 solubility increases as IECW of the electrode ionomer decreases. This is because the hydrophobicity provided by the Teflon-like backbone structure increased as IECW increased. The diffusion coefficient of oxy­ gen, on the other hand, decreases as IECW decreases since hydrophilicity provided by the sulfonic acid side chain end group decreases. HC-based polymer electrolytes have signifi­ cantly lower oxygen permeability compared to PFSA electrode ionomer since the chain rigidity is higher and the hydrophobi­ city of the aromactic backbone of the HC-based ionomer is not as strong as for PFSA. Figure 19 compares the fuel cell performance of MEAs using Nafion ionomer. When comparing electrodes with a fixed iono­ mer loading of 32 wt.% (i.e., the optimum loading of Nafion EW 1100), slight performance improvements in fuel cell polariza­ tion curves were observed, suggesting that improved proton conductivity only marginally impacts ORR kinetics. When com­ paring the fuel cell performance at the optimum loading of Nafion electrode ionomer for each IECW, further improvements in fuel cell performance are observed for ionomers having higher IECW. The optimum loading of ionomer decreases as IECW increases (i.e., as EW decreases). This indicates that the required amount of protons is satisfied with a smaller percentage of higher IECW Nafion in the electrodes, and the resulting thinner electrode (or more porous) electrode further contributes to improved oxygen transport in the electrodes. This result shows that fuel cell performance is limited by oxygen transport rather than proton conduction, even when using Nafion ionomers, which have relatively good oxygen permeability. For HC-based electrodes, fuel cell performance is more lim­ ited by oxygen transport due to the lower oxygen permeability. Figure 20(a) shows the effect of IECW of HC-based electrode ionomers on fuel cell performance. The fuel cell performance of

(b)

80

85 S-PEEK

Nafion 80

Pt utilization

Pt utilization

70

60

75

70

50 65

40

60 0.9

1.0

1.1 IEC (meq g –1)

1.2

1.3

1.0

1.2

1.4 1.6 IEC (meq g–1)

1.8

2.0

Figure 18 Effect of IECW on platinum utilization. (a) Nafion electrode ionomer (28 wt.%) measured from H2/air fuel cells at 120 °C/35% RH, H2/air with 1 atm total pressure (graph is redrawn from Xu, H., et al. J. Electrochem. Soc., 2007, 154 (2), B271–B27857) (b) S-PEEK electrode ionomer (10 wt.%) measured from half cell using 1.0 M H2SO4 aqueous solution.54 Reprinted with permission from J. Electrochem. Soc., 2005, 152, A752. Copyright 2005, The Electrochemical Society.

Polymers in Energy Applications | Polymers in Membrane Electrode Assemblies Solubility, diffusion coefficients, and permeability of O2 at 50 °C and 3 atm61–63

Table 7 Polymer H2SO4 (1M) BAM 407 Aciplex

Nafion 117 Nafion 120 (5bar) SPES-60 SPES-50 SPES-40 SPES-30 PTFE

(a)

IECW (meq. g−1)

Solubility (mol cm− 3)  10 6

Diffusion coefficient (cm2 s−1)  10 6

Permeability (mol cm− 2 s)  10 12

10.20 2.46 1.12 1.0 0.91 0.91 0.83 2.2 1.8 1.5 1.2 0.0

5 1.75 8 12 11 9.4 27 3.5 4.3 5.9 10.5 37

31 40.6 12 6.9 8.0 5.5 1.8 10.3 9.7 3.5 2.2 0.35

155 71 96 82 88 51 49 36 42 21 23 13

(b)

1.0 Binder loading = 32 wt.%

0.9

1.0 0.9

0.8

0.8

0.7

0.7

Cell potential (V)

Cell potential (V)

707

0.6 0.5 0.4 0.3 1100 EW Nafion 920 EW Nafion 800 EW Nafion

0.2 0.1

0.6 0.5 0.4 0.3 1100 EW Nafion (32 wt.%) 920 EW Nafion (28 wt.%) 800 EW Nafion (25 wt.%)

0.2 0.1

0.0

0.0 0.0

0.2

0.4

0.6

0.8

1.0

1.2

1.4

1.6

Current density (A cm–2)

0.0

0.2

0.4

0.6

0.8

1.0

1.2

1.4

1.6

Current density (A cm–2)

Figure 19 Effect of IECW of Nafion electrode ionomer on fuel cell performance; (a) at a fixed binder loading and (b) at an optimum binder loading. Cell performance measured at H2/air with 1 atm total pressure at 80 °C/100% RH conditions.57 Reprinted with permission from J. Electrochem. Soc., 2007, 154, B271. Copyright 2007, The Electrochemical Society.

the MEAs with HC-based SPEKK-bonded electrode is much inferior to the MEAs containing Nafion-bonded electrode and even to the MEA without using an electrode ionomer. The slight deterioration in performance seen in the SPEKK with higher IECW is not expected since the catalyst utilization typically increases with IECW and better fuel cell performance typically achieved with high catalyst utilization. This apparent contradic­ tion was attributed either to increased contact resistance at the membrane–cathode interface arising from the interfacial immis­ cibility of the HC-based electrode ionomer and the Nafion membrane, discussed further in Section 10.36.3.3 or to a denser electrode structure and further limited oxygen transport with the high IECw ionomer (density of HC-based polymers increases with IECw). The oxygen permeability of the electrode using HC-based polymer can be improved by incorporation of hydro­ phobic PTFE. Figure 20(b) shows the electrode performance of PTFE-incorporated, S-PEEK-bonded electrodes as a function of IECW. The PTFE-containing electrodes greatly enhance the fuel cell performance. This was not observed in PFSA-bonded elec­ trodes, which usually showed inferior performance with PTFE

incorporation. Several reasons may explain the better perfor­ mance of the electrodes using HC-based polymer that also contain PTFE: 1. the presence of PTFE aids water management within the catalyst layer through increasing hydrophobicity helps effec­ tive water removal and thus prevents water flooding, discussed further in Section 10.36.3.2.3; 2. the presence of PTFE aids in creating a more porous struc­ ture (will be discussed in Section 10.36.3.2.4); and 3. the presence of PTFE aids to increase oxygen solubility in the electrode and thus increase oxygen permeability. Note that all of these attribute result in better oxygen transport in the electrode. Figure 20(b) shows that the fuel cell performance of GDEs containing PTFE increases proportionally with S-PEEK IEC. Table 8 shows the proton conductivity and oxygen permeabil­ ity of S-PEEK electrode ionomers used in Figure 20(b) as a function of IECW. The physical properties were measured using

708

Polymers in Energy Applications | Polymers in Membrane Electrode Assemblies

(a)

(b) 1.0

1.0 a: IEC = 1.4 SPEKK b: IEC = 2.1 SPEKK c: No ionomer in the cathode layer

0.8

Potential (V)

Potential (V)

0.8

PTFE + S-PEEK (1.88) PTFE + S-PEEK (1.55) PTFE + S-PEEK (1.02) S-PEEK (1.88)

0.6

0.4

0.2

0.6

0.4

0.2

0.0

0.0 0

50

100

150

200

250

Current density (mA cm–2)

0

100 200 300 Current density (mA cm–2)

400

Figure 20 Effect of IEC of HC-based electrode ionomers on fuel cell performance (a) S-PEKK electrode ionomer at a fixed binder loading of 10 wt.%. Cell performance measured at H2/air with 1 atm total pressure at 80 °C/75% RH conditions.64 Reprinted with permission from J. Electrochem. Soc., 2008, 155, B532. Copyright 2008, The Electrochemical Society. (b) S-PEEK electrode ionomer at a fixed binder loading (12 wt.%) applied to PTFE (wt.% 10 wt.%) gas diffusion electrode. Cell performance measured at H2/O2 with ambient pressure at 35 °C/100% RH conditions.54 Reprinted with permission from J. Electrochem. Soc., 2005, 152, A752. Copyright 2005, The Electrochemical Society.

Table 8 Proton conductivity and oxygen permeability of S-PEEK used in Figure 22(b)54 IECW (meq. g− 1)

Proton conductivitya (mS cm− 1)

Oxygen permeabilityb (mol cm− 2 s)  1012

1.02 1.55 1.88

5.2 24.4 38.1

5.0 6.6 8.7

a

Measured using an impedance analyzer in water under ambient conditions. Measured using Pt microelectrode. Permeability was obtained from diffusion limiting current, Id and the equation.Id= 4nFADbCbr, where A is the geometric area of the electrode, r is the radius of the microdisk electrode, F is Faraday’s constant, Db is the oxygen diffusion coefficient, and Cb is the oxygen solubility. b

thin PEMs; the actual conductivity and permeability of analo­ gous electrodes should be less. Both proton conductivity and oxygen permeability of the S-PEEKs increased with IEC. This suggests that the performance improvement observed with high IECW S-PEEK ionomers in PTFE-incorporated electrode (Figure 20(b)) may be caused by improved kinetics through high conductivity, as well as by improved mass transport through high oxygen permeability, where fuel cell performance was not significantly limited by mass transport.

10.36.3.2.3

Hydrophobicity

During fuel cell operation, the rate of water accumulation (due to the combined effects of generation via electrochemical reac­ tion and electro-osmotic migration due to the electrical potential gradient) and removal at the cathode may not be equal. This imbalance leads to either PEM dehydration or electrode flood­ ing, as illustrated in Figure 21. When the water removal rate is faster than the generation rate, liquid water in the electrode pores is depleted and oxygen transport is improved; however, the PEM attached to the electrode can suffer from dehydration

due to the lack of water supply from the electrode. When the water accumulation rate is faster than the removal rate, the electrode pores are filled with liquid water. As a result, oxygen transport is slowed down, while the PEM is well hydrated due to the contact with liquid water in the electrode. PEM dehydration and electrode flooding can be detected by hysteresis in the polarization curves when measured with increas­ ing versus decreasing current.65 Figure 22 shows the polarization curves of MEAs using a HC-based polymer-bonded cathode as a function of cathode humidification. In Figure 22(a) under bypass conditions (i.e., no cathode humidification), the polarization curve recorded with decreasing current shows lower voltage at higher currents than the polarization curve with increasing cur­ rent. This difference indicates that the water content in the PEM is still high as the current is decreased; suggesting that the water produced at higher current densities is not removed fast enough to dehydrate the PEM during the cell current measurement. The result is higher cell potential in the backward polarization curve, as evidenced by a lower HFR with decreasing current than with increasing current. The polarization curves under saturated water vapor at 60 °C (Figure 22(b)) shows the opposite trend: the polarization curve recorded with decreasing current shows higher voltage at higher currents than the polarization curve with increas­ ing current. The lower performance with decreasing current in spite of the lower HFR is an indication of electrode flooding. Under fully hydrated conditions (i.e., saturated water vapor at 90 °C), severe flooding occurs and significantly poorer perfor­ mance is observed throughout the entire current density range with very little hysteresis. Note that even open circuit voltage can be affected by the electrode flooding. In this severe case, the flooding cannot be readily detected by hysteresis behavior so it can only be discerned when the performance is compared with the cell under less-flooded conditions such as lower humidity or increased cell temperature/gas flow rate. Significant performance deterioration is noted when cathode flooding takes place and

709

Polymers in Energy Applications | Polymers in Membrane Electrode Assemblies

(a)

(b)

Carbon catalyst supporter

Pt electro-catalyst

Electrode binder

PEM

PEM Water phase

Figure 21 Schematic illustration of (a) PEM dehydration and (b) electrode after flooding.

(b)

1.1 Cathode humidification

1.0

Cell potnetial (V)

Cell potnetial (V)

(a)

bypass

0.9 0.8 0.7 0.6

Cathode humidification 60 °C 90 °C

1.0 0.9 0.8 0.7 0.6 0.5

HFR (Ω cm2)

HFR (Ω cm2)

0.5

1.1

0.4 0.3 0.2 0.1 0.0

0.4 0.3 0.2 0.1 0.0

0.0

0.2

0.4

0.6

0.8

1.0

1.2

1.4

Current density (A cm–2)

0.0

0.2

0.4

0.6

0.8

1.0

1.2

1.4

Current density (A cm–2)

Figure 22 Effect of external hydration on fuel cell performance using HC-based electrode ionomer; (a) no cathode humidification (b) saturated water vapor temperature of 60 and 90 °C. Fuel cell performance was measured under 80 °C, 20 psig back pressure, and high stoic conditions. Anode–cathode Pt catalyst loading: 3 mg cm−2; anode ionomer: Nafion 1100; membrane: Nafion 1135; cathode ionomer: partially fluorinated poly(arylene ether sulfone).

reduction of HFR usually cannot compensate for the performance loss. Oxygen transport limited by flooding is not correlated to the oxygen permeability of the electrode ionomer based on the fact that the oxygen permeability of the electrode ionomer increases but oxygen transport problem due to the flooding becomes more severe as RH increases. For example, oxygen permeability of an electrode ionomer material similar to that used in Figure 20 increased 1.2–5 mol cm−2S as RH increased from 50% to 95% at 50 °C under atmospheric pressure. This indicates that the oxy­ gen permeability of the electrode ionomer plays little role in electrode flooding. The major structural factor related to electrode flooding is the hydrophobicity of the electrode ionomer. If the electrode ionomer has a more hydrophobic nature, the hydro­ phobicity will generate larger capillary forces to drive liquid water across the porous structure. For this reason, electrodes using PFSA electrode ionomer tend to flood less than HC-based electrode ionomers.

Figure 23 shows the effect of electrode ionomer hydropho­ bicity on H2/air fuel cell performance under zero humidified and fully humidified cathode conditions. Hydrophobicity of the electrode ionomer increased in the order of wholly aro­ matic HC(Biphenyl) < hexafluoro bisphenol A (Hexafluoro) < decafluoro biphenyl (Decafluoro) < perfluroro (PFSA) back­ bone. Fuel cell performance is improved as fluorine content in the cathode ionomer increased. The performance difference between MEAs is greater under fully humidified conditions where electrode flooding for less-fluorinated systems is much greater and extends to low potential regions. Although other factors related to PFSA electrode structure contribute to its excellent electrode performance, this example shows the importance of the ionomer hydrophobicity to electrode flooding and fuel cell performance. The other important fac­ tors, which are, porous structure of electrode derived from dispersing solvent are discussed in the next section.

710

Polymers in Energy Applications | Polymers in Membrane Electrode Assemblies

CF2CF2 Perfluoro:

n

CFCF2 m OCF2CF

x

CF3 O(CF2)zSO3H

CF3 F Decafluoro:

F F

CF3

F

F

F F

O

Biphenyl :

S

O

O

HO3S

F

SO3H

1.0 No humidifed cathode

0.9

Fully humidified cathode

0.9 0.8

0.7

Perfluoro

0.6 0.5

Decafluoro

0.4 Biphenyl

Cell potential (V)

0.8 Cell potential (V)

SO3H

O

1.0

0.3

O

HO3S

SO3H O

O

S

O

C

O

Hexafluoro:

O

Perfluoro

0.7 0.6

Decafluoro

0.5 0.4 Biphenyl

0.3

Hexafluoro

Hexafluoro

0.2

0.2

0.1

0.1

0.0 0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1.0

0.0 0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1.0 Current density (A cm–2)

Current density (A cm–2)

Figure 23 Effect of fluorine content in the cathode ionomer on fuel cell performance; H2/air fuel cell performance using Nafion (PFSA) binder is also shown for comparative purpose. Fuel cell performance was measured under 80 °C, 20 psig back pressure, and high stoic conditions. Anode–cathode Pt catalyst loading: 0.2 mg cm−2; anode ionomer: Nafion 1100; membrane: Nafion 212.

Dispersing solvent

The dispersing solvent used to make the catalyst ink strongly influences the final morphology and structure of the fuel cell electrode largely through polymer / solvent interactions. Since the sulfonated polymer electrolytes used in the electrode con­ sist of hydrophilic and hydrophobic polymer segments (i.e., amphiphilic structure), the polymer electrolytes typically do not truly dissolve but instead disperse to form micelles. Figure 24 shows the solvent uptake of Nafion (EW = 1100) in different dispersing solvents. The results show that there are two distinct swelling envelopes corresponding to dual cohesive energy densities for Nafion: one is ascribed to the hydrophobic PTFE-like backbone (20 J1/2 cm−3/2) whereas the other is attributed to the hydrophilic side chains (35 J1/2 cm−3/2). Similar solvent swelling behavior was observed in HC-based sulfonated poly(ether ether ketone) copolymers, with a some­ what higher cohesive energy density (i.e., 26.1 J1/2 cm−3/2) for the hydrophobic part and a similar cohesive energy density (i.e., 35 J1/2 cm−3/2) for hydrophilic part.67 PFSA polymer electrolytes are typically more difficult to disperse compared to HC-based polymer electrolytes, due to the partial crystalline structure that hinders the ability to disperse Nafion. In the early 1980s, Grot, from DuPont, dis­ covered an efficient method to obtain Nafion dispersions.

Typically the procedure involves the dissolution of Nafion in a water–alcohol mixture at 250 °C under pressure in an autoclave for a few hours.68 A similar method was published by Martin et al.29 In the water–alcohol dispersing solvent mix­ ture, the solvents likely partition such that the water interacts

80 Percent increase in weight

10.36.3.2.4

60

40

20

0

15

20

25

30

35

40

45

Solubility parameter of solvent (J1/2 cm–3/2) Figure 24 Solvent uptake of Nafion (EW = 1100) vs. solubility para­ meters of solvents.66 Reprinted from Yeo, R.S. Polymer, 1980, 432, Copyright (1980), with permission from Elsevier.

Polymers in Energy Applications | Polymers in Membrane Electrode Assemblies

well with the sulfonic acid group and the alcohol solvates the hydrophobic backbone. Efforts to make PFSA dispersions in a single solvent have also been made. In the early 1980s, Covitch proposed a method for dissolving PFSA membranes under reflux in high boiling point solvents such as amides. This method, however, leads to only partial dissolution, and to degradation.69 The same method was used later with dimethylsulfoxide, 70 or hexamethylphosphotria­ mide 71 as solvents. Moore and Martin developed a method to add high boiling solvent to the water–alcohol solvent dispersion and displaced the water–alcohol mixture with the high­ boiling-point solvent.25,29 Another method to disperse PFSA ionomers in a single solvent is a re-dispersing technique. Aldebert and Pinery prepared a PFSA single-solvent dispersion by drying the water–alcohol PFSA dispersion at low temperature (T < 80 °C) and then re-dispersing the white powder the most of the polar solvents available, excepting water and formamide.72 Later, Curtin further developed the method by freeze-drying and re-dispersing in a single solvent including water, alcohol, dimethylformamide, and so on.73,74 Recently, Kim et al. devel­ oped a method to disperse PFSA ionomers using high boiling point alcohols such as glycerol or ethylene glycol. In their method, either the Na+ or the H+ forms of PFSA (EW < 1200) was dried at low temperature (i.e., < 60 °C) and dispersed directly in high boiling point alcohols such as either glycerol or ethylene glycol, below the solvent’s boiling temperature.75 The amphiphilic nature of the PFSA implies that its disper­ sion morphology can be controlled by differential interactions with various solvents. Thus, the various solvents mentioned above are likely to yield dispersions with different microstruc­ tures. When these dispersions are mixed with catalyst particles for the preparation of electrode films, the polymer–solvent interactions will continue to strongly affect the morphology of the catalyst suspension, as well as the final microstructure of the electrode film after the solvent is removed. Hence, con­ trol of dispersion morphology is likely impact fuel cell performance and durability. Therefore, understanding the dis­ persion morphology of PFSAs is of great importance to understanding and manipulating the fuel cell performance and durability of the electrode films. While numerous studies exist on the morphology of PEMs, relatively few have been conducted on the dispersion morphol­ ogy. Among the studies that have been done, a variety of characterization techniques have been employed, including small-angle X-ray scattering (SAXS) and small-angle neutron scattering (SANS), light scattering, viscosity, conductivity, elec­ tron spin resonance, and NMR spectroscopy. SAXS and SANS dominate the literature, perhaps due to their unique ability to yield size, shape, and compositional information on aggregate structure. Early applications of these techniques for PFSA dis­ persions focused on Nafion 117 (EW 1100) in polar solvents. Aldebert et al.76,77 found that the H+ and Li+ forms of this ionomer formed cylindrical particles in solvents such as water, ethanol, formamide, and N-methyl formamide. Although the scattering curves were complicated by a maxi­ mum in the scattering intensity, the authors calculated cylinder radii that ranged from 18 to 31 Å, based on a Guinier analysis for cylindrical particles and from geometrical argu­ ments for cylinders organized in a hexagonal array. Loppinet et al.78 expanded on these studies to examine the effects of counterion and added salt for Nafion 117 dispersed in a larger

711

number of solvents. For all dispersions studied, acylindrical particle morphology was found. In addition to the Guinier and geometrical analyses, a Porod analysis was utilized to calculate cylinder radii, with all methods yielding fairly consistent results in the range of 12–26 Å. Again, the scattering profiles displayed a maximum in the scattering intensity, along with a higher order oscillation. By analyzing these peaks as pseudointerference terms originating from long-range electrostatic repulsions between the cylindrical particles, the authors conclude that the cylindrical particles possess some local ordering, rather than an isotropic arrangement, within the dispersion. The addition of salt screens the electrostatic interactions between particles, thereby reducing the degree of local order; interestingly, the addition of salt did not appear to affect the shape or size of the Nafion particles. In addition to the ‘modelless’ analyses of small-angle scattering data utilized above, more precise models can be used to describe the scattering from particle dispersions. The small-angle scatter­ ing intensity can be written in terms of the contrast (Δρ) between scattering entities, for example, polymer and solvent, the form factor (P(Q)) that describes the shape and size of the particle, and the structure factor (S(Q)) that describes interparticle interactions IðQÞα Δ ρ2 〈PðQÞSðQÞ〉

½11

For a polymer solution or dispersion, the contrast is given as the difference between the scattering length densities (SLDs) of the polymer and solvent, where the SLD can be calculated from the chemical composition and mass density. Whether or not the scattering data are fit to a precise model of the form given in the eqn [11], contrast variation studies, in which the contrast between polymer and solvent is varied through the use of deuterated/protonated solvent mixtures, can be used to gain further insights into particle composition. For example, Loppinet et al.78 examined Nafion 117 (H+ form) dispersions in protonated/deuterated methanol mixtures. In this contrast variation study, the authors were able to measure the SLD of the particle and subsequently extract information on its com­ position. By considering the SLD expected from the polymer’s chemistry and mass density, the authors propose that the cylindrical Nafion particles consisted of Nafion backbone and side chains only, with no solvent penetration and no solvated ionic groups. Further, they calculate a mass density intermedi­ ate between those expected for amorphous and crystalline PTFE and suggest high density, crystalline packing of the perfluori­ nated backbone within the particle. For dilute dispersions or solutions in which no interparticle interactions occur, S(Q) = 1 and the scattering data can be fit to a form factor to gain information on the polymer/particle shape and size. Figure 25(a) gives examples of SANS data for dilute dispersions of Nafion 212 (EW 1000, Na+ form) in glycerol, ethylene glycol, and NMP. The data from both the glycerol and ethylene glycol dispersions can be fit well to the form factor for cylindrical particles,79 and these fits yield the cylinder radius (22 Å for both), cylinder length (150 and 370 Å, respectively), and cylinder SLD (4.0  10−6 and 2.5 x 10−6 Å2, respectively). By comparing the cylinder SLDs to those expected for Nafion 212 and the respective solvents, we can calculate the volume fraction of solvent within the cylindrical particles. For the glycerol dispersion, this volume fraction is quite small, 0.05, while it is considerably higher, 0.42, for the ethylene

Polymers in Energy Applications | Polymers in Membrane Electrode Assemblies

Scattering intensity (cm–1)

(a)

(b) 1000

10

Scattering intensity (cm–1)

712

1

0.1 Glycerol Ethylene glycol NMP 0.01

0.01

0.1

Q (Å–1)

100% 2-propanol 75% 2-propanol 67% 2-propanol 50% 2-propanol 25% 2-propanol 0% 2-propanol

100

10

1

0.1

0.01

0.01

0.1 Q (Å–1)

Figure 25 SANS characterization of Nafion 212 dispersion in (a) nonaqueous single dispersing solvent systems and (b) water–2-propanol binary solvent system. Nafion solid content: 2.5 wt.%

glycol dispersion. These examples demonstrate that polymer– solvent interactions can significantly influence the size and composition of dispersed PFSA particles. In contrast to the glycerol and ethylene glycol examples in Figure 25(a), the SANS data for Nafion 212 in N-methyl pyr­ rolidone (NMP) does not fit well to the form factor for a cylindrical particle, nor does it fit to form factors for any other simple shape (e.g., sphere, ellipsoid, etc.). Instead, the best fit to this scattering curve is the Debye form factor, which describes a linear polymer chain in a theta solvent.80 This fit yields a radius of gyration of 40 Å, which agrees well with a Guinier analysis for these data. Thus, these SANS results show the power of the small-angle scattering technique in distin­ guishing between PFSA solution morphologies, that is, particle dispersions versus true solutions. Small-angle scattering can also be used to yield insights into the more complicated PFSA dispersion morphologies expected in mixed solvent systems, for example, the commonly used water–alcohol mixtures. Figure 25(b) shows SANS data for Nafion 212 (Na+ form) in water–2-propanol mixtures of var­ ious ratios that range from 0 to 100 vol.% 2-propanol. In 100 vol.% 2-propanol, the large upturn in scattering intensity at low Q suggests large particles or aggregates; furthermore, the scattering intensity exhibits a Q−4 dependence at high Q, which is indicative of a smooth particle–solvent interface. As the 2-propanol concentration is lowered from 100 to 50 vol.%, the scattering at low Q decreases and the formation of a peak becomes evident. In addition, the magnitude of the slope at high Q decreases, indicating greater solvent penetration into the Nafion particles. As the 2-propanol concentration is further lowered to 0 vol.%, a new low-Q upturn becomes evident, the peak grows in intensity and shifts to lower Q, and the magni­ tude of the high-Q slope increases again. Taken together, these data show that the Nafion particle morphology is strongly affected by changes in the water–2-propanol ratio. Even with­ out performing any precise model fits similar to those of Figure 25(a), the changes in scattering observed between these samples suggest that the large particles formed in 100% 2-propanol become more solvated with the addition of up to

50 vol.% water and then perhaps less solvated as the water content is further increased. Qualitatively, this is in good agree­ ment with expectations based on solubility parameters. As discussed above, 2-propanol should interact more favorably with the PFSA backbone, while water may solvate the ionic groups and side chains. Figure 26 shows the SEM images of electrodes prepared from glycerol single solvent and water–2-propanol mixed sol­ vent systems, after the catalyst ink was painted on a decal substrate and dried at 140 °C. The electrodes prepared from the water–2-propanol mixed solvent system clearly has a more open structure with a few microcracks, while the electrode prepared from the glycerol single-solvent system has a denser and more homogeneous structure. Further characterization of the electrodes after hot pressing at 210 °C for 6 min revealed that the structural differences still remained, while mechanical integrity of each electrode was improved. The structural differ­ ence between these two solvent-cast electrodes also appeared in solvent-cast stand-alone membranes. PEMs cast from the aqu­ eous dispersing solvent were fragile, while PEMs cast from the nonaqueous single-solvent system showed improved mechan­ ical integrity. Oxygen transport probably benefits from a more open structure, while the mechanical properties may degrade and thereby hinder the sustainability of the electrode during dynamic fuel cell operations. Figure 27 shows the fuel cell performance of MEAs using Nafion electrode ionomers prepared from different dispersing solvents. The electrode cast from the aqueous solvent system shows improved mass transfer behavior at high current density compared to the electrode cast from aprotic NMP solvent. The electrode cast from glycerol shows further mass transfer limita­ tions at high current density. The mass transfer limit is also affected by the composition of water/alcohol ratio. The mass transfer limit of the electrodes cast from the water–2-propanol solvent system increased with 2-propanol content. These results are consistent with SANS data that show denser and less phase-separated dispersion morphologies in glycerol and NMP, while the size of swollen particles in water–2-propanol dispersion increases with 2-propanol content.

Polymers in Energy Applications | Polymers in Membrane Electrode Assemblies

(a)

713

(b)

5.00 μm

VP2784-07 10.0 kV 10.0 mm ×10.0k SE

5.00 μm

VP2784-03 10.0 kV 10.0 mm ×10.0 k SE

Figure 26 SEM images of electrodes prepared from (a) water–2-propanol binary dispersing solvent and (b) glycerol single dispersing solvent. Drying temperature: 140 °C.

Figure 28 shows the fuel cell performance of MEAs using a HC-based electrode ionomers prepared from different disper­ sing solvents. The dispersing solvent of the HC-based electrode ionomer impacts the fuel cell performance in the entire range of current density, while the dispersing solvent of Nafion iono­ mer mainly impacts the fuel cell performance in the high current density regions. This is attributed to poorer oxygen transport in HC-based electrode ionomers than in PFSA-based electrodes. Although the extent of performance degradation is different, similar trends are observed: electrodes prepared from water–2-propanol showed the best performance, electrodes prepared from an aprotic solvent such as Dimethylacetamide (DMAc) showed intermediate performance and electrodes pre­ pared from glycerol showed the lowest performance. Presumably, this is correlated with the porosity of the electrode structure: the water–2-propanol solvent created the most por­ ous structure, while glycerol produced less porous and denser (a)

structures. However, electrodes prepared from water– 2-propanol did not have good mechanical integrity after an elevated temperature treatment (possibly due to the lack of crystallinity or chain flexibility), resulting in a fast disintegra­ tion of electrode after only a few hours of operations. In contrast, MEAs using Nafion electrode ionomer prepared from glycerol dispersions showed good durability.

10.36.3.3

Interfacial compatibility between the PEM and the electrodes is one of the most important factors affecting fuel cell perfor­ mance. Most commercial MEAs have good interfacial compatibility, as they are fabricated with optimized processing and operating conditions and with a limited materials set. However, interfacial failure due to lack of interfacial compatibil­ ity has often been observed under transient operating conditions (b)

1.0 0.9

1.0 0.9

iR corrected cell potential (V)

iR corrected cell potential (V)

Membrane–Electrode Interface

0.8 0.7 0.6 0.5 Dispersing solvent 0.4

Water/2-propanol NMP Glycerol

0.3

0.8 0.7 0.6 0.5

Water-to-2-propanol ratio 1:3 1:1 2:3 3:1

0.4 0.3

0.2

0.2 0.0

0.2

0.4

0.6

0.8

1.0

1.2

Current density (A cm–2)

1.4

1.6

1.8

0.0

0.2

0.4

0.6

0.8

1.0

1.2

1.4

1.6

1.8

Current density (A cm–2)

Figure 27 H2/air fuel cell performance of MEAs using PFSA Nafion electrode ionomer as a function of (a) dispersion solvent and (b) water/2-propanol ratio; fuel cell performance was measured at 80 °C, under fully humidified conditions, 30 psig back pressure. Membrane: Nafion 212, Catalyst loading: 0.2 mg/cm2(Pt/c)

714

Polymers in Energy Applications | Polymers in Membrane Electrode Assemblies

1.0

0.30

0.9

Water/2-propanol (1:1) DMAc Glycerol

0.8 0.7 0.6 0.5 0.4 0.3 0.2 0.00

Y = 7.2 × 10–4 + 0.035

0.25

HFR (Ω cm2)

iR corrected cell potential (V)

Cathode dispersing solvent

0.20 0.15 0.10 0.05

0.05

0.10

0.15

0.20

0.25

0.00

0.30

–2)

0

50

Current density (A cm

Figure 28 H2/air fuel cell performance of MEAs using HC-based electrode ionomer (BPSH-35) as a function of dispersion solvent; fuel cell perfor­ mance was measured at 80 °C, under fully humidified conditions, 30 psig back pressure. Membrane: Nafion 212, Catalyst loading: 0.2 mg/cm2(Pt/c)

and with nonconventional fuel cell materials. Interfacial delami­ nation was observed with conventional platinum-based electrode particularly when harsh operating conditions were applied. Several researchers observed that the performance of cells processed at low temperatures degraded after freeze–thaw cycles that resulted in localized membrane–electrode delamina­ tion.43,81,82 In liquid feed DMFCs, the interfacial delamination between anode and membrane often limits the fuel cell perfor­ mance, probably due to the greater water activity and high metallic volume fraction of the Pt–Ru alloy catalyst,43 On the other hand, interfacial delamination in H2/air fuel cells is usually observed between the PEM and cathode. Interfacial failure/dela­ mination is also observed in MEAs that utilize alternative metal catalysts or PEMs.43,83,84 Wu observed substantial interfacial delamination when a Co/Fe carbon composite was used in the cathode catalyst layer after 450 h steady state fuel cell operation at 0.4 V.85 Many researchers 83,84,86–89 reported that inferior fuel cell performance was obtained from alternative PEMs due to interfacial incompatibility/delamination between the mem­ branes and PFSA-bonded electrodes. Initial efforts toward improving interfacial compatibility of PFSA-based MEAs focused on changing the counter-ion. Dynamic Mechanical Analyzer (DMA) and Thermo Gravimetric Analysis (TGA) results showed that as the size of the counter-ion increases, the temperature at which the storage modulus drops by several orders of magnitude decreases; further these organic counter-ions are stable in Nafion at temperatures in excess of 300 °C.90,91 Consequently, changing the acid or sodium form of Nafion to a bigger counter-ion such as tetrabutyl ammonium allows the polymer to flow at a higher processing temperature and enhances interfacial adhesion with­ out degradation. Figure 16(b) (in Section 10.36.3.1) shows the effect of counter ion on long-term stability of H2/air fuel cell performance.

10.36.3.3.1 electrode

Interfacial resistance between PEM and

The interfacial resistance between PEM and electrode can be used as a metric to quantify interfacial compatibility. Interfacial

100

150

200

250

300

350

Wet membrane thickness (μm) Figure 29 High frequency resistance (HFR) of the single cells using 1100 EW Nafion membranes as a function of membrane thickness at 80 °C under DMFC conditions.89 Reprinted with permission from J. Electrochem. Soc., 2007, 154, B739. Copyright 2007, The Electrochemical Society.

resistance can be obtained from a cell resistance plot as a function of membrane thickness under fully hydrated condi­ tions, where HFR increases due to PEM dehydration can be ruled out. Figure 29 shows an example for an MEA using a Nafion-bonded electrode. In the plot, a straight-line fit of the data is obtained, as PEM resistance is expected to vary propor­ tionally with thickness, while all other sources of resistance have been held constant. The y-intercept of the plot represents the nonmembrane resistance, and the difference between the observed nonmembrane resistance and the electronic resis­ tance represents the membrane–electrode interfacial resistance. Nonmembrane resistance and electronic resistance for the plot is 35 and 27 mΩ cm2, respectively, and interfacial resistance is 8 mΩ cm2. Table 9 shows the interfacial resistance of a few selected MEAs derived from different PEMs. Values of the interfacial resistance of the MEAs range from 8 mΩ cm2 for Nafion and 57 mΩ cm2 for recast Nafion. The two wholly aro­ matic membranes have interfacial resistances in between those of Nafion and recast Nafion. Another way to evaluate interfacial compatibility is by monitoring cell resistance under full hydration during extended-term fuel cell operations. While initial cell resistance

Table 9 Proton conductivity (σ), nonmembrane resistance, and interfacial resistance (RINT) of selected PEMs with Nafion-bonded electrode89

PEM Nafion Sulfonated poly(arylene ether nitrile) Sulfonated poly(arylene ether sulfone) Recast Nafion

σ (mS cm−1)

Nonmembrane resistance (Ω cm2)

RINT (mΩ cm2)

139 102

0.035 0.044

8 17

98

0.070

43

143

0.084

57

HFR gain (mΩ cm2),

after 200 h under DMFC operating conditions

Polymers in Energy Applications | Polymers in Membrane Electrode Assemblies

10.36.3.3.2

50

40

30

20

10

0 0

10

20

30

40

Initial interfacial resistance (mΩ

50

cm2)

Figure 30 Relationship between HFR gain after 200 hours under DMFC operating conditions and initial interfacial resistance.92 Reprinted with permission from J. Electrochem. Soc, 157, B1602 (2010). Copyright 2010, The Electrochemical Society.

reflects the initial state of conductivity of the fuel cell compo­ nents, changes in cell resistance during long-extended-term fuel cell operations (e.g., HFR gain) indicate interfacial failure, assuming that the other component’s conductivity is main­ tained. Hundreds of hours of fuel cell operation (often in accelerated conditions such as liquid feed DMFC) were typi­ cally applied to measure the HFR gain. While interfacial resistance measurements require at least three MEAs (often more than five MEAs for accuracy) to measure the nonmem­ brane and electronic resistances, the measurement of HFR gain during long-term fuel cell operations needs only one MEA, though for extended times. Also, the degradation of other MEA components should be negligible during the measure­ ment. Figure 30 shows an HFR gain versus initial interfacial resistance for several PEM and Nafion-bonded electrode sys­ tems. A linear correlation between initial interfacial resistance and HFR gain during long-term operation suggests that both methods may be used for quantifying interfacial compatibility between membrane and electrodes.

(a)

715

Origin of interfacial failure

There are three mechanisms known for interfacial failure depending on material characteristics (1) wetting/adhesion, (2) water transport mismatch between polymer electrolytes at the membrane–electrode interface, and (3) dimensional mis­ match between membrane and electrode, as illustrated in Figure 31. Interfacial failure through the wetting/adhesion mechanism assumes that physical interactions between the PEM and the electrode ionomer lead to interfacial issues. The physical inter­ actions of different polymers at the interface could result in poor adhesion and a tendency to delaminate. Alternatively, an inefficient interfacial contact between proton conducting domains, either at the atomic level or due to some sort of low conductivity skin, could result. Gubler et al. reported interfacial failure of radiation-grafted fluorinated ethylene propylene (FEP) and ethylene tetrafluoroethylene (ETFE)-based copoly­ mers using Nafion-based electrodes.83,84,94,95 They suggested that the tendency of these PEM materials to repel water caused difficulties in obtaining good membrane–electrode adhesion. The somewhat higher surface energy (10%) of radiation-grafted copolymers compared to Nafion was pro­ vided in support of this hypothesis. Several other PEM systems (with a wide range of surface energies) were also found to suffer from interfacial failure during fuel cell opera­ tion: PSSA84 sulfonated polysulfone,86 sulfonated polyimide,96 sulfonated polyphenylene,87 and composite membranes.88 Interfacial failure through water transport mismatch at the membrane–electrode interface is based on the fact that forces can build at the interface due to differences in electro-osmotic drag coefficient (ED), the number or water molecules carried per proton in a given polymer electrolyte, and diffusional/ hydraulic permeability limitation. These forces at the interface result because materials with ED mismatch want to either build up or deplete water at the interface depending on the direction of current flow. The ability or inability of water to move to or away from this interface can increase local forces and could play a role in membrane–electrode interfacial degradation. The ED of fully hydrated Nafion ranges between 2.6 and 6, depend­ ing on temperature and acid concentration83,84,86,87 and tends to be higher than other PEMs,88 including polystyrene sulfonic acid (ED = 2.0),96 sulfonated polyarylene (ether sulfone)s (ED = 1–3.5)89 and phosphoric acid doped membranes (ED < 1.0).81 Specific combinations of materials with very

(b)

(c) MEA

MEA

H2O

H2O

Dry state

MEA

Proton conducting channel

High drag polymer

Low drag polymer

Hydrated state

mechanical stress at the electrode– membrane interface may lead to delamination

Figure 31 Schematic illustration of proposed origin of interfacial failure; (a) poor wetting-adhesion, (b) water transport mismatch and (c) dimensional mismatch between membrane and electrode.93 Reprinted with permission from J. Electrochem. Soc., 2010, 157, B1616. Copyright 2010, The Electrochemical Society.

716

Polymers in Energy Applications | Polymers in Membrane Electrode Assemblies

Water uptake of cathode Water uptake of anode 60

HFR gain (mΩ cm2)

50 40 30 20 10 Sulfonated poly(arylene ether) Nafion Recast nafion

0 –10 –20

0

20

40

60 80 100 Water uptake (vol.%)

120

140

Figure 32 HFR gain after 100 h DMFC operations vs. water uptake of PEMs.93 Reprinted with permission from J. Electrochem. Soc., 2010, 157, B1616. Copyright 2010, The Electrochemical Society.

different ED could potentially lead to large forces at the inter­ face. As an analogy, microporous layers, and GDLs are almost always used for water management purposes within fuel cells in order to maintain membrane hydration and prevent flood­ ing. These hydrophobic materials can generate water pressures of several atmospheres under operating conditions. By reason­ able analogy, forces of a similar magnitude could arise at the membrane–electrode interface and result in interface-based performance losses. The third mechanism for interfacial failure presented is based on a dimensional mismatch between the PEM and electrode, where dimensional changes during hydra­ tion/dehydration result in membrane–electrode interfacial delamination that can subsequently lead to poor performance and durability. Yang et al. investigated adhesion between Nafion and Pt–Ru catalyst layers using microscratch tests. They suggested that interfacial fracture occurred in these sys­ tems by displacement of the membrane and catalyst layers since the PEM tends to swell but the catalyst layer tends to contract due to the coalescence of anode catalyst particles.82 Figure 32 shows the HFR gain versus water uptake for the PEMs. A strong correlation between cell HFR gain and water uptake was obtained with minimal data scatter and no extreme outlying data points. This correlation, and the lack of correla­ tion with other hypothesized parameters, supports the idea that interfacial failure primarily results from issues involving water uptake for the systems presented.

10.36.3.3.3

Methods to improve interfacial compatibility

There have been efforts to mitigate interfacial delamination. Wilson found that a thin-film electrode using the tetrabutylam­ monium (TBA) salt form of PFSA and processed at a higher pressing temperature (200 °C) maintained the fuel cell per­ formance longer than the same electrode processed at a lower pressing temperature (150 °C), as discussed above. Interestingly, the MEAs with high-pressing temperature showed stable performance during several freeze–thaw cycles, which is consistent with other results showing that stable interfaces of high-temperature-processed MEAs were maintained after 100 freeze–thaw cycles from -40 to 80 °C.97 According to Wilson,

the improved long-term durability for the highertemperature-processed PFSA electrode was attributed to the increased flow of the TBA+ form PFSA above its melting tem­ perature, which rendered a better electrode structure and better adhesion to the membrane. Xiang demonstrated improved interfacial adhesion with a modified thin-film method in which catalyst ink was applied directly onto each side of the pretreated membrane, followed by heating under vacuum at 80 °C. They compared this system with a lowtemperature-processed (130 °C) thin-film MEA. SEM images of the MEA formed by conventional low temperature pressing (130 °C) clearly showed interfacial delamination, while the MEA using the modified thin film method showed intimate membrane–electrode contacts.51 Narayanan and Surampudi introduced novel MEA fabrication method that involved soak­ ing the membrane and electrode with water and/or isopropanol before hot pressing to give better electrode transfer and interfacial adhesion. They observed better electrode trans­ fer and improved catalyst utilization for DMFC applications.98 However, the pressing process must be carried out at a tem­ perature of > 140 °C and no long-term data were provided. The second approach for improving interfacial adhesion between membrane and electrode is to use a buffer layer between the membrane and electrode. Choi et al. reported that interfacial delamination in DMFC mode could be pre­ vented by inserting thin electron-conducting polymers such as poly(N-vinyl carbazole) between the membrane and elec­ trode.99 Cho et al. improved interfacial adhesion of Nafion-based MEAs using a nonconductive thin layer coating of PVDF.100 Kim et al. also reported a similar approach, but used a low water-swollen ion conducting polymer such as sulfonated poly(arylene ether sulfone) between the membrane and electrode.75 This approach is preferred in DMFC applica­ tions, where methanol crossover can be further reduced; however, it may not be effective in PEMFC applications since the additional layer impedes proton transfer. The third approach involves using a solvent plasticizer. This method involves applying a volatile solvent on the catalyst layer before the hot-pressing step; the solvent acts to plasticize the PFSA ionomer and thereby improve interfacial adhesion with the membrane. The solvent plasticizer is evaporated dur­ ing the hot-pressing procedure. This method allows for the reduction of hot-pressing temperature of conventional Nafion-based MEAs down to 40 °C, depending on the solvent used, without cell performance loss and integrity. This method is applied to the non-Nafion PEMs with Nafion-bonded elec­ trodes without any interfacial failures, as confirmed by excellent fuel cell performance. All three approaches improve the interfacial adhesion to a certain degree, However, one should note that the enhancement of interfacial adhesion can vary depending not only on chemistry of PEMs, electrode, and buffer layers, but also on MEA processing and fuel cell operat­ ing conditions.

10.36.4 Summary This section discussed the structural effects of electrode iono­ mer on fuel cell performance and interfacial issues between PEM and electrode. Since electrochemical reactions in fuel cell electrodes are affected by many factors such as electrochemical

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717

activity, electronic conduction, ionic conduction, gas transport, mechanical stability, and membrane–electrode interfacial sta­ bility, several variables for electrode ionomer are selected to explain the effect of electrode ionomer on fuel cell perfor­ mance. The general structural effect of the electrode ionomer is summarized here.

systems. This section deals only with known electrode iono­ mers, to elucidate the relationship and advance the research progress for specific systems that promise better fuel cell performance.

1. Increased loading of electrode ionomer improves proton conduction, mechanical stability, and interfacial stability, but has detrimental effects on oxygen transport (through reducing porosity) and on catalytic activity (through redu­ cing electron conductivity). 2. Increasing the ion exchange capacity of the electrode iono­ mer improves proton conduction. However, mechanical stability and interfacial stability can be reduced due to the excessive water uptake. Oxygen transport of high IECW elec­ trode ionomers can be improved by electrode ionomer with a higher diffusion coefficient; however, the decreased hydro­ phobicity of high IEC binder may lead to a flooded electrode. 3. Increasing the hydrophobicity of the electrode ionomer can greatly improve oxygen transport within the electrode by increasing the capillary force to remove water. When a fluorinated PEM is used, increasing the hydrophobicity of the electrode ionomer can also improve interfacial stability by matching the water diffusion/electro-osmotic drag coef­ ficient between the PEM and electrode. 4. The dispersing solvent used for the electrode ionomer also significantly impacts electrode properties, most promi­ nently through affecting electrode morphology. Single dispersion solvent generates a denser and less porous struc­ tures compared to a water-based multisolvent systems. Consequently, MEAs using a single dispersion solvent have lower oxygen transport, while mechanical stability is rela­ tively good. Further control of electrode structure can be accomplished by changing the solvent composition of mul­ tisolvent systems or by changing the solubility parameter of single solvent system. 5. Chain mobility has a kinetic effect on fuel cell performance. An electrode ionomer having higher chain mobility tends to form favorable chain configurations during electrochemical reactions in a relatively short time, which is beneficial to fuel cell performance.

Acknowledgment

In general, PFSA electrode ionomers are superior to HC-based electrode ionomers largely due to their hydrophobicity, and resulting high oxygen transport. However, there are a few fuel cell applications in which PFSA electrode ionomers cannot be used or are difficult to utilize. For example, fluorinated poly­ mers in alkaline fuel cells are less stable than nonfluorinated polymers due to the electron-withdrawing nature of the fluor­ ine atom. Another example is high-temperature fuel cell applications (100–200 °C) where oxidative stability of fluori­ nated polymer electrolytes is questionable. Also, there are electrode systems that require HC-based electron-conducting polymers for catalytic stability. Therefore, it is critical to under­ stand the structure–property–performance relationship of electrode ionomers in order to develop advanced fuel cell

The authors at LANL thank the US Department of Energy Fuel Cell Technologies Program (Program Manger: Dr. Nancy Garland) for supporting this work (No. DE-AC52-06NA25396). MDG acknowledges partial support from the WCU program through the National Research Foundation of Korea funded by the Ministry of Education, Science and Technology (No. R31­ 2008-000-10092-0).

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Biographical Sketches Dr. Dae Sik Kim received his PhD in chemical engineering from Hanyang University (South Korea) in 2005, studying polymer membranes for PEMFC and DMFC. He continued working in the area of polymer electrolyte fuel cells at National Research Council in Canada (Dr. Michael D. Guiver’s group). In 2008, he joined the Los Alamos National Laboratory, where he expanded his research interest on alkaline fuel cell membranes as well as membrane electrode assembly of alkaline fuel cell. He has published over 35 papers in peer-reviewed journals and holds three US patents in the field of polymer electrolyte membrane fuel cells.

Cynthia Welch received her PhD in polymer science and engineering from the University of Massachusetts, Amherst, in 2001. She is currently a staff member in the Polymers & Coatings Group of the Materials Science & Technology Division of Los Alamos National Lab. She is an expert in polymer physics and in the use of small-angle scattering techniques. Recently, she has been applying these techniques to understand the morphological evolution of polymer morphology during the fabrication of nanocomposite electrodes for fuel cells.

Dr. Rex Hjelm is a senior scientist Los Alamos National Laboratory. He is also an adjunct professor of chemical and environmental engineering, University of California, Riverside. He has 145 articles and 146 invited presentations in scattering instrumentation, methods and computer code development, polymeric materials structure, materials defects, colloids, viral structure, structural biology of molecular motors, and the molecular response of polymers to stress. He has edited major reports in neutron scattering facility instrumentation and neutron scattering instrument design and edited three volumes on composite polymer materials. Hjelm has received six United States Department of Energy Defense Programs Award of Excellence in Research and the United States Department of Energy, Outstanding Mentor Award. Hjelm received a BSc in genetics from the University of California, Berkeley and a PhD in biology/biochemistry from the Johns Hopkins University. He was a Jane Coffin Childs and NATO Postdoctoral Fellow and Science Research Council Senior Fellow at Portsmouth University (UK), a Howard Hughes Fellow at Johns Hopkins in biophysics and medicine, a member of the faculty of the University of Illinois Schools of Medicine and Pharmacy, and scientist at Argonne National Laboratory divisions of Biology and the Intense Pulsed Neutron Source.

Dr. Yu Seung Kim is a staff scientist at Los Alamos National Laboratory. He received his MS and PhD degrees in chemical engineering from Korea Advanced Institute of Science and Technology, Daejon, Korea. He joined Professor James McGrath’s research group at Virginia Tech (1999) and co-invented a novel, hydrocarbon-based proton exchange membrane system (2003). He moved to the Los Alamos National Laboratory fuel cell team as a postdoc. research fellow (2003) and currently is a project leader for the LANL fuel cell program. He was recipient of Richard A. Glenn Award from the ACS Division of Fuel Chemistry in 2002 and LANL Distinguished Postdoctoral Award (Honorable mention) (2005). His present research is primarily focused on proton exchange membranes for PEMFC and DMFC, ionomer dispersion, advanced materials for alkaline fuel cell systems, and membrane–electrode interface.

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Dr. Michael Guiver is a principal research officer at the National Research Council Canada (NRC). He obtained his BSc from London University in 1977, MSc in chemistry (Carleton) in 1980, and PhD in polymer chemistry from Carleton University in 1988. He has worked in designing specialty polymer materials for membrane separations since1981. He is an editor for the J. Membr. Sci. and on the International Advisory Board of Macromolecular Research and the Editorial Board of Polymers. In 2009, he joined the WCU Department of Energy Engineering at Hanyang University, Seoul, Korea, teaching and conducting research 4 months per year. He is known for his published work in polysulfone modification by lithiation, and novel polymeric materials for membrane separations and fuel cell membranes. His present research is focused primarily on two areas: (1) polymer electrolyte membranes for fuel cell application (PEMFC and DMFC) and (2) the development of specialized intrinsically microporous polymers, which are high free-volume ladder polymers for membrane gas separation applications such as carbon dioxide (greenhouse gas) capture and oxygen enrichment from air.