hybrid inorganic-organic nanocomposite polymer electrolytes based on nafion and fluorinated tio2 for...

i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n en e r g y 3 7 ( 2 0 1 2 ) 6 1 6 9e6 1 8 1

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Hybrid inorganic-organic nanocomposite polymer electrolytesbased on Nafion and fluorinated TiO2 for PEMFCs

Vito Di Noto a,*,1, Mauro Bettiol b, Fabio Bassetto b, Nicola Boaretto a, Enrico Negro a,Sandra Lavina a, Federico Bertasi a

aDipartimento di Scienze Chimiche, Universita di Padova, Via Marzolo 1, Padova (PD), I-35131, ItalybBreton Research Centre, BRETON S.P.A., Via Garibaldi 27, Castello di Godego (TV), I-31030, Italy

a r t i c l e i n f o

Article history:

Received 2 March 2011

Received in revised form

20 July 2011

Accepted 29 July 2011

Available online 27 August 2011

Keywords:

Hybrid inorganic-organic proton-

conducting membranes

Nafion

Polymer electrolyte membrane fuel

cells

Dynamical mechanic analyses

Vibrational spectroscopy

Fabrication and testing of

membrane-electrode assemblies

* Corresponding author. Tel./fax: þ39 498275E-mail address: [email protected] (V. D

1 Active ACS, ECS and ISE member.

0360-3199/$ e see front matter Copyright ªdoi:10.1016/j.ijhydene.2011.07.131

a b s t r a c t

In this report, three hybrid inorganic-organic proton-conducting membranes based on

a novel fluorinated titania labeled TiO2F dispersed in Nafion were prepared. The mass

fraction of TiO2F nanofiller ranged between 0.05 and 0.15. The water uptake and the proton

exchange capacity of the membranes were determined; the membranes were further

characterized by TG, DMA and FT-IR ATR investigations. Finally, the hybrid membranes

were used in the fabrication of membrane-electrode assemblies (MEAs), which were tested

in operating conditions as a function of the back pressure and of the hydration degree of

the reagents streams. It was demonstrated that, with respect to pristine recast Nafion, at

25%RH the MEA fabricated with the membrane including a mass fraction of TiO2F equal to

0.10 yielded a higher maximum power density (0.206 W cm�2 vs. 0.121 W cm�2). Finally, it

was proposed a coherent structural model of this family of hybrid membranes accounting

for both the properties determined from “ex-situ” characterizations and for the perfor-

mance obtained from measurements in a single fuel cell in operating conditions.

Copyright ª 2011, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights

reserved.

1. Introduction ionomeric membrane between two gas diffusion electrodes,

Polymer electrolyte membrane fuel cells (PEMFCs) are a class

of electrochemical devices operating at temperatures below

130 �C characterized by a very high energy conversion effi-

ciency (as high as 55% or more), a high energy density and

a good compatibility with the environment [1e3].

The heart of a PEMFC is its membrane-electrode assembly

(MEA), a five-layer system obtained by sandwiching an

229.i Noto).

2011, Hydrogen Energy P

each covered by a suitable electrocatalytic layer to promote the

various electrochemical processes involved in the operation of

the device [4,5]. The ionomeric membrane used in the fabrica-

tionof theMEAisnecessary to transport theprotonsobtainedat

the anode from the oxidation of the fuel to the cathode, where

they are recombinedwith the products of the oxygen reduction

reaction (ORR). Water is obtained as the final reaction product

[3,6]. Despite significant research efforts, nowadays the most

ublications, LLC. Published by Elsevier Ltd. All rights reserved.

mailto:[email protected]

http://www.sciencedirect.com

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http://dx.doi.org/10.1016/j.ijhydene.2011.07.131



i n t e rn a t i o n a l j o u r n a l o f h y d r o g e n en e r g y 3 7 ( 2 0 1 2 ) 6 1 6 9e6 1 8 16170

widely used ionomeric proton-conducting membranes for

application in PEMFCs are based on perfluorinated polymers

such as Nafion�, Aquivion�, Aciplex� and others2 due to their

high proton conductivity, good mechanical properties and

excellent chemical and electrochemical stability [7e9].

However, these materials require a high hydration level to

transport protons effectively and their maximum operating

temperature is below 100 �C. These latter properties hinder the

technological application of these materials, since: (a) the final

fuel cell plant requires bulky and expensive water and heat

management modules; and (b) it is necessary to use pure

hydrogen as the fuel due to the poor tolerance of platinum

electrocatalysts toward even small traces of common contam-

inants. The latter (which include CO and H2S) are standard

reaction byproducts of the steam-reforming process typically

usedtoobtaincheaphydrogen fromcarbon,methaneandother

hydrocarbons [10,11]. The preparation of hybrid inorganic-

organic materials where a suitable filler is dispersed in a per-

fluorinated ionomer is one of the most promising routes to

address the drawbacks of the state-of-the-art proton-conduct-

ing membranes [12e24]. Several families of nanofillers have

beenexperimentedwith, including: (a)heteropolyacids, suchas

silicotungstic acid, phosphotungstic acid, molybdophosphoric

acid and others; (b) zirconium phosphate; (c) organically

modifiedsilicates andsilane-basedfillers; (d) zeolites; and (e) Pt,

PteSiO2 and PteTiO2. Fillers are generally chosen from mate-

rials with a strongly acid and/or hydrophilic character,

assuming that these featureswouldenhancewater retention in

the membrane at high temperatures and low hydration levels.

Furthermore, an acid filler is expected to provide additional

mobile protons as charge carriers. In the past few years, our

researchgrouphasexecutedanextensive studyof the interplay

between the structure and the proton conductivitymechanism

of hybrid inorganic-organic membranes obtained by doping

a Nafion hostmatrix with amass fraction of nanofiller equal to

0.15 or lower. Nanofillers were chosen among: (a) single inor-

ganic oxoclusters such as SiO2, TiO2, ZrO2, HfO2, WO3, Ta2O5

[9,25e28]; (b) single inorganic oxoclusters doped with an ionic

liquid [29]; (c) “coreeshell” oxoclusters suchas [(ZrO2)$(SiO2)0.67]

and [(TiO2)$(WO3)0.148] [30e32]; and (d) amorphous silica func-

tionalized with perfloroalkylated chains [33]. Significant inter-

actions between the inorganic nanofiller and theNafionmatrix

were developed. Indeed, the mechanical properties of the

hybrid materials were markedly improved at temperatures up

to ca. 200 �C; this evidence was ascribed to the formation of

dynamic crosslinks between the Nafion host polymer and the

surface of nanofiller particles [27,29e32]. In addition, in hybrid

membranes the presence of the nanofiller triggered extensive

modifications in the secondary structure of the Nafion host

polymer. In some systems, it was observed that a significant

fraction of the fluorocarbon backbone chains underwent a 157/ 103 conformational transition; hydrophilic domains were

alsoheavily influencedby theNafion-nanofiller interactions, as

witnessed by the FT-IR spectra of the hybrid membranes

[28,30e32]. Finally, the dielectric response of the hybrid

membranes was strongly affected by the addition of the

2 The DuPont Oval Logo, DuPont�, The miracles of science�and all products denoted with a � are trademarks or registeredtrademarks of DuPont or its affiliates.

nanofiller, with the development of dielectric polarization

events which were attributed to the formation of additional

interfaces between the various phases of the hybrid materials

[9,34,35]. Very surprisingly, it was observed that an improve-

ment in the proton conductivity of some hybrid membranes

over pristine Nafion is not necessarily obtained at the highest

values ofwater uptake [33], and that this improvement canalso

be triggered by basic nanofillers such as HfO2 [27,28] or by

nanofillers showing both hydrophobic and hydrophilic func-

tionalities [33]. These evidences, together with accurate inves-

tigations carried out by broadband dielectric spectroscopy

studies [25,27,29], led us to hypothesize that the proton

conductivity of hybrid inorganic-organic systems based on

perfluorinated ionomers such as Nafion may be improved in

those systems where the polymerenanofiller interactions

promote the coupling between the relaxation modes of the

hydrophilic and the hydrophobic domains [9]. In a continuing

effort to improve theunderstandingof theeffectofafilleronthe

structure and the proton conductivity mechanism of hybrid

inorganic-organic systems based on perfluorinated ionomers

such as Nafion, in this study a novel nanofiller made of fluori-

nated titania labeled TiO2F and obtained with a proprietary

procedure was adopted [36,37]. Three membranes were

prepared by a solvent-casting procedure varying the mass

fraction of the nanofiller between 0.05 and 0.15. The

membranes were characterized by TGA, DMA and FT-IR ATR;

their water uptake and proton exchange capacity (PEC) were

measured. In addition, the membranes were used in the fabri-

cation of membrane-electrode assemblies (MEAs), which were

tested in operating conditions. One of the major goals of this

workwas to devise a coherentmodel capable to account for the

structural featuresof thehybridmembranesand their interplay

with the proton conduction mechanism, allowing to interpret

the performance of theMEAs in single fuel cell tests in a variety

of operating conditions, with a particular reference to the back

pressure and the hydration degree of the reagents streams.

2. Experimental

2.1. Reagents

A 5 wt% solution of a Nafion� ionomer with proton exchange

capacity (PEC)of1mequiv$g�1 (Liquion1000EW, IonPower)was

used as received. Submicrometric fluorinated titania, labeled

TiO2F, was provided as a courtesy by Breton S.p.A. and was

synthesized according to a proprietary procedure [36]. All the

other reagents and solvents were provided by SigmaeAldrich

and further purified by standard methods [38]. The C2-20

electrocatalyst used in the preparation of all the MEAs had

aplatinumloadingequal to20wt%,waspurchasedbyBASFand

used as received. Doubly distilled water was used in all the

procedures.

2.2. Membrane preparation

Hybrid inorganic-organic membranes of formula [Nafion/

(TiO2F)x] were prepared by a solvent-casting procedure as

follows [31,33]. The nanofiller mass fraction x was set equal to

0.05, 0.10 and 0.15. Furthermembraneswere preparedwithout



i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n en e r g y 3 7 ( 2 0 1 2 ) 6 1 6 9e6 1 8 1 6171

adding the TiO2F nanofiller, labeled “pristine recast Nafion”

and used as the reference for all themeasurements. A suitable

amount of TiO2F was added to a suspension of Nafion in DMF

prepared as described elsewhere [31,33]. The total dry weight

of the Nafion ionomer and the TiO2F nanofiller was set equal

to 1 g. The resultingmixture was homogenized by a treatment

in ultrasonic bath for 2 h and recast in a Petri dish with

a diameter of 6 cm at 100 �C for 5 h. The resulting membranes

were characterized by a thickness of ca. 200 mm and were

further treated, purified and activated as described elsewhere

[33]. The membranes used in the fabrication of MEAs were

prepared with exactly half the amount of reagents, yielding

a thickness of ca. 100 mm. All the membranes were dried for

one night at room temperature under a dry air flow (l z 2)

before carrying out the thermogravimetric analyses, the

dynamic mechanical analyses, the FT-IR ATR investigations

and the fabrication of MEAs.

2.3. Fabrication of membrane-electrode assemblies(MEAs)

MEAs were prepared with a catalyst-coated substrate (CCS)

procedure as described elsewhere [39]. The platinum loading

in both the anodic and the cathodic electrocatalytic layer

was equal to 0.4 mg cm�1; the Nafion/C ratio was equal to 0.3

[4]. The electrocatalytic layers were deposited on GDS1120

carbon paper obtained by Ballard Power Systems. The

resulting gas diffusion electrodes (GDEs) were hot-pressed

on the membranes according to a protocol detailed else-

where [40].

2.4. Instruments and methods

Thermogravimetric analyses were performed with a High

Resolution TGA 2950 (TA Instruments) thermobalance. A

working N2 flux of 100 mL min�1 was used. The TG profiles

were collected in the temperature range between 20 and

900 �C, using an open platinum pan loaded with ca. 7 mg of

each material. Dynamic mechanical analyses (DMA) were

carried out with a TA Instruments DMA Q800 instrument,

using the film/fiber tension clamp. The temperature spectra

were measured by subjecting a rectangular dry film sample of

ca. 25 (height) mm � 6 (width) mm � 0.2 (thickness) mm to an

oscillatory sinusoidal tensile deformation at 1 Hz with an

amplitude of 4 mm, and with a 0.05 N preload force. The

measurements were carried out in the temperature range

from �100 to 210 �C at a rate of 4 �C min�1. The mechanical

response of the materials was analyzed in terms of the elastic

(storage) (E0) and viscous (loss) modulus (E00). tan d ¼ E00/E0 was

analyzed as a function of temperature to measure the mate-

rial damping characteristics such as vibration and sound

damping phenomena. The proton exchange capacity (PEC) of

the membranes was determined as follows [25,26]. 100 mg of

each sample was suspended in 100 mL of a 1 M KCl solution.

The suspension was stirred overnight; afterward, the solid

fraction was allowed to settle and the remaining aqueous acid

solution was titrated with a 10�3 M KOH solution using phe-

nophtalein as indicator. The water uptake of fully hydrated

sample films was determined by measuring the TG profiles of

the isothermal mass elimination vs. time as reported

elsewhere [30,31]. The number of moles of water per equiva-

lent of acid groups of Nafion l was determined from the

isothermal TG profiles (data not shown) using Eq. (1):

l ¼ 1000$

�wt0 �wtN

wtN$

1MWH2O$PEC$ð1� xÞ

�(1)

where wt0 and wtN are the weight of the fully hydrated and

the dry sample, respectively, MWH2O is the molecular weight

of water, PEC the proton exchange capacity of Nafion� and x is

the mass fraction of TiO2F. The water uptake was determined

as follows. First, water was eliminated from membrane

isothermally at 35 �C for 100 min. Second, the membrane was

kept at 130 �C for 40 min, in order to remove residual water.

The initial instant of the desorption process was determined

as reported elsewhere [30]. The FT-IR ATR spectra in the

medium infrared region (MIR) were collected with the

instrumentations and the methods described elsewhere

[25,26,30,31]. Briefly, the spectra were collected using a Nico-

let FT-IR Nexus spectrometer equipped with a triglycine

sulfate (TGS) detector at a resolution of 4 cm�1 and a Per-

kineElmer Frustrated Multiple Internal Reflections accessory

186-0174. FT-IR ATR measurements were obtained by aver-

aging 1000 scans. Each nanocomposite membrane was

squeezed between the surface of a prismatic germanium

crystal of 18 (height) � 51 (width) � 2 (thickness) mm3 and

a pressing counterpart device in order to achieve a perfect

contact. 50% ca. of the crystal surface was covered with the

membrane. An incident light angle of 45� with 25 total

internal reflections was adopted. The baseline correction of

FT-IR ATR profiles was carried out with a Nicolet FT-IR Nexus

spectrometer software.

2.5. Tests in a single-cell configuration

Single fuel cell tests were carried out in a 5 cm2 single cell

with a two-channel serpentine flow field for both the anodic

and the cathodic sides, using pure hydrogen as the fuel and

both pure oxygen and air as the oxidant. All the MEAs were

mounted in the single cell with their opaque side (see below

in Section 3.1.) facing the cathodic electrode. The tempera-

ture of the reagents streams and of the cell were kept

constant at 85 �C. The hydrogen flow rate was set equal to

800 mL min�1; air and oxygen flow rates were set to 1700 and

500 mL min�1, respectively. Polarization curves were first

collected with fully-humidified reagents streams at a back

pressure of 4 bar. Subsequently, polarization curves were

registered with fully-humidified reagents streams at a back

pressure of 1 bar. In a successive step, the relative humidity

of both reagents streams was lowered to 75%, and the cor-

responding polarization curves were measured after the

system reached stability. Similarly, further polarization

curves were determined with both reagents streams having

the same relative humidity, set equal to either 50%, 25%,

12.5% or 5%. All the measurements with reagents streams

having a relative humidity lower than 100% were collected at

a back pressure of 1 bar. Both pure oxygen and air were used

as the oxidants in every set of operating conditions. The

polarization curves were not corrected for internal resis-

tance losses.



Fig. 1 e (a) Dependence of water uptake (WU) and l vs.

nanofiller mass fraction x for pristine recast Nafion and

[Nafion/(TiO2F)x] membranes; (b) values of the measured

PEC as a function of x.

Fig. 2 e Thermogravimetric measurements of [Nafion/

(TiO2F)x] nanocomposite membranes; I, II and III indicate

the main thermal degradation events.


3. Results and discussion

3.1. Membrane preparation

The TiO2Fy nanofiller used in the preparation of the hybrid

membranes is based on TiO2 including ca. 2.3 wt% of fluorine,

which corresponds to an F/Ti molar ratio y equal to ca. 0.09.

The nanofiller was labeled TiO2F [36,37]. The concentration of

TiO2F particles on the bottom side of each hybrid membrane

after the solvent-casting procedure is higher, giving so rise to

the “nanofiller-rich” Side B; the upper side of the membrane

was labeled Side A. Side A and Side B are characterized by

different textures: Side A is smooth and glossy, while Side B is

opaque. The various characterizations here detailed are

carried out onmembranes sharing the same thickness, i.e. ca.

200 mm, as reported elsewhere [31,33]. On the other hand, all

the membranes used in MEA fabrication are ca. 100 mm thick,

in order to match more closely the usual standard for MEAs

fueled with hydrogen [8]. In addition, with a membrane

thickness equal to ca. 100 mm, good single fuel cell perfor-

mance is obtained, minimizing the crossover of the reagents

and the risk to develop fractures and pinholes in the final

MEAs.

3.2. Water uptake (WU), and proton exchange capacity(PEC)

Fig. 1 shows the evolution on x of: (a) the water uptake (WU);

(b) the number of water molecules per equivalent of acid

groups of Nafion l; and (c) the proton exchange capacity (PEC)

as a function of the nanofiller mass fraction x for pristine

recast Nafion and [Nafion/(TiO2F)x] membranes. The PEC

values determined as described in the Experimental Section

by titration procedures are very similar to the expected

nominal values. As x is raised, the PEC of the membranes

decreases as witnessed by Fig. 1(b). In addition it is observed

that, with respect to pristine recast Nafion, the WU and the l

values of [Nafion/(TiO2F)x] membranes are distinctively

smaller and reach a minimum at x ¼ 0.1. Both theWU and the

l values are decreasing following the same trend. This

evidence is not attributed entirely to the decreased PEC of the

hybrid membranes, but also to a reduction on x of the free

volume in the polar domains of the hybrid [Nafion/(TiO2F)x]

membranes, as will be better clarified in Section 3.7.

3.3. TGA analysis of the membranes

The TGA profiles reported in Fig. 2 highlight the thermal

decomposition processes typically observed in this class of

Nafion-based materials [26,29,31,33]. Three main events are

observed, labeled I, II and III. At the lowest temperatures

(T < 100 �C) an elimination is evidenced, which is ascribed to

the desorption of the residualwater from themembranes. The

initial water desorption is more pronounced in the pristine

recast Nafion, in accordance with the larger water uptake of

this membrane in comparison with the hybrid [Nafion/

(TiO2F)x] materials. I, II and III are attributed to the

degradation of sulfonic acid groups (100 �C< T< 250 �C), to the

degradation of the perfluoroetheral side chains of Nafion

(300 �C < T < 380 �C) and to the decomposition of the per-

fluorinated backbone chains of Nafion (T > 400 �C), respec-tively. The high temperature residue (Tz 900 �C) is consistentwith the weight percentage of the inorganic TiO2F nanofiller

component in the hybrid membranes. I and II are also studied

by evaluating the derivative of the TG profiles shown in Fig. 2

in the appropriate temperature ranges. Results for I and II are

shown in Fig. 3(a) and (b), respectively.With respect to pristine

recast Nafion, I and II of hybrid [Nafion/(TiO2F)x] materials



Fig. 3 e Derivatives of the TG profiles of [Nafion/(TiO2F)x]

membranes vs. temperature. (a) event I; (b) event II.


peak at a temperature which is ca. 40 �C lower (I: w170 �C of

hybrid materials vs. w210 �C of pristine recast Nafion; II:

w320 �C of hybrid materials vs. w355 �C of pristine recast

Nafion). This evidence is interpreted as follows: the TiO2F

nanofiller acts as a catalyst, promoting the I and II thermal

degradation events. Thus, it can be assumed that the TiO2F

Fig. 4 e Temperature spectra of storage (E0) and loss modulu

nanofiller is interacting with the eSO3H-tipped per-

fluoroethereal side chains of the Nafion host polymer. A

reduced thermal stability of the eSO3H groups in the hybrid

[Nafion/(TiO2F)x] materials in comparison with the pristine

recast Nafion may also explain the slightly lower measured

PEC in comparison with the nominal values. Indeed, with

respect to the pristine recast Nafion, in the hybrid [Nafion/

(TiO2F)x] materials stronger interactions between the eSO3H

groups and the particles of the TiO2F nanofiller are consoli-

dated during the thermal treatment at 130 �C, which are

responsible of the decrease of the PEC of these materials [31].

3.4. DMA analysis of the membranes

The DMA analysis is carried out in the temperature range

�100 �C< T< 210 �C. The trends of the storagemodulus E0 andof the loss modulus E00 as a function of the temperature are

plotted in Fig. 4. The typical Nafion behavior is observed [31],

characterized by: (a) a decrease in both E0 and E00 as T is raised;

and (b) a collapse of themechanical properties of both pristine

recast Nafion and [Nafion/(TiO2F)0.05] at T > 100 �C; on the

other hand, [Nafion/(TiO2F)0.10] and [Nafion/(TiO2F)0.15] main-

tain appreciable mechanical properties at temperatures as

high as 210 �C. Fig. 5 shows the trends of the elasticmodulus E0

of thematerials as a function of themass fraction of nanofiller

at T ¼ 25 �C and T ¼ 100 �C. With respect to the pristine recast

Nafion, the hybrid [Nafion/(TiO2F)x] membranes are charac-

terized by higher E0 values at both T ¼ 25 �C and T ¼ 100 �C. AtT ¼ 25 �C, E0 z 500 MPa for all the hybrid [Nafion/(TiO2F)x]

membranes, while the corresponding E0 value for pristine

recast Nafion is much lower, and equal to ca. 300 MPa. At

T ¼ 100 �C, E0 increases in the order: pristine recast

Nafionw [Nafion/(TiO2F)0.05] << [Nafion/(TiO2F)0.15] < [Nafion/

(TiO2F)0.10]. This evidence is attributed to the development of


TiO2F nanofiller in the hybrid membranes [27,29e32,41]. The

plots of the loss modulus E00 (right panel of Fig. 4) and of tand

s (E”) vs. temperature for [Nafion/(TiO2F)x] membranes.



Fig. 5 e Dependence of storage (E0) modulus vs. TiO2F

nanofiller mass fraction x at T [ 25 �C and T [ 100 �C.


(Fig. 6) highlight three main thermomechanical relaxations at

T z �50 �C, T z 50 �C and T z 100 �C. These three relaxation

types are typical of hybrid inorganic-organic proton-con-

ducting membranes based on Nafion, and are attributed as

follows [31]. The event observed at: (a) T z �50 �C is ascribed

to the thermomechanical relaxation of perfluoroethereal side

chains (b relaxation); (b) T z 50 �C is attributed to the

conformational transitions occurring in hydrophobic PTFE-

like domains of the Nafion host polymer (a relaxation of

fluorocarbon backbone chains corresponding to segmental

motions); and (c) T z 100 �C is assigned to the long-range

motion of both the backbone and the side chains which

results facilitated when a weakening of the electrostatic

interactions within the ionic aggregates occurs (aPC relaxation

peak). The inset of Fig. 6 shows clearly that the intensity of the

maximum of the aPC relaxation peak decreases as x is raised

from0 to 0.1. This is an indication that a smaller fraction of the

overall mechanical energy provided to the system is dispersed

by this relaxation mode as x is raised. In addition, the

temperature TPC of the maximum of the aPC relaxation peak

Fig. 6 e tan d vs. temperature of [Nafion/(TiO2F)x]

membranes. The inset shows the dependence on x of the

intensity of the maximum tan d and of TPC. TPC is the

temperature of the peak maximum. The lines in the inset

are intended as guides to the eye.

also decreases slightly as x is raised from 0 to 0.1. Both of these

evidences can be interpreted admitting that the density of


particles of the TiO2F nanofiller increases as x rises, thus

reducing the size of the hydrophobic domains of Nafion.

3.5. FT-IR ATR analysis

Fig. 7 shows the spectra of both sides of the [Nafion/(TiO2F)x]

membranes as a function of the mass fraction of the TiO2F

nanofiller. It is observed that the spectral profiles shown in

Fig. 7 are very similar in comparison to those reported else-

where for a similar class of hybrid inorganic-organic Nafion-

basedmaterials [31]. Thismakes the correlative assignment of

the FT-IR ATR spectra of the hybrid [Nafion/(TiO2F)x]

membranes very easy. It is observed that all the spectra of the

Side A of the membranes are similar to the spectrum of pris-

tine Nafion (left panel of Fig. 7). On the other hand, all the

spectra of the Side B of the [Nafion/(TiO2F)x] membranes are

quite different from the spectrum of the pristine recast Nafion

membrane (right panel of Fig. 7). This evidence is interpreted

admitting that during the solvent-casting process the TiO2F

nanofiller developed a concentration profile. With respect to

Side A, Side B became more enriched in the TiO2F nanofiller,

giving rise to an opaque texture. At the same time, the top side

of the membrane (Side A) remained smooth and gloss. The

study of the FT-IR ATR spectramust take into account that the

perfluorinated backbone chains of theNafionhost polymer are

present as a blend of two possible conformations, i.e., 103 and

157 [28,31,32]. In general, it is observed that the 157 confor-

mation predominates; the ratio between the two conforma-

tions is affected by the presence of a nanofiller, as discussed in

detail elsewhere [31]. Fig. 8 highlights the influence of the

TiO2F nanofiller on the helical conformation of the fluoro-

carbon backbone chains of the Nafion host polymer in the

hybrid [Nafion/(TiO2F)x] membranes. All the spectra and the

difference spectra reported in Fig. 8 arenormalizedon thepeak

at 980 cm�1, ascribed to the antisymmetric CeOeC stretching

of the perfluoroetheral side chains attached to a fluorocarbon

backbone chain characterized by the 157 helical configuration

[31]. Fig. 8(a) shows clearly that, with respect to pristine recast

Nafion, the concentration of the 103 fluorocarbon backbone

chains in the Side A of the hybridmembranes is larger. Indeed,

the intensity of the peaks at ca. 1420 and 1455 cm�1, ascribed to

the vibrational modes of the eSO3H groups belonging to

a Nafion macromolecule whose main fluorocarbon chain is in

the 103 configuration, is higher [31]. With respect to pristine

recast Nafion, the concentration of 103 fluorocarbon backbone

chains also increases on the Side B of the hybrid membranes,

as revealed by the difference spectra shown in Fig. 8(b). Indeed,

as the concentration of the TiO2F nanofiller is raised, the

relative intensity of the peaks ascribed to the 103 helical

configuration is significantly increased and reaches

amaximumfor the [Nafion/(TiO2F)0.10]membrane [31]. Thus, it

can be concluded that the presence of TiO2F nanofiller parti-

cles triggers a 157 / 103 conformational transition in a signif-

icant fraction of the fluorocarbon helices of the Nafion host

polymer. The final result is an enrichment of 103 fluorocarbon

chains and a reduction in the crystallinity of PTFE domains

throughout the [Nafion/(TiO2F)x] membranes.



Fig. 7 e FT-IR ATR spectra of pristine Nafion and [Nafion/(TiO2F)x] membranes. A is the upside surface of the membrane after

solvent-casting process; B is the bottom side.


3.6. Tests in a single-cell configuration

The performance in single-cell configuration of the MEAs

assembled with the hybrid [Nafion/(TiO2F)x] membranes is

shown in Fig. 9 and Fig. 10. In both instances: (a) the reagents

streams are fully humidified; and (b) pure oxygen is used as

the oxidant. The polarization and power curves shown in Figs.

9 and 10 are collected with reagents streams having a back

pressure of 4 and 1 bar, respectively. One way to gauge the

overall performance of MEAs is to determine the maximum

Fig. 8 e (a) FT-IR ATR spectra shown in Fig. 7 in the wavenumbe

of Nafion/(TiO2F)x] membranes; each spectrum is obtained by su

B reported in Fig. 7. All the spectra are normalized to the peak

power density they yield. Indeed, if MEAsmount conventional

gas diffusion electrodes (GDEs) including Pt/C electrocatalysts,

as in this work, in general the maximum of the power density

curve falls at a cell potential of 0.4e0.6 V. In these conditions,

mass transport issues are usually not important; thus, the

sources of overpotential arise from electrode kinetics and

ohmic losses. In a series of MEAs mounting the same GDEs

and having electrocatalytic layers sharing the same formula-

tion, the electrode kinetics and the ohmic losses arising from

the interfaces between the various layers of the MEA can be

r range between 1510 and 1390 cmL1; (b) difference spectra

btracting the FT-IR ATR spectrum of side A from that of side

at 980 cmL1.



Fig. 9 e (a) Polarization curves; and (b) power curves

obtained from MEAs assembled with a pristine recast

Nafion membrane (solid symbols) and the [Nafion/(TiO2F)x]

hybrid membranes (open symbols). Oxidant: pure oxygen;

back pressure of the reagents: 4 bar; RH [ 100%.

Fig. 10 e (a) Polarization curves; and (b) power curves

obtained from MEAs assembled with a pristine recast

Nafion membrane (solid symbols) and the [Nafion/(TiO2F)x]

hybrid membranes (open symbols). Oxidant: pure oxygen;

back pressure of the reagents: 1 bar; RH [ 100%.

Fig. 11 e Trends of themaxima of the power density curves

in MEAs as a function of the wt% of TiO2F nanofiller. Data

derived from Figs. 9 and 10. Oxidant: pure oxygen;

RH [ 100%. The lines are intended as guides to the eye.


assumed to be the same. Thus, the main factor distinguishing

the performance of the various MEAs is the ohmic loss due to

the proton resistivity of the ionomeric membrane. In a first

approximation, the maximum of the power density curve is

directly proportional to the proton conductivity of the iono-

meric membrane. Fig. 11 shows the maximum power density

derived from the power curves plotted in Fig. 9(b) and 10(b),

obtained by the MEAs as a function of the wt% of the TiO2F

nanofiller. It is clearly noticed that the maximum power

density increases as the wt% of the TiO2F nanofiller is raised,

both at a back pressure of 4 bar and of 1 bar. Fig. 9 shows

clearly that, with respect to the MEA assembled with the

pristine recast Nafion membrane, the MEAs assembled with

the hybrid [Nafion/(TiO2F)x] membranes are characterized by

significantly improved mass transport properties. This

conclusion is based on the larger maximum current densities

observed in the latter case (2.3e2.5 A$cm�2 vs. ca. 1.4 A$cm�2

of the MEA assembled with pristine recast Nafion). This

evidence is also observed as the back pressure of the reagents

streams is set to 1 bar, as reported in Fig. 10. Figs. 12 and 13

show the envelope of the polarization curves at different RH

% values obtained from the MEAs assembled with the pristine

recast Nafion membrane and with the hybrid [Nafion/

(TiO2F)0.10] membrane, respectively. As a general trend, it is

observed that the slopes of the polarization curves at inter-

mediate cell potentials (V z 0.6e0.4 V) increase as the RH% of

the reagents streams is lowered. This evidence is attributed to

the progressive decrease in proton conductivity both in the

membrane and in the GDEs. This latter phenomenon arises



Fig. 12 e Envelope of: (a) the polarization curves; and (b) the

power curves of the MEA assembled with the pristine

recast Nafion membrane as a function of the relative

humidity of the reagents streams. Oxidant: pure oxygen;

back pressure of the reagents: 1 bar.

Fig. 13 e Envelope of: (a) the polarization curves; and (b) the

power curves of the MEA assembled with the [Nafion/

(TiO2F)0.10] membrane as a function of the relative

humidity of the reagents streams. Oxidant: pure oxygen;

back pressure of the reagents: 1 bar.

Fig. 14 e Trends of the maxima of power curves as

a function of the relative humidity of the reagents streams.

Back pressure of the reagents: 1 bar; oxidant: (a) pure

oxygen; (b) air. The lines are intended as guides to the eye.


from the dehydration of the Nafion ionomer used in the

formulation of the electrocatalytic layers and gives rise to

a more difficult transfer of the ionic species at the interfaces

between the membrane and the GDEs. Thus, the main overall

result of the progressive dehydration of the MEAs is the

increase in the ohmic losses. Since all the MEAs mount the

same GDEs, it can be assumed that the effect arising from the

dehydration of the Nafion ionomer included in the electro-

catalytic layers is the same in all the systems. As a conse-

quence, the main discriminating factor in the performance of

the different MEAs arises from the contribution of the

membrane. It is to be noticed that water is produced at the

cathode side of MEAs during operation; if the current density

is large enough, this water may be sufficient to partially re-

hydrate the system. This causes a sudden decrease in the

resistivity and, as a consequence, the appearance of “bumps”

in the polarization curve, such as in the case of the MEA

assembledwith the [Nafion/(TiO2F)0.10]membrane at a relative

humidity of the reagents streams equal to 12.5% (see Fig. 13).

Fig. 14 details the effect of the reduction of the hydration

degree of the reagents streams used to feed the MEAs on the

maximum of the power density curves. It is observed that,

with respect to the pristine recast Nafion membrane, the

hybrid [Nafion/(TiO2F)0.10] membrane is more tolerant to

dehydration, showing a higher maximum power density as it

is fed with reagents streams with a relative humidity lower

than 100%. In particular, at a relative humidity of the reagents

streams equal to 25%, themaxima of the power density curves

of the hybrid [Nafion/(TiO2F)0.10] membrane and of the pristine




recast Nafion membrane are 0.206 W cm�2 and 0.121 W cm�2,

respectively. On the other hand, the performance of the

hybrid [Nafion/(TiO2F)0.05] and [Nafion/(TiO2F)0.15] membranes

is about as severely affected by dehydration as the pristine

recast Nafion membrane. The discussion carried out so far is

centered on experiments performed using pure oxygen as the

oxidant, to prevent complications arising from “blanketing”

effects caused by the nitrogen included in air. However, all the

major trends observed are still evident if air is used as the

oxidant, as witnessed by Fig. 14(b) and the figures reported in

the Supplementary Information. The complete envelopes of

the polarization and power density curves of the MEAs

assembled with the hybrid [Nafion/(TiO2F)0.05] and [Nafion/

(TiO2F)0.15] membranes are also reported in Supplementary

Information.

3.7. Interplay between the structure and the protonconductivity

The experimental evidence highlighted with the techniques

discussed above allows to propose a coherent model of the

structure of the [Nafion/(TiO2F)x] membranes, and how it

relates to the proton conductivity. If the nanofiller mass

fraction x is 0.15, aggregation between the nanofiller particles

occurs; at lower values of x, no aggregation is expected. TiO2F

particles are assumed to interact with all the components of

the Nafion host polymer, i.e., the eSO3H groups, the per-

fluoroethereal side chains and the fluorocarbon backbone

chains. Indeed, TiO2F particles prompt the thermal degrada-

tion ofeSO3H groups and of perfluoroethereal side chains (see

Fig. 3); at the same time, they trigger a 157 / 103 conforma-

tional transition in the secondary structure of the per-

fluorinated backbone chains (see Figs. 7 and 8). The

interactions between the TiO2F nanofiller and the Nafion host

polymer are expected to give rise to dynamic crosslinks, as

witnessed by the improvedmechanical performance of hybrid

[Nafion/(TiO2F)x] membranes in comparison with the pristine

recast Nafion (see Figs. 4 and 5). It is assumed that dynamic

crosslinks are formed between the eSO3H groups of Nafion

and the particles of the TiO2F nanofiller, and arisemainly from

dipolar interactions. In summary, it is expected that in

[Nafion/(TiO2F)0.05] the concentration of dynamic crosslinks is

relatively small, providing only a minor reinforcing effect in

comparison with pristine recast Nafion at T ¼ 100 �C. In

[Nafion/(TiO2F)0.10], the higher concentration of the TiO2F

nanofiller and its good dispersion in the Nafion host polymer

give rise to the highest concentration of dynamic crosslinks

occurring between the two phases and, as a consequence, the

highest E0 at T ¼ 100 �C [41]. On the other hand, in [Nafion/

(TiO2F)0.15] the concentration of the TiO2F nanofiller is prob-

ably large enough to give rise to aggregations in the inorganic

phase resulting in a somewhat poorer interaction between the

Nafion host polymer and the nanofiller. This results in

a slightly inferior E0 value in comparison with [Nafion/

(TiO2F)0.10]. The formation of dynamic crosslinks in the hybrid

[Nafion/(TiO2F)x] membranes is coherent with the particular

features of the surface of TiO2F particles. Indeed, metal oxides

are well-known to develop dynamic crosslinks with the

Nafion host polymer [27,29e32]. Such an interpretation is also

coherent with previous results obtained with other hybrid

inorganic-organic systems [33]. As the overall concentration

of dynamic crosslinks increases, the mechanical properties of

the materials are enhanced [27,29e32,41]. Thus, with respect

to pristine recast Nafion, themechanical properties of [Nafion/

(TiO2F)x] membranes are improved, as shown in Figs. 4 and 5.

Indeed, the presence of a sufficiently large concentration of

dynamic crosslinks between the TiO2F nanofiller and the

Nafion host polymer in [Nafion/(TiO2F)0.10] and [Nafion/

(TiO2F)0.15] materials may explain the reason the latter mate-

rials show appreciable mechanical properties at temperatures

beyond 120 �C and as high as 210 �C. In addition, with respect

to [Nafion/(TiO2F)0.10], it can be expected that [Nafion/

(TiO2F)0.15] includes a smaller concentration of dynamic

crosslinks between the TiO2F nanofiller and the Nafion host

polymer due to the formation of aggregations of nanofiller

particles, which give rise to a partial segregation of the inor-

ganic phases. The inclusion of TiO2F particles in the hybrid

inorganic-organic materials is also expected to modify

significantly the secondary structure of the hydrophobic

domains of the Nafion host polymer. In particular, since the

TiO2F particles can interact with all the main components of

the Nafion host polymer, it is assumed that theymay promote

the coupling between the hydrophobic and the hydrophilic

domains, boosting the proton conductivity of thematerials [9].

In addition, it is proposed that these Nafion-TiO2F interactions

give rise to smaller hydrophobic domains. This hypothesis

would lead to a smaller free volume in the hydrophilic

domains, explaining the reduced water uptake of the [Nafion/

(TiO2F)x] membranes in comparison with pristine recast

Nafion (see Fig. 1(a)). The presence of inorganic aggregates in

[Nafion/(TiO2F)0.15] is expected to give rise to interfaces

between nanofiller particles acting as preferential proton

percolation pathways in fully-humidified conditions. This

would explain the improved fuel cell performance of the MEA

assembled with [Nafion/(TiO2F)0.15] (see Figs. 9e11). Figs. 7 and

8 witness a 157 / 103 conformational transition in the per-

fluorinated backbone of the Nafion host polymer, which is

consistent with the development of dynamic crosslinks

between the TiO2F nanofiller and the organic matrix. The 157conformation is characterized by perfluoroethereal side

chains always facing the same direction of the fluorocarbon

backbone helix; on the other hand, in the 103 conformation

the side chains are distributed all around the backbone helix

[42]. As a consequence, a high percentage of 157 fluorocarbon

chains increases the degree of crystallinity and the size of the

hydrophobic domains, while a high concentration of 103chains promotes the increase of disordered phases, which are

capable to promote the interaction with the nanofiller. Thus it

can be hypothesized that, with respect to pristine recast

Nafion, a larger concentration of 103 fluorocarbon chains in

the hybrid [Nafion/(TiO2F)x] membranes is an indication of

smaller hydrophobic domains and a low water uptake. In

addition, this model is consistent with the improved mass

transport properties observed in the performance of MEAs

assembled with [Nafion/(TiO2F)x] membranes in comparison

with the reference MEA based on pristine recast Nafion (see

Figs. 9 and 10). Taken together, with respect to pristine recast

Nafion, in the hybrid [Nafion/(TiO2F)0.10] membrane a smaller

free volume in the hydrophilic domains and the lack of

interfaces between nanofiller particles may explain the lower




water requirements to achieve a good proton conductivity.

This model is consistent with the enhanced performance in

single fuel cell tests even with reagents streams at a relative

humidity as low as 12.5%, as witnessed by Figs. 12e14.

4. Conclusions

In this work, three hybrid inorganic-organic proton-conduct-

ing membranes are prepared by a solvent-casting procedure

dispersing a fluorinated titania (TiO2F) obtained through

a proprietary procedure [36,37] in a Nafion matrix. The

resulting [Nafion/(TiO2F)x]membranes include amass fraction

x of the TiO2F nanofiller equal to either 0.05, 0.10 or 0.15. It is

shown that, with respect to the pristine recast Nafion used as

the reference, the [Nafion/(TiO2F)x] membranes are charac-

terized by a markedly lower water uptake which reaches

a minimum at x ¼ 0.10. TGA investigations show that the

TiO2F nanofiller prompts the thermal degradation of: (a) the

eSO3H groups; and (b) the perfluoroetheral side chains of the

Nafion host polymer. Indeed, with respect to the pristine

recast Nafion, the associated thermal decomposition events

are shifted to lower temperatures by ca. 40 �C. The presence in

hybrid membranes of Nafion e TiO2F dynamic crosslinks is

witnessed by a marked improvement in the mechanical

properties of the hybrid membranes in comparison with

pristine recast Nafion at T ¼ 25 �C. The density of R-

SO3H$$$TiO2F$$$HSO3-R dynamic crosslinks is expected to

contribute significantly to the improvement of themechanical

properties of [Nafion/(TiO2F)x] membranes in comparison

with pristine recast Nafion. At T ¼ 100 �C, It is observed a 10-

fold increase in the storage modulus E0 of [Nafion/(TiO2F)0.10]

over the reference (ca. 50 MPa vs. ca. 5 MPa); in addition,

[Nafion/(TiO2F)0.10] and [Nafion/(TiO2F)0.15] show appreciable

mechanical properties at temperatures as high as 210 �C,while the reference undergoes an irreversible elongation at

T z 120 �C. The interactions between the TiO2F nanofiller and

the Nafion host polymer are also witnessed by: (a) the marked

decrease in the tand relaxation peak observed at T z 100 �Cand ascribed to the development of dynamic crosslinks

between the Nafion host polymer and the particles of the

TiO2F nanofiller; and (b) the increased concentration of fluo-

rocarbon chains with a 103 helical conformation, as evidenced

from the FT-IR ATR spectra of both sides of the hybrid [Nafion/

(TiO2F)x] membranes. Since the TiO2F nanofiller interacts with

all themain chemical features of the Nafion host polymer, it is

expected to improve the coupling between the hydrophilic

and the hydrophobic domains, reducing the size of the latter

and boosting the proton conductivity of the hybrid materials.

However, at x ¼ 0.15 a partial aggregation of TiO2F particles

may occur, resulting in a decreased concentration of dynamic

crosslinks and in the formation of interfaces between the

nanofiller particles. The latter may act as preferential proton

percolation pathways in fully hydrated conditions, resulting

in an improved performance in single fuel cell tests. The

hybrid [Nafion/(TiO2F)x] membranes are used to fabricate

membrane-electrode assemblies (MEAs), which are tested in

operative conditions in single-cell configuration. The results

are consistent with the functional and structural model

proposed above. Indeed, it is observed that the best

performance in fully hydrated conditions is achieved by the

MEA assembled with the [Nafion/(TiO2F)0.15] membrane, with

a maximum power density equal to 0.625 W cm�2 vs.

0.429W cm�2 of the reference. However, as the hydration level

is reduced below 100%, the best fuel cell performance is

registered for the MEA fabricated with the [Nafion/(TiO2F)0.10]

membrane: indeed, as the relative humidity of the reagents

streams is set to 25%, a maximum power density of

0.206 W cm�2 is obtained, vs. 0.121 W cm�2 of the reference in

the same operating conditions. This result is consistent with

a material characterized by hydrophilic domains requiring

a lower amount of water to operate effectively in comparison

with the pristine recast Nafion reference. Taken together, the

proposed hybrid inorganic-organic [Nafion/(TiO2F)x] materials

are promising candidates for the development of proton-

conducting membranes for application in PEMFCs operating

with reduced hydration and at a higher temperature in

comparison with state-of-the-art materials.

Acknowledgments

Research was funded by the Italian MURST project PRIN2007,

“Passive direct methanol fuel cells: electrocatalysts for the

oxygen reduction reaction based on carbon nitride supports

and hybrid inorganic-organic membranes based on fluori-

nated ionomers and nanoparticles of mixed oxoclusters”. The

authors are grateful to BRETON S.p.A. (Castello di Godego,

Italy, www.breton.it), to have provided the TiO2F nanofiller.

The author N. B. would like to thank Texa S.p.A. for the Ph. D.

grant. The authorswould also like to extend theirmost sincere

thanks to the staff of the electronic workshop of the Depart-

ment of Chemical Sciences of the University of Padova for the

skillful technical assistance, provided byMr. Claudio Comaron

and Alberto Doimo, M.S.

Appendix. Supplementary information

Supplementary information associatedwith this article can be

found, in the online version, at doi:10.1016/j.ijhydene.2011.07.

131.

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