Phosphate Containing Solid Electrolyte for ITFC

8/3/2019 Phosphate Containing Solid Electrolyte for ITFC

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Ref.1. New intermediate temperature fuel cell with ultra-thin proton conductor electrolyte - J. Power Sources 152,

(2005), 200.

A new type fuel cell called as hydrogen membrane fuel cell (HMFC).A much thinner electrolyte can be easily realized because it is formed on a solid, non-

porous membrane.Another advantage of this approach is an ease of high density stacking because thephysical base of this fuel cell is a metal film, not ceramics as in the case of SOFCs.The HMFC can use only proton conductors as an electrolyte because the hydrogenmembrane layer only permeates hydrogen.The BaCe0.8Y0.2O3 perovskite was chosen for the electrolyte material. The electrolytelayer, which is 0.7 m in thickness, was coated on Pd film by pulse laser deposition. Testcells were operated from 400 to 600 C. The open circuit voltages were close to theoreticalvalue in all operating temperatures. The power density was 0.9 and 1.4 W cm2 at the

operating temperature of 400 and 600 C.

Schematic diagram of hydrogenmembrane fuel cell structure.

IV characteristics of single cell at varioustemperatures. Anode gas was moist H2 andcathode gas was moist air (both 40 Chumidified).

Temperature(C)

Anode gas Cathode gas MeasuredOCV (mV)

Theoreticalvoltage(mV)

440 H2 = 100% Air 1103 1140

440 H2 = 50% Air 1082 1120

440 H2 = 10% Air 1036 1073

530 H2 = 100% Air 1081 1120

610 H2 = 100% Air 1051 1100

Table . Open circuit voltage of single cell in various conditions


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Ref.2. Electrochemical analysis of hydrogen membrane fuel cells- J. Power Sources 185, (2008), 922.

Electrochemical analysis was conducted with respect to a hydrogen membrane fuel cell (HMFC) with SrZr0.8In0.2O3electrolyte.Most of the voltage loss derives from the cathode and the electrolyte, and a small amount of anode polarization was observed only inregions with high current density.The cathode polarization was nearly an order of magnitude lower than that of SOFCs.The conductivity of the film electrolyte was almost identical to that of the sinter at 600 C; however, it was several times as large at 400 C.

AC impedance spectra of test cell withelectrolyte thickness of 2 m at 400 C. (a)Effects of the bias current, (b) effects ofconcentration of H2 in anode gas, and (c)

effects of concentration of O2 in cathodegas.

VI characteristics of test cells

with various electrolytethickness at 400 C.

Electrolyte, anode and cathode polarization asa function of current density at 400 C inhumidified H2 and air: (a) d = 0.7 m; (b) d =

2.0 m; (c) d = 4.0 m; (d) d = 6.0 m.


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Electrolyte resistance as a functionof electrolyte thickness at 400 C.

VI characteristics of test cell withelectrolyte thickness of 4 m at varioustemperatures.

TEM micrograph revealed that the film electrolyteconsists mainly of long columnar crystals whichcan be related to the conductivity enhancementbelow 600 C.

Temp.(C)

ElectrolyteThickness (m)

MeasuredOCV (mV)

Theoreticalvoltage (mV)

400

0.7

980

1142

2.0 1020 1142

4.0 1100 1142

6.0 1110 1142

500 4.0 1070 1123

600 4.0 1048 1102

Table- Open-circuit voltage of singe cells in variousconditions (Anode gas: H2; cathode gas: air.)

TEM cross-section image and schematic model ofthe film electrolyte.

Most of the voltage loss of the HMFC is due to cathode activation polarizationand electrolyte overpotential, and small anode concentration polarization wasobserved only in regions with high current density. The cathode polarizationwas approximately an order of magnitude lower than that of SOFCs, anddisplayed strong dependence on the temperature. Although the conductivity ofthe film electrolyte was almost identical to that of the sinter at 600 C, it was

several times as large at 400 C. In addition, the conductivity and the thicknessof the film electrolyte displayed a non-proportional relation.


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Ref. 3. Efficient, Anhydrous Proton-Conducting Nanofilms of Y-Doped Zirconium Pyrophosphate at Intermediate

Temperatures-Adv. Mater. 20,2008, 2398.

Nanofilms of amorphous Y-doped zirconium phosphate (ZrYP-120 film annealed at 120 0C) and pyrophosphate(ZrYP-400 film annealed at 400 0C) by a novel approach (a combination of the layer-by-layer solgel process and

annealing). The obtained defect-free ZrYP-400 nanofilm shows high proton conductivity with ASR as low as0.085Vcm2 at 340 0C under anhydrous conditions.The conductivity of ZrYP-120 film depends on moisture content (conductive in moist air at 20 0C but non-conductivein dry air) but the conductivity of ZrYP-400 film is independent of moisture content. In ZrYP-120 film conduction maybe due to H3O+ or H5O2+ but in ZrYP-400 film it is transport of hydrogen-bonded protons through transformation ofintermolecular hydrogen bond

The effective proton concentrationis enhanced by Y-doping.

The impedance response of the Zr0.95Y0.05P3-400 pyrophosphate nanofilm at differenttemperature in the flow of dry air.


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The conductivity () value of the ZrP3-400 film increases from 1.6x10-7 to4.2x10-5 S cm-1, from 200 to 400 0C. With the doping of 5 mol % yttrium(Zr0.95Y0.05P3-400) increases from 7.4x10-7 to 1.5x10-4 S cm-1 from 200 to340 0C. At 340 0C, its proton conductivity is 12.5 times higher than that of

the ZrP3-400 nanofilm. Further increase in the doping (Zr0.90Y0.10P3-400),leads to lesser improvement of proton conductivity due to the formationof segregated yttrium phosphate leading to partial disconnection of theP2O7 network for proton conduction.

The turn around 280 0C in the curve of the Zr0.95Y0.05P3-400 film implies thatthere might be a change of conducting mechanisms or conducting species.The turn around is most probably ascribed to hydrogen-bond sites withhigher(OH), which require higher activation energy, that contribute to theproton conduction above 280 0C, while those with lower(OH) are mainproton conducting species below 280 0C.


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Ref. 4. Thickness-Induced Proton-Conductivity Transition in Amorphous Zirconium Phosphate Thin Films

Chem. Mater. 22, (2010), 5528.

Temperature dependency ofdry of ZPfilm with different thicknesses in dry air.

Plots ofdry at 150 C (O) and Eadry() of ZP

films and H

2

O

at 150 C () and EaH

2

O

() ofZP-H films.

amorphous ZrP2.5Ox thin films were a thermally stable, anhydrous proton conductor since these retain the abundant protons bound tometaphosphate glass framework. ZrP2.5Ox films reveal unique conductivity transition as followed by decrease of thickness, which is relatedto the water content inside film. The proton conductivity in dry air is not dependent on the thickness in d> 60 nm, but it increases asreducing thickness into below 60 nm. As a consequence, the conductivity of 40 nm thick film is 200 times higher than that of 60 nm thick

film. Furthermore, the conductivity of films with dof 100 nm increases by 2 orders of magnitude by hydration treatment and becomes ashigh as that of the as-prepared 40 nm thick film, though the conductivity of 40 nm thick film is kept at high level regardless of hydration.The elevated conductivity of thin films can attribute to the structural modification of metaphosphate glass framework by hydration. Thehydrated, high-conductive phase is thermally stable and does not restore to the original unhydrated phase by annealing at 400 C in dry air,but it is stabilized only as the thickness is less than a few tens of nanometers. This findings strongly suggest that the hydration ofmetaphosphate films is encouraged in the nanometer-thick layer in proximity to the electrode rather than in the surface layer exposeddirectly to moistures. Consequently, the highly conductive, hydrated phase is confined within the region of a few tens of nanometersthickness from electrode interface, so that the relatively thick films reveal the phase separation between the inner hydrated nanolayer and theouter unhydrated layer.

This is the first example of the proton-conductivity transition triggered by the size-confined hydration effect. The origin for this size-confinement is still unclear, and is not readily explained by the exponential decay of structure induced by the fixed charge or free energy atthe interface. These phenomena reported here are speculated to be operative in various phosphate-based amorphous proton conductors, sincethe Gibbs energy of amorphous compounds must be larger than that of the crystalline materials under same composition. The current resultspose an opportunity to create the high-proton-conductivity electrolyte based on the hydrated metaphosphate glass.


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Thermal Decomposition Spectra (TDS) of H2O (m/z = 18)of ZP films with thickness of 100 nm () and 40 nm (- -).

The amount of desorbed CH4 (m/z = 16), CO (m/z = 28), and CO2 (m/z = 44) ismuch smaller than that of desorbed H2O and OH in the measured temperaturerange the water does not evolve by combustion of organic contaminants.

There is no clear difference in the intensity of peak at around 400 C between the40 nm and the 100 nm the concentration of water at elevated temperature inthe 40 nm is larger than that in the 100 nm because the volume of the former is

much smaller than that of the latter.

physically adsorbed and/orchemically bound waters

proton stable atelevated temperatures

Temperature dependency of protonconductivity of 40 and 100 nm thick -ZrP2.5Oxfilms treated under various conditions.

The conductivity of 40 nm thick film is not varied by annealing in H2O/air forseveral hours. On the other hand, H2O of a 100 nm thick ZP-H film is nearly 2orders of magnitude larger than dry of the corresponding ZP film at

temperatures below ca. 200 C and EaH2O of the former is apparently smallerthan that of the latter. Above 200 C, ASR of ZR-H films is too small to detectprecisely with present setup.dry of 100 nm thick ZP-H/dry film is almost same as H2O of the ZP-H,indicating that the enhanced conductivity of ZP-H is not restored to theoriginal value of ZP even though the ZP-H was annealed again at 400 C indry air for overnight. These suggest that the relatively thick a-ZrP2.5Ox filmreacts with H2O at elevated temperatures to form the hydrated, high-conductive phase and the hydrated films do not recover to the original phase,

but the high value of of a thin film is maintained regardless of hydrationtreatment.The D2O of 100 nm thick ZP-H/D is lower than H2O of the ZP-H by a factorof 1.21.4, showing that the conductivity decreases by H/D isotope exchange.This result confirms that the equilibrium between film and moisture actuallytakes place although of the hydrated ZP-H is not sensitive to the humidity.This ratio of H2O / D2O is closer to the square root of MD/MH 1.4,suggesting that a-ZrP2.5Ox film dominantly conducted protons by thermallyactivated hopping process.


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Ref. 5. Thin Film Fuel Cell Based on Nanometer-Thick Membrane of Amorphous Zirconium Phosphate Electrolyte

- J. Electrochem. Soc. 158, (2011), B866.

Figure-(a) I-V characteristics of a fuelcell, H2 (pH2 = 0.5 atm), Ni@Pd|ZrP2.6Ox| Pt, air, at 400C. (b) Impedanceresponse of the fuel cell at OCV. Thesolid line indicates the calculated curvewith an equivalent circuit model depictedin (c). The semicircle drawn by reddashed line and blue dashed line is fitted

to the high-frequancy semicircle (HS)and low-frequency semicircle (LS),respectively. (c) Equivalent circuit model.R: resistance and Q: pseudo capacitance(constant phase element).

SEM images of the surface of Pd anode support (a) before and (b) after exposing to 50%-H2/Ar gas at 400 0C for 2 h. (c) Cross-sectional TEM image of HMFC of Ni@Pd/ZrP2.6Ox/Pt.

The inset is area-selected electron diffraction pattern from the ZrP2.6Ox film. The scale bar in(a) and (b) is 1 m, and that in (c) is 150 nm.

The hydrogen permeable membrane fuel cell (HMFC) based on amorphous ZrP 2.6Oxelectrolyte operates at 400C. The electrolyte reveals the protonic transport number of unitywithout electronic leakage in FC conditions with stable OCV of 1.0 V similar to thetheoretical value. The maximum power density is still limited and is about 2 mW cm2

because of the large polarization related to the proton transfer and the oxidative reaction ofmetal hydride at anode/electrolyte hetero interfaces. Ni interlayer was introduced betweenthe Pd anode and the ZrP2.6Ox electrolyte in order to suppress the deterioration of theelectrolyte nanofilm by the deformation of the Pd anode during hydrogen absorption. In theZrP2.6Ox electrolyte the transport number of proton was unity at 400C as determined by anEMF measurement. The modification of the Ni anode surface by an ultrathin Pt or Pd layereffectively decreased the anode/electrolyte interfacial polarization.


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Figure 3. I-V characteristics of the Ni/Pd|ZrP2.6Ox| Pt cells with various thicknessesof the Ni interlayer. TheI-V relationship isnot varied by changing the thickness,indicating that the low diffusivityof hydrogen through Ni ((i)) is notresponsible for the poralization resistancesof the HMFC.

Fig. 4.(a) Impedance responses at OCV ofH2, Ni@Pd |ZrP2.6Ox| Pt,pO2(Pt) fuel cell at400C with changing oxygen partialpressureat cathode,pO2(Pt), from 0.1 to 1.0.(b) Impedance responses at OCV of H2,Ni@Pd |ZrP2.6Ox| Pt, air fuel cellat 400Cwith changing hydrogen partial pressure atanode,pH2(Pd),from 0.2 to 1.0.

The resistances of the cell clearly changewith the gas concentrations at cathode andanode which indicate that the anodicprocesses significantly contribute to a largepolarization resistances of the cell.

Ni/PdPt@Ni/Pd

Pd@NiPd

Fig. 5. (a) I-V characteristics at 400C of the HMFC formed on anodes of Ni/Pd, Pt@Ni/Pd and

Pd@NiPd. (b) impedance responses of the cell at OCV. The HS and LS of the cell with an ultrathin

Ptanode surface layer (ca. 5 nm thickness), Pt@Ni/Pd |ZrP2.6Ox|Pt, become smaller than those of the cellwithout thePt surface layer and the maximum power density increases by40% even though OCV isslightly lower than 1.0 V.When an ultrathin Pd layer is implemented, the semicircles at the higherfrequency (HS) andlower frequency (LS) are effectively depressed and the maximum power densityofthe cell with the configuration of Pd@Ni/Pd |ZrP2.6Ox| Pt (1.8mW cm2) is twice higher than that ofNi/Pd |ZrP2.6Ox|Pt. These results strongly suggest that the electrochemical processes atthe anode,which can be assigned to the proton transferacross the hetero interface between the metal anode andtheoxide electrolyte and to the hydrogen dissociation reaction of theanode, are responsible for both

HS and LS.


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Fig. 6. Electromotive force of hydrogen concentration cell with aconfiguration of H2 (pH2 = 1), Ni@Pd |ZrP2.6Ox| Pt, pH2(Pt) = 0.10.8 atm.

The black solid line indicates the theoretical EMF given by Eq. (3). Themeasured (o) and the calculated slope (- - -) of EMF in wet condition. Themeasured () and the calculated slope () of EMF in dry condition.

These results suggest that the contribution of an oxide ion and an electron tothe charge carriers in the ZrP2.6Ox is negligibly small in comparison withthat of a proton. The large anodic polarization in our HMFC is not related to

the water evolution at the solid-solid hetero interface.


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Ref. 6. Proton-Conductivity of Amorphous Aluminum Phosphate Thin Films under Anhydrous Conditions - J.

Electrochem. Soc. 158, (2011), P41.

Fig. 2. Colecole plot of 110-nm thickAl1P2 films, in dry Ar (O). Solid line iscalculated with equivalent circuitmodel shown in inset. Rox = resistanceand Qox = capacitance of film bulk.

Fig. 3. Temperature dependency of of Al3P2 (d = 125 nm), Al1P1 (d= 105 nm), Al1P2 (d = 110 nm),and Al1P3 (d = 90 nm), measuredin dry Ar.

Al-rich films, Al3P2 and Al1P1, reveal the linear Arrhenius relationship in the whole temperature range with an activation energy Ea of 1.0 and0.9 eV, respectively. The values of Al1P1 are higher than those of Al3P2 by a factor of about 4.

On the other hand, P-rich films, Al1P2 and Al1P3, reveal nonlinear temperature dependency and exhibit the apparent change of slope at round200C. of the P-rich films linearly increases with an activation energy Ea of 0.7 eV below 200C with a break at around 200C and tends toincrease with an small Ea of 0.20.3 eV above 220C. The value of Al1P2 is very similar to that of Al1P3 in the whole temperature range.

Furthermore, of P-rich films is at least 2 orders of magnitude higher than that of Al-rich films at temperatures below 200C, but the latterbecomes close to the former at around 400C.

Summary: 100-nm thick films of amorphous aluminumphosphate were prepared in various Al/Si molar ratios bymultiple spin-coating with the mixed alkoxide solutions.

These were efficient proton conductor at temperaturesabove 200C under anhydrous conditions. The P-rich filmsconsisting of the aluminum metaphosphate glass phase keptthe relatively high conductivity of the order of 105 S cm1in the temperature region.

On the other hand, Al-rich films made of the Al2O3P2O5mixed glass phase could exhibit such high conductivityonly at around 400C due to the large activation energy.

The current results strongly suggest that the conductivity ofa-AlPnOx glassy films can be improved by forming theoptimal glass network structure.

Al3P2

Al1P1

Al1P2

Al1P3


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Fig 4 Time decay curve of theresistance of Al1P1 () and Al1P2 (X)films at 270C as switching the gasflow from 1%-H2/Ar to 1%-D2/Ar.

The resistances in Fig. 4 are normalized with resistance at 0 s in 1%-H2/Ar, R0. When theatmosphere is altered from 1%-H2/Ar to 1%-D2/Ar, both films reveal the abrupt increase ofresistivity.The resistivity change exponentially decays and saturates for less than 1 h in bothspecimens. Consequently, the ratio of resistance in dry 1%-D2/Ar to that in 1%-H2/Ar isabout 1.28 for Al1P1 and 1.22 for Al2P2. These values are relatively smaller than the

expected by classical diffusion model (MD/MH = 1.42).This may indicate that the protonic conduction in the films obeys nonclassical hoppingtransport where the dissociation of OH bond is the rate determining step for proton hopping.Hence, it is concluded that a-AlPnOx films predominantly transport the proton even undernonhumidified atmosphere.

Conduction in a-AlPnOx films

Local environment structure of the aluminum atoms

Fig. 5. Al K-edge XAS of a-AlPnOxfilms of 100 nm thickness. (a) AlPO4powder, (b) Y-zeolite powder, (c)

Al3P2, (d) Al1P1, (e) Al1P2, (f)Al1P3, and (g) -Al2O3 powder.

Al-rich films reveal only a strong peak at 1565.5 eV assigned totetrahedrally coordinated AlO4.P-rich films show the peak at 1565.5 eV with a shoulder at 1567eV and a clear peak at around 1572 eV, indicating that AlO4 andAlO6 coexist inside these films.The intensity of peaks at 1567 and 1572 eV in Al1P3 is larger thanthat in Al1P2, indicating that the molar ratio of AlO6 to AlO4increases with the increasing Al content.

Fig. 6. FTIR spectra of a-AlPnOx films and ofreferences of aluminumorthophosphate AlPO4 and

aluminum pyrophosphateAl(PO3)3.

The phosphate groups in a-AlPnOx mainly take the form of

metaphosphate in every composition, but the Al-rich filmsinclude phosphate group belonging to POAl chain built upby alternative linkage of PO4 and AlO4 tetrahedra via vertexsharing.

Al K-edge XAS indicates that most of the Al atoms in these filmstake tetrahedral configuration. The concentration of octahedralAlO6 coordinates clearly increases with the increasingconcentration of P, suggesting that the local structure of a-AlPnOx films tends to be close to the aluminum metaphosphateAl(PO3)3 where Al3+ ions coordinate to the oxygen ofmetaphosphate chain in octahedralconfiguration.


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Ref. 7. Proton conductors of cerium pyrophosphate for intermediate temperature fuel cell Electrochem. Acta 56,

2011, 6654.

Abstract The crystal structure & proton conductivity ofcerium pyrophosphate are investigated for potentialelectrolyte applications in intermediate temperature fuel cell.The CeP2O7 thin plate sintered at 450 C exhibits superior

proton conductivity (3.0 102 S cm1 at 180 C) underhumidified conditions .Doping with 10 mol% Mg on the Ce site of CeP 2O7, theconductivity was raised to 4.0 102 S cm1 at 200 C. TheMg doping also shifts and widens its temperature window forelectrolyte applications. Ce0.9Mg0.1P2O7 is considered a moreappropriate composition, with conductivity >102 S cm1between 160 and 280 C.A hydrogenair fuel cell with Ce0.9Mg0.1P2O7 electrolyte

generates electricity up to 122 mA cm2 at 0.33 V using 50%H2 at 240 C (Peak power 40 mWcm-1 at 240 OC).

Fig. (a) Arrhenius plots for proton

conductivity of sintered CeP2O7 atPH2O=0.114atm. (b) Arrhenius plots forproton conductivity of 450 C-sinteredCeP2O7 at PH2O=0.0040.114atm.

Fig. (a) Proton conductivities of theCe1xMgxP2O7, x = 0.0, 0.1, 0.2, and 0.3,

calcined at 300 C and sintered at 450 C(measured in moist air ofPH2O=0.114atm). (b) XRD pattern of themost conductive sample Ce0.9Mg0.1P2O7.

Fig.- Performance of the FC based on Ce0.9Mg0.1P2O7The peak power values are much lower than that of Sn0.9In0.1P2O7 cell, 264 mW cm2 at 250 C andSn0.95Al0.05P2O7 and SEPS polymer cell, 163 mW cm2 at 225 C (Hibino et.al.). The low power ofthis FC could result from the less-ideal cell potential. The origin of low cell potential can be the

porosity of sintered Ce0.9Mg0.1P2O7 disk (12%), which led to certain pores connecting the fueland air sides, such that hydrogen diffused through the electrolyte membrane.


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Ref. 8. Proton conductivity of CeP2O7 for intermediate temperature fuel cells Solid State Ionics 179,2008, 1138.

Abstract A single phase new proton conductive electrolyte ofCeP2O7 was synthesized.The conductivity of CeP2O7 increased with the increasing oftemperature and kept above 10 2 S/cm in the temperature range of150250 C.

A single H2/O2 fuel cell with CeP2O7 as electrolyte membraneexhibited a reasonable power density of 25 mW/cm2 at 200 C,showing potential applications.

Fig. 1. XRD patterns of CeP2O7calcined at different temperature.

Fig. 2. TGA-DSC of CeP2O7in air atmosphere (phasetransformation at 455 C fromCeP2O7 to Ce(PO3)4

Fig. 3. Electrical conductivities ofCeP2O7

Fig. 4. EMF value of (a) H2/O2concentration cell (b) H2 (1 atm)/H2 (0.1atm) concentration cell.

Fig. 5. H2/O2 fuel cellperformance with CeP2O7(1.2 mm) electrolyte.

The OCV value is lower than the theoretical value calculatedfrom Nernst's equation, which may be caused by the mechanicalleakage of gas through the electrolyte and the electron hole inthis material. The proton transfer number calculated base on the

EMF value to theoretical value of cell b was about 0.870.9,indicated that CeP2O7 is mainly a proton conductor.


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Ref. 9. Proton Conduction in In3+-Doped SnP2O7 at Intermediate Temperatures-J. Electrochem. Soc. 153(8), (2006),

A1604.

Proton conduction in undoped SnP2O7 -For galvanic cell H2 (1 atm),Pt/C|SnP2O7 |Pt/C, H2 + Ar(0.1 atm) the tH+ of SnP2O7 was in therange 0.970.99, indicating that this material is substantially a pure proton conductor in H 2 atmospheres. Cell H2,Pt/C|SnP2O7 |Pt/C, (air)showed EMF values deviating from the theoretical values (tH+ = 0.890.92), although the EMF values were as high as 920 mV. This may be dueto electron holes in the SnP2O7 bulk. However, mechanical leakage of gas through the electrolyte may also be responsible for the lower EMFvalues compared to the theoretical values, because the EMF value was affected by the thickness of the electrolyte used.

Fig.3 pO2 dependence ofconductivity of SnP2O7 at 250Cand pH2O of 0.0008, 0.0062, and0.12 atm.

AtpH2O = 0.0008 atm, conductivity increases with increasing pO2

under oxidizing conditions indicating that this material showsmixed electron-hole and proton conduction. This tendencydecreased with increasing pH2O & at pH2O = 0.12 atm, theconductivity was almost independent of pO2, indicating thedisappearance of electron holes from the bulk. It appears that thereis interaction between water vapor and electron holes. Thefollowing equilibrium has been proposed as a mechanism of protonincorporation in perovskite oxides such as SrCe0.95Yb0.05O3H

2O + 2h2H + 1/2 O

2

H2O + 2 Oox + 2h2OHo + 1/2 O2 (3)There is almost no variation of the conductivity with pO2 underreducing conditions pO2 = 10201027 atm was observed,excluding the reduction of Sn4+ and P5+ to lower valences.

Fig.4 Isotope effect onconductivity of SnP2O7.

Isotope effect on conductivity of SnP2O7- SnP2O7 yielded a 1.061.44 times higher conductivity and a lower activation energy of 0.03eV for H2O-containing atmospheres than for D2O-containing atmospheres. This result can be interpreted by a nonclassic H/D isotope effect.When the dissociation of the O-H bond is a rate-determining step for proton conduction, the activation energy for D+ is higher than that for H+

by a difference in zero point energy of 0.05 eV, which is near the difference in activation energy shown above. It is thus proposed that protonsmigrate via dissociation of O-H bonds hopping mechanism.

Abstract- SnP2O7 -based proton conductors were characterized. Undoped SnP2O7 showed conductivities 102 S cm1 in temperature range of75300C. The proton transport numbers (tH+) of this material at 250C to be 0.970.99 in humidified H2 and 0.890.92 under fuel cellconditions. Partial substitution of In3+ for Sn4+ led to increase in proton conductivity from 5.56 x102 to 1.95 x101 S cm1 at 250C. FTIR andTPD measurements revealed that the effects of doping on the proton conductivity could be attributed to an increase in the proton concentrationin the bulk Sn1xInxP2O7. The deficiency of P2O7 ions in the Sn1xInxP2O7 bulk decreased the proton conductivity by several orders ofmagnitude. The mechanism of proton incorporation and conduction is examined and discussed in detail.


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XRD patterns of Sn1xInxP2O7

Fig. 6 Temperature dependence of conductivityof Sn1xInxP2O7 in unhumidified air pH2O =

0.0075 atm.

Effect of In3+doping on proton conduction- maximum conductivity at In3+ = 10mol % which corresponded well with the substitution limit for In 3+ from the XRDmeasurements.

IR spectra The absorbance ratios of Sn0.9In0.1P2O7 toSnP2O7 for (OH, 1655 cm1) and (OH,3410 cm1 ) were 5.2 and 5.7 respectively,

which are comparable to their conductivityratio of 4.6 at 50C Fig. 6. This suggest thatthe absorption bands are mainly attributableto protons incorporated in the bulk. It is alsolikely that the protons interact with thelattice oxide ions to form hydrogen bonds.

TPD spectra

The proton concentrations in Sn0.9

In0.1

P2O

7

and SnP2O7 were determined by assumingthat all the evolved water vapor and H2 canbe attributed to the incorporated protons.The resulting proton concentration valueswere 10.4 and 2.5 mol % for Sn0.9In0.1P2O7and SnP2O7, respectively. These values werein good agreement with the protonconcentration predicted from the In3+ content

of 10 mol %. It thus appears that the protonswere fully introduced as point defects by thesubstitution of In3+ for Sn4+.


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Another possible reaction is the following interaction between water vapor and an oxygen vacancyH2O(g) + V 2Hi + Oxo H2O(g) + Oxo + V 2OHo (4)

The degree of Reaction (4) cannot be entirely demonstrated at this stage. However, when the proton conductivity of Sn0.9In0.1P2O7was measured at different pH2O values at 250C, it slightly increased with increasing pH2O. This result is associated with theprocess of proton incorporation through Reaction 4 rather than 3, because the order of mobility is oxygen vacancy < proton< electron hole. Therefore, Reaction 4 as well as Reaction 3 are possible mechanisms of proton incorporation.

Effects of P2O7 deficiency in Sn0.9In0.1P2O7 on proton conduction

Fig.10 Temperature dependence ofconductivity of Sn0.9In0.1(P2O7)1-y in

unhumidified airpH2O = 0.0075 atm.

TPD spectra

IR spectra

The proton conductivity ofSn0.9In0.1(P2O7)0.85 was about 2orders of magnitude lower than thatof Sn0.9In0.1P2O7, indicating that theconductivity is strongly affected bythe number of P2O74 ions in the

lattice.

The Sn0.9In0.1(P2O7)1-y IR spectrum showed peaks atalmost same wave numbers as those observed forSn0.9In0.1P2O7. The absorbance ratios of Sn0.9In0.1P2O7 toSn0.9In0.1(P2O7)1-y were 1.4 and 1.5 for(OH) and (OH),respectively, which are much smaller than theirconductivity ratio shown in Fig. 10. A similar behaviorwas obtained for the TPD spectra.

The proton concentration in Sn0.9In0.1(P2O7)0.85 wasestimated to be about 8.1 mol % per unit, which is notsignificantly different from the value of 10.4 mol % forSn0.9In0.1P2O7. Therefore, the large difference in protonconductivity between Sn0.9In0.1P2O7 andSn0.9In0.1(P2O7)0.85 can be attributed to the difference inproton mobilities rather than proton concentrationsbetween them.A possible speculation on the proton mobility of

Sn0.9In0.1(P2O7)0.85 is that the P2O7 deficiency causes apartial disconnection of the P2O7 network for protonconduction, resulting in a large energy barrier for proton

jumps between sites.


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Ref. 10. Intermediate-Temperature Proton Conduction in Al3+-Doped SnP2O7J. Electro. Soc. 154(12), 2007, B1265.

Abstract- Al3+-doped SnP2O7 proton conductors were prepared by controlling the initial composition of the reactants SnO2, AlOH3, and H3PO4.Sn1xAlxPyOzwithy


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The values for dc conductivity were approximately in agreement with the values for ac conductivity at all temperatures tested. The slightdifference in conductivity between the dc and ac measurements is probably due to the contact resistance between the electrode and the current

collector. The close agreement between the ac and dc conductivity values means that the high proton conductivity of Sn0.95Al0.05P2O7 is alsodemonstrated under fuel cell conditions.

The ionic transport numbers was in the range of 0.97 - 0.99, which was similar to the electrical conduction behavior of Sn0.9In0.1P2O7. Thissuggests that SnO2 contained in Sn0.95Al0.05P2O7 does not affect ion conduction. The migration of specific ions that contribute to the high ionictransport number of Sn0.95Al0.05P2O7 cannot be identified with this measurement, but it can be assumed that protons are major charge carriers inthis material because the ionic size of oxide anions or other cations is too large to achieve high conductivity under the present conditions.

Effect of Al3+

substitution on the conductivity ofSn1xAlxP2O7- In XRD patterns ofSn1xAlxP2O7 with x values of 0.10 and 0.15, somepeaks of AlPO4 were present. Therefore, conductivity ofthese samples was less.

Conductivity of thesample prepared from Al2O3 were lower than those of SnP2O7, indicatingthat the use of Al2O3 had a negative effect ontheconductivity.Also, there is a possibility that NO3- and Cl- ions remained in the samples prepared with Al(NO3)3 and AlCl3, causing a deterioration in theconductivity. The sample prepared with Al2O3 showed diffraction peaks assigned to AlPO4, besides SnP2O7 and SnO2 which led to decrease inconductivity as reported by Matsuda.


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Ref. 11. Intermediate temperature ionic conduction in Sn1xGaxP2O7 - J. Power Sources 195, (2010), 5596.

Abstract- A novel series of samples Sn1xGaxP2O7 (x = 0.00, 0.01, 0.03, 0.06, 0.09, 0.12, 0.15) are synthesized by solid state reaction. XRDpatterns indicate that the samples ofx= 0.00 0.09 exhibit a single cubic phase structure, and the doping limit of Ga3+ in Sn1xGaxP2O7 isx =0.09. The protonic and oxide-ionic conduction in Sn1xGaxP2O7 are investigated using some electrochemical methods at intermediatetemperatures (323523 K). The samples exhibit appreciable protonic conduction in hydrogen atmosphere, and a mixed conduction of oxide-ion and electron hole in dry oxygen-containing atmosphere. The highest conductivities are observed for the sample ofx = 0.09 to be 4.6

102

S cm1

in wet H2 and 2.9 102

S cm1

in dry air at 448 K, respectively. The H2/air fuel cell usingx = 0.09 as electrolyte (thickness: 1.45mm) generates a maximum power density of 19.2 mW cm 2 at 423 K and 22.1 mW cm2 at 448 K, respectively.

Fig. Temperature dependence of electricalconductivity of SnP2O7 and Sn1xGaxP2O7 (x = 0.01

0.12) in (a) dry air and (b) wet H2.

The higher conductivities of the Ga3+ doped samples are resulted from the higheroxygen vacancy concentration. While partially substituting Sn4+ with Ga3+ ions, chargecompensation is achieved through the formation of oxygen vacancies as indicated bythe Eq. (1).

Mixed conduction of oxygen ion and electron hole is resulted from Eq. (2) in dry

oxygen-containing atmosphere,

When water vapor is introduced, as shown in Eq. (3) and (4), the conduction of electronhole and oxygen vacancy decreases, at the same time, the protonic conduction appears.In wet H2, the protonic conduction may be further improved according to Eq. (3), (4)and (5), resulting in the prevailing protonic conduction in hydrogen atmosphere.

The influence of the doping level of Ga3+ at Sn4+ sites on the conductivities may be attributed to the effective concentration of oxygen vacancyin the samples. On one hand, the total concentration of oxygen vacancy Vo mainly increases with the increasing of the doping level of Ga3+.On the other hand, the concentrations of point defect pairs, GaSnV0, GaSnV0GaSn, and GaSnOH0, which resulted from the coulombicattraction among the point defects with opposite charges, also increase at the same time. Considering the opposite two factors above, theeffective concentrations of oxygen vacancy may reach its largest value at x = 0.09. Moreover, the impurity phase of SnO

2in the samples for x >

0.09 may be also responsible for the decrease in the conductivities. Therefore, these factors result in the highest conductivity at x = 0.09.


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Ref. 12. Ionic conduction in Sn1xScxP2O7 for intermediate temperature fuel cells J. Power Sources 196,2011, 683.

A novel series of samples Sn1xScxP2O7 (x = 0.03, 0.06, 0.09, 0.12, 0.15) were prepared. The doping limit of Sc3+ in SnP2O7 was atleastx = 0.09.The conductivities increased with various Sc3+ doping levels in the order: (x = 0.12) < (x = 0.03) < (x = 0.09) < (x = 0.06). Thehighest conductivity was observed to be 2.76 102 S cm1 for the sample ofx = 0.06 under wet H2 atmosphere at 473 K.The ionic transport numbers (tion = 0.950.99) were close to unity, and the relatively low electronic conductivity in wet hydrogenatmosphere. The maximum proton conductivity (2.24 102 S cm1) at 473 K is higher than that in BaCe0.85Y0.15O3 (1.04 102 Scm1) under wet H2 atmosphere at 873 K.The H2/air fuel cells using Sn1xScxP2O7 (x = 0.03, 0.06, 0.09) as electrolytes (thickness: 1.7 mm) generated the maximum powerdensities of 11.16 mW cm2 forx = 0.03, 25.02 mW cm2forx = 0.06 and 14.34 mW cm2 forx = 0.09 at 423 K, respectively.Sn1xScxP2O7 may be a promising solid electrolyte system for intermediate temperature fuel cells.

Fig. Temperature dependence ofelectrical conductivity ofSn1xScxP2O7 in wet H2 atmosphere at323523 K.

Fig. XRD patterns of theSn1xScxP2O7 (x = 0.03 0.12)samples.

ff f 3 d l l h d


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Fig. 5. Partial conductivities of

charged species in wet H2 atmosphere.

Fig. 6.IVP curves for a hydrogen/air fuelcell using Sn1xScxP2O7 (x = 0.03-0.09) aselectrolytes at 423 K. Electrolyte thickness:

1.7 mm.

Fig. 4. EMFs of the H2 concentration cell: H2, Pt|Sn1xScxP2O7 (x = 0.06)| Pt, H2Ar (pH2=1.01410Pa) andwater vapor concentration cell: H2 (pH2O=2.34310Pa), Pt|

Sn1xScxP2O7 (x = 0.06)| Pt, H2(pH2O=1.23410Pa).

Effect of Sc3+ doping level on the conductivities- The influence of Sc3+ doping on theconductivities may be attributed to the effective concentrtion of oxygen vacancy and impurityphase of Sc(PO3)3. The higher Sc

3+ doping level resulted in higher oxygen vacancy concentration,nevertheless, the concentrations of point defect pairs, ScSnVO, ScSnVOScSn, and ScOHOalso increase at the same time. Considering above two opposite factors, the effectiveconcentrations of oxygen vacancy may reach its maximum value at x = 0.06. As a result, thehighest conductivity 2.76 102 S cm1 was observed for the sample ofx = 0.06 under a wetH2atmosphere at 473 K. The conductivities of samples ofx > 0.06 decreased with the increasing

doping levels due to the formation of the secondary Sc(PO 3)3 phase and lowereffective oxygen vacancy concentrations. The effect of the formation of the pseudo-cubic 333superlattice on the conductivities is still unclear in present stage.

f 13 S i i i f S O S O i i i i


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Ref. 13. Synthesis and characterization of dense SnP2O7SnO2 composite ceramics as intermediate-temperature

proton conductors- J.Mater. Chem. 21, (2011), 663.

Abstract: Dense SnP2O7-SnO2 composite ceramics wereprepared by reacting a porous SnO2 substrate with an85% H3PO4 solution at elevated temperatures.Comparison of the observed EMF with the theoreticalvalue in two gas concentration cells demonstrated that

this composite ceramic is a pure ion conductor, whereinthe predominant ion species are protons. FT-IR andproton magic angle spinning NMR analyses revealed thatthe protons interacted with lattice oxide ions in theSnP2O7 layer to form hydrogen bonds. An H/D isotopeeffect suggested that proton conduction in this compositeceramic was based on a proton-hopping mechanism. Theproton conductivity is 102 S cm1 in the temperaturerange of 250600 C.

Fig. 1 Influence of temperature on the growth of anSnP2O7 layer, (a) XRD patterns at the ceramic surface; (b)XRD patterns in bulk; (c) relationship between the electrical

conductivity of the ceramic at 400 C and the intensity of themain peak (200) for SnP2O7 on ceramic surface and bulk.

Fig. 2 Influence of carbon content in mixtures of carbon and SnO2 powderduring the preparation of SnO2 substrate on the growth of SnP2O7 layer. (a)XRD patterns at the ceramic surface; (b) XRD patterns in bulk; (c) relationship

between the electrical conductivity of the ceramic at 400 C and the intensityof the main peak (200) for SnP2O7 ceramic surface and bulk.

The electrical conductivity as well as the growth of theSnP2O7 layer increased with increasing carbon content.Unfortunately, a crack-free SnO2 substrate with carboncontent above 8 wt% could not be prepared. Thus, theoptimal carbon content was concluded to be 8 wt%.


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Fig. 3 Temperature dependence of the electricalconductivity of the SnP2O7SnO2 compositeceramic with various thicknesses. For comparison,data for the SnO2 substrate are also included.

Cross-sectional SEM and EDX images of the SnO2 substrate (Fig. 4) & SnP2O7SnO2 composite ceramic (Fig 5).(a) SEM; (b) Sn element mapping; (c) P elementmapping.

Fig. 6 Hydrogen permeation properties of theSnP2O7SnO2 composite ceramic at various

temperatures. For comparison, data for theSnO2 substrate are also included.

Denseness and robustness of SnP2O7SnO2 composite ceramic

(Fig. 4) (Fig. 5)


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10 vol% H2, AuSnP2O7SnO2Au, 1 vol% H2

E = RT/2F ln(PH2(II)/PH2(I))

Fig. 7 EMF values of (a) hydrogen concentrationcell and (b) H2O vapor concentration cell using

the SnP2O7SnO2 composite ceramic as anelectrolyte at various temperatures.

3 vol% H2O, AuSnP2O7SnO2Au, 0.6 vol% H2O

E = RT/2F ln(PH2O(II)/PH2O(I))

H2O concentration in eachchamber 3 vol%.

H2 concentration in eachchamber 10 vol%.

Fig. 8 Temperature dependence of the conductivity of the SnP2O7SnO2

composite ceramic in various atmospheres. Atmospheric gases used werewet air (PH2O = PD2O = 0.026 atm) and dry air (PH2O = ca. 0.001 atm).

Proton conduction in SnP2O7SnO2 composite ceramic


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Proton environment in SnP2O7SnO2 composite ceramic

Fig. 9 TPD spectrum of the SnP2O7SnO2composite ceramic powder. Fig. 10 FTIR spectrum of the SnP2O7SnO2composite ceramic powder. Fig. 111

H MAS NMR spectrum of theSnP2O7SnO2 composite ceramic powder.

The protons are present in the SnP2O7 layer. The inserted protons can easily jump between adjacent oxide ions by a series ofbreaking and making of OH bonds.


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Ref. 14. Ionic conduction in undoped SnP2O7 at intermediate temperatures- Solid State Ionics 181, (2010), 1521.

Abstract- SnP2O7 is prepared with various initial molar ratios of phosphorus vs. metal ions, P ini/Sn, and different temperatures from 573 to923 K. The preparation conditions are optimized giving consideration to the influence of H3PO4 concentration, Pini/Sn molar ratio and heat-treating temperature. The ionic conduction behaviors indicate that the samples obtained from 85% H3PO4 and SnO2 nanopowders with Pini/Sn 2.4 are a single cubic phase, and that in wet hydrogen atmosphere, the samples are almost pure ionic conductors, the ionic conduction iscontributed mainly by proton and partially by oxide ion. An ionic conductivity of 2.17 10 2 S cm 1 is achieved for the sample prepared

from Pini/Sn = 2.8 under wet hydrogen atmosphere at 448 K.

Single cubic phase of samples is obtained when Pini/Sn =2.4, 2.6, 2.8 and 3.0, corresponding to Pfin/Sn = 2.02, 2.16,2.28 and 2.50, respectively. When Pfin/Sn > 2, excess P asamorphous PmOn existed in the grain boundary of SnP2O7crystal. When Pini/Sn = 2.0, except cubic structure ofSnP2O7 as a main phase, a SnO2 impurity phase also isobserved. Therefore, in order to obtain a single cubic

phase of SnP2O7, Pini/Sn should be controlled to be 2.4 ormore. However, moisture absorbability of the samplesincreases with increasing Pini/Sn molar ratio.

Fig. Temperature dependence of conductivity of SnP2O7 prepared fromdifferent Pini/Sn molar ratios and heat-treating temperatures.

Excess P as amorphous PmOn exists in the grain boundary of SnP2O7 crystal. Thisintergranular PmOn is considered to produce more channels for proton conductionand serves as superionic highways resulting in the increase of the conductivitieswith increasing Pini/Sn molar ratio. Under the other same conditions, the

conductivities increase with the order: (dry air) < (wet air) < (wet H2); (heat-treated at 873 K) < (heat-treated at 773 K). The highest conductivity is observed tobe 2.17 10 2 S cm 1 for the sample prepared from Pini/Sn = 2.8 under wetH2 atmosphere at 448 K. However, the stability of sample prepared from P ini/Sn of3.0 became poor, which may be attributed to its higher moisture absorbabilitythough its conductivities are higher.


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Fig. Partial conductivities of chargedspecies in the sample prepared fromPini/Sn of 2.4 (in wet hydrogenatmosphere).

Fig. Isotope effect on conductivity ofSnP2O7 prepared from different Pini/Snmolar ratios in argon saturated with H2Oor D2O vapor at 298 K.

Fig. Transport numbers of the sampleprepared from Pini/Sn of 2.4 in wethydrogen atmosphere.

The total conductivity is much higher than that of electronic conductivity, the protonic conductivity dominates ionic conduction,whereas the oxide-ionic conductivity reaches a certain extent in wet hydrogen atmosphere over the entire range of temperatures tested.

The oxide-ionic conduction is relevant to oxygen vacancies (VO

) in the sample.One may speculate about the reason for the presence of oxygen vacancies (VO ). H3PO4 and SnO2 as reactants for preparing the samplemay be assumed to contain trace of metal ions (Mn+) with lower ion valence than Sn4+ ions. While the sample is prepared, oxygenvacancies (VO ) are formed due to substitution of a small amount of Mn+ on the Sn4+-site, accordingly, resulting in the oxide-ionicconduction to a certain extent.

Ref 15 Intermediate temperature stable proton conductors based upon SnP O including additional H PO -J


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Ref. 15. Intermediate temperature stable proton conductors based upon SnP2O7, including additional H3PO4 J.

Mater. Chem. 20, (2010), 7827.

Abstract-In order to examine the influence of phospate impurities upon the conduction properties of SnP2O7, SnP2O7-H3PO4 compositeswere synthesised through different methods with varying starting P:Sn molar ratios. It was found that cubic SnP2O7 is the main crystallinephase and amorphous phases were observed when starting P:Sn ratios exceeded 2:1. Solid State 31P NMR confirmed residual phosphoricacid in samples with high starting P:Sn ratios whilst impedance spectroscopy indicated these to be good proton conductors. The highestconductivity observed was 3.5 102 S/cm at 300 C in air for samples with high starting P:Sn ratios and calcined at higher temperatures.The conductivity stability of the composites was found to be promising.

The most significant distinctions between different samples are the presence of an amorphous layer when the starting P:Sn ratio exceeds2:1. The amorphous layer should be rich in phosphorus as the thickness of layer increased from samples SP1 to SP3. Particles withrelatively clean edges and grain boundaries were observed for samples with stoichiometric ratio, irrespective of the starting materials. Thegrid distances marked on Fig. 4 correspond to (211), (210), (200), (111), (210) and (220) planes for cubic SnP2O7.

P:Sn molar

ratio

Raw Material Synthetic

history

Nomenclatur

e

2:1 SnO2.nH2O, H3PO4 300 OC, 2.5 h SP01

2:1 SnO2.nH2O, H3PO4 300 OC, 2.5 h650 OC, 2.5 h

SP02

2.6:1 SnO2.nH2O, H3PO4 300 OC, 2.5 h SP1

2.8:1 SnO2.nH2O, H3PO4 300 OC, 2.5 h SP2

4:1 SnO2.nH2O, H3PO4 300 OC, 2.5 h650 OC, 2.5 h

SP3

2:1 SnCl4.5H

2O, (NH

4)2HPO

4300 OC, 2.5 h650 OC, 2.5 h

SP4

Table Synthetic history and nomenclatures of all samples

Fig. 4 TEM images of SP02 (a) and (b); SP1 (c); SP2 (d);SP3 (e) and SP4 (f). The face distances marked of 3.22, 3.51 , 3.98 , 4.59 , 3.55 and 2.81 correspond to (211), (210), (200), (111), (210) and (220)planes respectively for cubic SnP2O7.

DP-MAS-NMR spectra show a number of characteristic peaks indicative of free phosphoric acid and tin coordinated groups
http://pubs.rsc.org/en/content/articlehtml/2010/jm/c0jm01089hhttp://pubs.rsc.org/en/content/articlehtml/2010/jm/c0jm01089hhttp://pubs.rsc.org/en/content/articlehtml/2010/jm/c0jm01089hhttp://pubs.rsc.org/en/content/articlehtml/2010/jm/c0jm01089hhttp://pubs.rsc.org/en/content/articlehtml/2010/jm/c0jm01089h


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DP-MAS-NMR spectra show a number of characteristic peaks indicative of free phosphoric acid and tin coordinated groups .

Conductivity-The samples with non-stoichiometric ratios, SP1, SP2 and SP3, are highlyconductive in both air and wet reducing conditions. A slightly enhanced conductivity in wetconditions implies that these samples are essentially proton conductors. Stoichiometric samplesshow a much lower conductivity and the conductivity is highly dependent on thermal treatment.A different treatment at 300 and 650 0C can yield a variation in conductivity of almost threemagnitudes (sample SP01 and SP02). The samples with poor conductivity are either mainly

SnP2O7 (SP02) or even pure cubic SnP2O7(SP4). The slightly higher conductivity in sampleSP02 is attributed to the residual acid and the presence of HPO42- groups. The significantlyhigher conductivity in samples SP01, SP1, SP2 and SP3 is probably due to the presence ofconsiderable amounts of H3PO4 and the SnP2O7 per se is not directly contributing to the highconductivity. This is consistent with the results of our previous report but demonstrates largediscrepancies with other groups. This is probably because of the excellent thermal stability andversatility of H3PO4 that is generally difficult to be removed under normal conditions. However,the observed high proton conductivity suggests a new route to develop novel proton conductorssince all samples are basically in solid state. The highest conductivity achieved was 3.5 x 10 -2S/cm at 300 0C in air for sample SP3.

Fig. Conductivity versus temperature ofSP01; SP1; SP2; SP3 and SP4 in different

atmospheres.

Fig. Conductivity of SP1, SP2 and SP3 at250 0C in different atmospheres as a

function of time.

Conductivity Stability- The conductivity of sample SP1 remains unchanged in both air and wet5% H2 for more than 50 h, confirming good stability. However, the conductivity of sample SP2keeps on decreasing over the first 50 h and tends to be stable thereafter which could be due tothe slow evaporation of H3PO4 as the initial acid loading in SP2 is relatively high. Anotherpossible explanation is the condensation effect: the slow generation of pyrophosphoric acidcould take place at high temperatures and decrease the conductivity to some extent. However,

the conductivity of sample SP3 was stable which was probably due to a better thermal stabilityafter the treatment at a higher temperature (650 0C). The SEM images of samples SP1, SP2 andSP3 taken after the conductivity stability testing show no obvious microstructural changescompared to those before the conductivity tests indicating that these samples are quite stable. ATEM picture of sample SP3 after the conductivity stability tests indicates that the core-shellmicrostructure remains.

Conclusion- The large difference in conductivity between composites and essentially purecubic SnP2O7 suggested that the conductivity mainly originates from the residual phosphoric

acid and cubic SnP2O7 itself is not a good proton conductor.

Ref 16 Sn In P O BasedOrganic/Inorganic Composite Membranes: Application to Intermediate Temperature


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Ref. 16. Sn0.9In0.1P2O7-BasedOrganic/Inorganic Composite Membranes: Application to Intermediate-Temperature

Fuel Cells- J. Electrochem. Soc. 154(1), 2007, B63.

Abstract- An anhydrous proton conductor, 10 mol % In3+-doped SnP2O7 was composed by 1,8-bis(triethoxysilyl)octane (TES-Oct) and 3-(trihydroxysilyl)-1-propanesulfonic acid (THS)Pro-SO3H. The composite membrane with 90 wt % Sn0.9In0.1P2O7 showed high protonconductivities of 0.04 S cm1 or more between 150 and 200C in dry air. The packing of the Sn0.9In0.1P2O7 particles in the matrix was relativelyuniform, with no formation of pinholes observed. Fuel cell tests verified that the OCV was maintained at a constant value of ~970 mVregardless of the electrolyte thickness (60200 m), while the Ohmic resistance was decreased to 0.24 cm2 by reducing the electrolytethickness to 60 m. The peak power densities achieved with dry H2 and air were 109 mW cm2 at 100C, 149 mW cm2 at 150C, and 187 mWcm2 at 200C. Furthermore, fuel cell performance was improved by hotpressing an intermediate layer consisting of Sn 0.9In0.1P2O7, Pt/C, TES-Oct, and (THS)Pro-SO3H between the electrolyte and cathode.

The composite membranes with only TES-Oct showed a lower conductivity compared to that in thepresence of the (THS)Pro-SO3H component. The observed difference can be ascribed to the -SO3Hgroups available for proton transfer. The SO3H group likely formed a proton conducting pathwayfrom one Sn0.9In0.1P2O7 cluster to another.

The conductivities of all the composite membrane samples decreased with increasing temperaturefrom 200 to 250C, which was irreversible upon heating and cooling. While the former behavior isdue to the dehydration of the membrane below 200C, the later behavior is due to the thermaldecomposition of components such as PTFE above 200C. Evidence for this reasoning is provided bythe TG analysis of the composite membrane and pellet samples.

XRD peaks observed for the compositemembrane sample were almost identical tothose of the pellet sample, suggesting that thepolymer components exist as an amorphousphase in the matrix.SEMs of (a) Sn0.9In0.1P2O7 compositemembrane and(b) pellet samples. Thecomposite membrane is more compact,

suggesting that the polymer functioned as abinder at the interface.

Measured in dry air

Fig. EMFs of H2 concentration cells with Sn0 9In0 1P2O7 composite membrane and pellet samples,


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g 2 0.9 0.1 2 7 p p p ,and their proton transport number as a function of temperature in dry atmosphere. It indicates thatthe polymer does not retard the proton transport number of Sn0.9In0.1P2O7.

Fig. (a) OCVs and (b) Ohmicresistance of fuel cells withSn0.9In0.1P2O7 composite membraneand pellet samples at 200C as afunction of electrolyte thickness. BothH2 and air were dry.

A decrease in the OCV with decreasing electrolyte thickness was observed for the pellet sample,implying that H2 or O2 crossover through the electrolyte increases with decreasing electrolytethickness.In contrast, the OCVs for the composite membrane were almost independent of the electrolyte

thickness, indicating that both H2 and O2 crossovers through the electrolyte are negligible, which isconsistent with the SEM.The lower OCV values for the composite membranes may be due to partial electron-holeconduction in the electrolyte, causing an internal short circuit of the fuel cell.The Ohmic resistances of two samples decreased linearly with decreasing electrolyte thickness. Inparticular, the composite membrane showed linearity over the thickness range from 60 to 150 m,assuming that the Sn0.9In0.1P2O7 powders were homogeneously distributed at ~10m level in thematrixes.The Ohmic resistance values of the composite membrane samples were always higher than those of

the pellet samples at the same electrolyte thickness, reflecting the difference in the protonconductivity for the two samples. However, the composite membrane sample showed a relativelylow resistance of 0.24 cm2, while maintaining a high OCV of 986 mV, which could not beachieved for pellet sample.

(a)

(b)At all of the tested temperatures, OCVs above950 mV were obtained, and no limiting currentbehavior was observed at high currentdensities. In addition, the current-voltage

slopes became lower as the operatingtemperature increased. The peak power densitythus reached 109 mW cm2 at 100C, 149 mWcm2 at 150C, and 187 mW cm2 at 200C.However, the power densities were muchlower compared to those expected from theOhmic resistances as an example, 0.24 cm2at 200C of the electrolyte, which may beascribed to the large polarization resistance ofthe fuel cell.

Dry H2 & air flow rate = 30 mL/minelectrolyte thickness = 60 m

Th t ti l l d i t d b th th d i th t t f


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Dry H2 & air, flow rate = 30 mL/minelectrolyte thickness = 60 m

Dry H2 & air, flow rate = 30 mL/minelectrolyte thickness = 60 m

The overpotentials were always dominated by the cathode in the temperature range of100200C; the cathodic overpotentials account for 8385% of the overall voltagedrops during the cell discharge. The anodic overpotentials as well as the IR dropsnegligibly affected the voltage drops especially at higher temperatures.The water formed electrochemically did not affect the subsequent cathode reactionbecause the discharge at high current densities gave rise to no limiting currents. It islikely that the charge-transfer reaction of oxygen reduction at the electrolyte/electrodeinterface proceeded at a very slow rate.Also, only cathode was in physical contact with the electrolyte, probably resulting inlow-density three phase boundaries (TPBs). Thus, it is necessary to optimize theelectrolyte/electrode interface.

no intermediatelayer

Withintermediatelayer

An attempt was made to hot-press an intermediate layer constructed fromSn0.9In0.1P2O7, Pt/C, TES-Oct, and (THS)Pro-SO3H between the electrolyte andcathode. As a result of the optimization of the weight ratio of the Sn 0.9In0.1P2O7

electrolyte to the Pt/C catalyst, the cathodic overpotential was shown to be the mostimproved when Sn0.9In0.1P2O7 :Pt/C = 20:1. The voltage drop was reduced by usingthe intermediate layer, so that the peak power density increased from 149 to 197mW cm2 at 150C. Similar effects were obtained at other temperatures, althoughthese were smaller at higher temperatures. One can consider that such performancegains are due to an increase in the area of the three-phase boundary for the cathodereaction. Indeed, the total polarization resistance decreased by 1.3 cm2 at 150Cby applying the intermediate layer. One may also expect that the fuel cellperformance is further enhanced by improving the microstructure, catalytic activity,

and proton conductivity of the intermediate layer.

Ref 17 Proton Conduction in SnP O LaP O Composite Electrolytes- Electrochem Solid-State Lett 12(2) 2009


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Ref. 17. Proton Conduction in SnP2O7LaP3O9Composite Electrolytes- Electrochem. Solid-State Lett. 12(2),2009,

B11.

Abstract-Theproperties of proton-conducting composite SnP2O7 electrolytes containing LaP3O9 are reported.These electrolytes demonstrateboth improved stability and enhanced proton conductivitieswhen compared to their individual constituents. Two different methods wereusedto prepare samples: sintering or warm-pressing. A conductivity of1.7104 S cm1 was obtained for warm-pressed SnP2O7LaP3O9 (Sn:La,82:18) at 350C.Under FC conditions the composite with composition Sn:La, 82:18exhibited the highest OCV of 0.983 V at 350C.TheOCV values confirmed that the electrolytes were predominantly ionically-conducting.

Fig. XRD patterns of SnP2O7LaP3O9,(Sn:La = 97:3, 82:18, 70:30, 50:50)composite electrolytes and SnP2O7 andLaP3O9. The arrowed reflections aretentatively associated with LaPO4.

Fig. Cross-sectional SEM images ofSnP2O7LaP3O9 composite electrolytes

(Sn:La = 50:50 (a), 70:30 (b), and 18:12 (c),and warm-pressed electrolyte with Sn:La =82:12(d).

At high La content the structure appears tocontain an amorphous component. For theSn:La 82:18, well-defined crystallites of size200 nm were observed. EDX showed that thecrystallites were mainly SnP2O7. In contrast,SEM image of warm-pressed SnP2O7LaP3O9,Sn:La = 82:18 (Fig. d) shows a nonporousstructure with a component that has an

amorphous morphology. This may be due to thelower preparation temperature that does notallow the full crystallization of SnP2O7.

The impedance plot shows the typical large electrode dispersion which is attributable to the

partially blocking nature of the Pt-paste electrode in a hydrogen containing environment.

Th d ti it d ith i i g L P O t t b t th


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Fig. Temperature dependencies ofac bulk conductivities for sintered

SnP2O7LaP3O9

The conductivity decreases with increasing LaP3O9 content but thetemperature stability range is significantly extended. The bestcombination of temperature/conductivity and stability wasobtained for a composition Sn:La = 82:18. For the warm-pressedsamples, a lower La content resulted in a porous microstructurewith a decreased stability under humidified environments. Theincreased conductivity in humidified, 4% hydrogen supports thepresence of proton conductivity. At temperatures below 350C, thewarm-pressed sample with Sn:La com position of 82:18 showedsignificantly higher conductivity, 1.7 x10-4 S cm-1, than thecorresponding sintered pellet at that temperature. This higherconductivity might be due to residual phosphoric acid that was notremoved by the warm pressing, which would be expected to leadto a rapid conductivity decline when held for a prolonged period atelevated temperatures. To verify this, the sample was held at300C. Instead, the conductivity remained steady for over 50 h,

after an initial drop of less than 5%. This increased stability at300C compares favorably to the rapid decline of conductivityabove 250C for the tin metaphosphates doped with In or Al.

Fig. Time dependency of theconductivity of SnP2O7LaP3O9(Sn:La = 82:18) prepared by

warm-pressing at 300C.

Fig. OCVs of fuel cell withsintered SnP2O7LaP3O9composite electrolytes. Thedashed line indicates thetheoretical OCV of ahumidified 4% hydrogenairfuel cell.

The measured OCVs are all lower than the theoretical OCV of1.18 V vs a standard hydrogen electrode. However, the OCVvalues are all higher than 0.9 V below 450C, while the cellwith electrolyte composition Sn:La = 82:18 exhibited thehighest OCV, 0.983 V, at 350C. This confirms that thematerials are predominantly ionic conductors. The differencebetween the theoretical and measured OCVs is likely to be due

to some H2 crossover because of a small amount of residualporosity in the sample, or possibly to some electronicconductivity. DC polarization measurements on other rare earthphosphate proton conductors did not indicate, however, asignificant electronic contribution to the conductivity.These results show that SnP2O7LaP3O9 composite electrolytescan display promising proton conductivities, where theconductivity/temperature stability can be favorably manipulatedby rare-earth additions.

Superprotonic KH(PO3H)SiO2 composite electrolyte for intermediate temperature fuel cells- Journal of Power


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Superprotonic KH(PO3H) SiO2 composite electrolyte for intermediate temperature fuel cells Journal of Power

Sources 194 (2009) 843846

Novel thin film composite electrolyte membranes, prepared by dispersion of nano-sized SiO2 particles in the solid acidcompound KH(PO3H), can be operated under both oxidizing and reducing conditions. Long-term stable protonconductivity is observed at 140 C, i.e., slightly above the superprotonic phase transition temperature of KH(PO3H),under conditions of relatively lowhumidification (pH2O0.02 atm).

Powders of KH(PO3H) were prepared by slow evaporation of an aqueous solutionobtained by mixing of H3PO3 and KOH in molar ratio 1:1. The powders were dried

in an oven at ~105 .C for 20 h, ground in an agate mortar and stored in a desiccatordue to hygroscopicity of the pure salt. Powders of KH(PO3H).SiO2 composites wereprepared by dispersing of SiO2 powder in an aqueous solution of KH(PO3H),followed by drying at ~105 .C for 20 h.

Fig- Dependence of proton conductivity on temperature for (a) pure KH(PO3H)(first heating scan under dry N2), and (b) KH(PO3H)SiO2 composites (air;pH2O 0.02 atm) at different mass fractions, , of SiO2 (particle size 14 nm) in thesamples. Also shown in (a) are data for pure KH(PO

3H) from differential thermal

l i (DTA) d d it S lid d d h d li id t th

Unlike sulphate and selenate solid acid electrolytes, KH(PO3H) can be operated inboth oxidizing and reducing atmospheres. Slight humidification of the gases in liquidwater at room temperature, with an equivalent pH2O of 0.02 atm, is sufficient toprevent impairment of the proton conductivity due to dehydration. Dispersion withnano-sized SiO2 particles leads to improved mechanical properties. The dispersion-strengthened composite electrolyte can be easily made into a thin film in the m

range by dip-coating from a suspension. Most important, the elevated temperature ofoperation and low humidification requirement of the thin film solid acid electrolytemembranes will lead to significant system simplifications in comparison withpolymer electrolyte fuel cells.

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Phosphate Containing Solid Electrolyte for ITFC