A Review of Mixed Ionic and Electronic Conducting Ceramic Membranes as Oxygen Sources for High-Temperature Reactors

Chapter 11

A Review of Mixed Ionic and Electronic Conducting Ceramic Membranes as Oxygen Sources for High-Temperature Reactors Qiying Jiang1, Sedigheh Faraji2, David A. Slade1 and Susan M. Stagg-Williams1,* 1

Chemical and Petroleum Engineering Department, University of Kansas, Lawrence, Kansas, USA Chemical Engineering Department, California State University, Long Beach, California, USA * Corresponding author: E-mail address: [email protected] 2

INTRODUCTION Mixed ionic–electronic conducting (MIEC) ceramic membranes have received substantial interest in recent decades for various applications requiring gas separations [1–3]. Dense (i.e., nonporous) oxygen-conducting MIEC ceramic membranes allow oxygen ion diffusion through the solid ceramic lattice, resulting in theoretically infinite selectivity for oxygen. Unlike the oxygenconducting ceramic membranes used in solid-oxide fuel cells (SOFCs) such as yttria-stabilized zirconia (YSZ), the ability of MIECs to conduct electron and oxygen ions simultaneously enables these membranes to operate without an external electrical circuit. An oxygen-MIEC ceramic membrane requires only different gas-phase environments on its two surfaces and a sufficiently high temperature to transport oxygen from the high oxygen content surface to the low oxygen content surface [1–4]. Under an oxygen potential gradient and above some threshold activity temperature, a fully densified oxygen-conducting membrane would exhibit perfectly selective oxygen production into the low oxygen environment [5]. At temperatures above their threshold activity temperature (typically greater than 600
C [6]), dense oxygen-conducting ceramics can be characterized in general as exhibiting (1) highly mobile lattice oxygen ions; (2) continuously variable oxygen content (i.e., oxygen nonstoichiometry) via the ability to support lattice oxygen defects and/or undergo gradual, dispersed phase changes; and (3) the surface abilities to dissociate molecular oxygen while incorporating oxygen ions and Membrane Science and Technology, Vol. 14. # 2011, Elsevier B.V. All rights reserved.

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to reassociate oxygen ions while evolving molecular oxygen [7]. Between their respective activity threshold temperatures and degradation temperatures, oxygen-MIEC ceramic materials exhibit oxygen fluxes that generally increase with both temperature and the magnitude of the imposed oxygen gradient. Oxygen-MIEC membranes have been explored for a variety of applications, including oxygen supply via air separation [2,8–11], energy production [4,12,13], SOFC anodes [14–20], and oxygen sources for high-temperature reactions. The high-temperature reactions of interest include partial oxidation of methane (POM) [21–26], other hydrogen production reactions [13,27–29], oxidative dehydrogenation of light alkanes [30–32], and oxidative coupling of methane [33–38]. Current commercial barriers for utilizing membranes in high-temperature hydrocarbon conversion reactors include low membrane oxygen flux, poor mechanical stability, and high membrane fabrication costs. The area of greatest interest in the research community to date, and of particular focus in this chapter, is the use of oxygen-MIEC membrane reactors for synthesis gas (syngas) production, particularly via POM. POM requires pure oxygen as a feedstock, and the high energy and safety-related costs associated with industrial-scale oxygen production provide a substantial incentive to develop oxygen-MIEC membranes as alternative oxygen sources for syngas production reactors [5,6,10,21,24,25,39–43]. The major challenge in developing membranes for this application involves the historically incompatible requirements of high oxygen transport and high material stability. These requirements have proven particularly difficult to reconcile in the strongly reducing environments of methane conversion reactors. After providing an overview of oxygen-conducting MIEC ceramic membrane materials, this chapter discusses recent work to overcome the unique challenges involved in incorporating these membranes into syngas production reactors. The scope includes the relationships between crystal structure, oxygen permeation, stability, and reactor performance.

GENERAL ATTRIBUTES OF OXYGEN-CONDUCTING MIEC CERAMIC MATERIALS Oxygen Nonstoichiometry The oxygen conduction phenomenon exhibited by oxygen-MIEC membranes is attributed to their ability to support oxygen vacancies and lattice disorder, which allows the relatively rapid and sustainable transport of oxygen ions and holes under the appropriate conditions [2–4,19,31]. The number of oxygen vacancies is a function of material composition, initial lattice structure, temperature, and ambient gas composition. Membrane material composition, temperature, and oxygen gradient are the dominant factors in determining oxygen flux [14,15,44–46]. The oxygen transport mechanism of oxygen-MIEC membranes is shown in Figure 11.1. First, an oxygen molecule from the gas phase must adsorb on the high oxygen potential surface of the ceramic, at which time it dissociates into

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Reaction on low [O2-] surface 4e- + 2O2- ® O2

Reaction on high [O2-] surface O2 + 4e- ® 2O2Air

e-

Sweep gas

O2Oyxgen-depleted air

Sweep gas enriched with

FIGURE 11.1 O2 transport mechanism in a dense MIEC membrane.

two oxygen ions by receiving electrons from the membrane surface. Oxygen ions then migrate from the high to low oxygen potential surface via the oxygen vacancies in the lattice. In an inert environment, the oxygen ions lose electrons at the low oxygen potential surface, recombine to form molecular oxygen, and desorb into the gas phase. In a reactive environment, the low oxygen surface reactions will depend on the other chemical species present in the gas phase and adsorbed on the low oxygen surface. It is well established that equilibrium oxygen ion content in these materials varies with both temperature and ambient oxygen partial pressure, and oxygen content losses beyond a material’s phase stability limit can cause chemical decomposition. Under a constant oxygen partial pressure, equilibrium oxygen ion content decreases with increasing temperature. In some cases, certain constituents can even be reduced to their metallic state by temperature alone [47]. A reducing atmosphere (e.g., one containing hydrogen and/or methane or, according to Ma and Balachandran [48], one with an oxygen partial pressure < 10 6 atm) can produce the same effect at even lower temperatures [35,49].

Self-Adjusting Phase Equilibria For metal oxide materials with labile oxygen ions, changes in oxygen content following changes in temperature and/or oxygen environment can be manifested as phase adjustments. Studies using in situ X-ray diffraction (XRD) have confirmed that responsive phase adjustments in oxygen-MIEC ceramics are common. Depending on both the material and the new environment, individual phase changes can occur throughout the entire membrane or they can be highly localized (e.g., only at a surface or as distributed pockets of a newly formed phase in an equilibrium mixture of phases) [35,42,49]. For a material with a large variety of phase options that is exposed to a continuous oxygen partial pressure gradient, it is reasonable to expect the formation of a phase composition gradient that reflects the electrochemical oxygen potential distribution in the membrane.

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Numerous researchers have confirmed that certain materials, most notably SrFeCo0.5O3d (SFC), can consist entirely of intimately intermixed phases. Such mixtures are referred to as “phase assemblages” or “solid solutions,” and they will adjust their phase distribution according to their environment [50–53]. Phase distribution changes in oxygen-MIEC ceramics have been observed to be readily reversible [25,48,54], with the consequence that equilibrium phase distributions can be highly sensitive to both temperature and gaseous environment [51]. However, phase transition kinetics for solid-phase equilibrium transitions are slow enough at lower temperatures that phase change reversibility can be masked or suspended [55]. This allows nonequilibrium compositions to be maintained at temperatures below the activity threshold. Another consequence of the phase lability of MIEC ceramics is their potential to exhibit known phase changes at lower temperatures than expected in response to gas-phase oxygen content manipulation [53]. This phenomenon confirms that equilibrium phase composition and phase distribution are determined by membrane oxygen content, which depends on both temperature and gaseous environment. This feature increases both the opportunities for, and the challenges associated with, oxygen-MIEC membrane applications. Membrane fracture, which is the common shortcoming of ceramic oxygen-MIEC materials, particularly in partial oxidation applications [5], can result from chemical decomposition in the form of phase conversion and segregation or from lattice expansion mismatch within either a single phase or a set of phases [39,54].

Chemical Expansivity Oxygen-MIEC ceramics generally expand as their oxygen content decreases. A decrease in oxygen content increases lattice oxygen vacancies and reduces the oxidation states of a portion of the membrane’s metal ions, both of which diminish the overall binding forces within the lattice and thus allow it to expand even under isothermal conditions [56–59]. To illustrate the potential extent of lattice expansion from compositional changes, the perovskite-to-brownmillerite phase transition has been observed to produce a unit-cell volume increase of up to 6% for SrCo0.8Fe0.2O3d (SCF) [57]. Adler has proposed that this “chemical expansivity” should be considered a new physical property for these materials. He concludes that chemical expansion from oxygen content decrease can be significantly greater than thermal expansion alone and argues that labeling as thermal expansion— the volume increases of oxygen-conducting MIECs during temperature increases—oversimplifies the phenomenon [56]. Oxygen-MIEC nonuniform chemical expansion is a much more likely explanation for a fracture-inducing density difference than nonuniform thermal expansion. At high temperatures, electronic conductivity is a characteristic feature of these mixed-conducting materials and is necessary to their oxygen

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transport function. The Wiedemann–Franz law states that the ratio of thermal conductivity to electronic conductivity increases with temperature for solid materials. At temperatures high enough to conduct electrons (i.e., approaching 1000 K), these materials should therefore exhibit sufficient thermal conductivity to support the assumption of nearly isothermal behavior at steady state. The membrane’s lattice oxygen content, on the other hand, is clearly not isocratic, so a chemical expansion gradient is inevitable. The likelihood that an imposed chemical gradient will be more extreme than a possible thermal gradient supports the argument that chemical expansion under operating conditions is both greater and more important than thermal expansion.

Microstructure of Oxygen-MIEC Ceramics Researchers have been trying to relate oxygen transport through MIEC membranes to membrane microstructure and then use this understanding of the relationship to design membranes with higher oxygen flux [60–66]. The microstructure properties of interest include grain size, the distribution of the grain boundary, and inhomogeneous grain boundary composition (i.e., differences between grain boundary and bulk compositions). Recent studies have suggested that the grain boundary provides the pathway for oxygen transport in the membrane material and that sintering conditions can significantly impact grain boundaries. One study on La0.5Sr0.5FeO3d showed that as sintering temperature and time increased, grain size increased while grain boundary decreased [61]. The resulting oxygen permeation experiments showed that membranes with larger grain sizes exhibited lower oxygen fluxes. A similar phenomenon was observed on LaCoO3d membranes [62]. It was suggested in these studies that the grain boundary acts as a pathway for oxygen transport, so that a membrane with smaller grain boundaries produces lower oxygen flux than a membrane of the same bulk material with more extensive grain boundaries. Similar to the observations with the lanthanum-based membranes, studies on CaTi0.8Fe0.2O3d, Ba1xSrxCo0.8Fe0.2O3d, and Ba0.5Sr0.5Fe0.8 Zn0.2O3d membranes have shown that increasing sintering temperature and time produces larger grain size [60,65,66]. However, in contrast to the lanthanum-based membrane studies, increased grain size was correlated with significantly increased oxygen fluxes with the Ca- and Ba-based membrane materials. In these cases, the grain boundary seemed to act as a barrier to oxygen transport. The seeming contradiction in the effect of sintering conditions on the oxygen flux can be explained by differences in the role of the grain boundary in oxygen transport. For materials in which the grain boundary acts as a barrier to diffusion, larger grains and thus, smaller grain boundaries, will result in increased oxygen flux [64]. In contrast, for materials in which the oxygen transport is facilitated by the grain boundaries, smaller grains with larger grain boundaries will increase oxygen flux. While this explanation suggests that the grain

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boundary has a profound effect on the oxygen flux, it does not reveal the true mechanism of oxygen transport through the grain boundary or provide a simple way to the effect of grain boundary on flux. More research in this area is necessary to answer these questions.

COMMON OXYGEN-MIEC MEMBRANE MATERIALS Oxygen-MIEC membrane materials can be classified into two main groups based on crystal structure: fluorite and perovskite. In recent years, dual-phase composite MIEC membranes have also been developed in an attempt to create membranes with both high oxygen permeability and good mechanical stability. These dual-phase membranes are the focus of Chapter 12 and will only be introduced briefly here.

Fluorites The fluorite oxide structure is represented by AO2, where A is the large fourvalent cation such as Zr4þ and Ce4þ [67]. Fluorite contains oxygen anions in simple cubic packing, with half of the interstices occupied by metal cations. The metal cations occupy the cube centers and the oxygen anions are tetrahedrally coordinated to the metal cations [68]. Two fluorite oxides known to have particularly high oxygen ion mobility, ZrO2 and Bi2O3, have been extensively investigated for use as oxygen-MIEC membrane materials. At room temperature, ZrO2 has a monoclinic crystal structure. As the temperature increases, the crystal structure of ZrO2 transforms to the tetragonal (> 1000
C) and cubic structures (> 2300
C) [69]. It is generally believed that cubic ZrO2 has a high oxygen ionic conductivity [70]. To maintain the cubic phase at room temperature and thus eliminate the mechanical stresses created by volume expansion from phase transition, dopant materials such as CaO, MgO, and Y2O3 can be added to pure ZrO2 [69,71,72]. In this scenario, the solid solution is referred to as “stabilized ZrO2.” Bi2O3 has four polymorphs, designated a-, d-, b-, and g-Bi2O3. Among the four phases, the d-phase Bi2O3 exhibits high oxygen ionic conductivity [73]. As with ZrO2, compounds such as Y2O3 and Er2O3 are added to pure Bi2O3 to maintain the d-phase at low temperatures [74,75]. Compared to ZrO2, Bi2O3 shows higher oxygen ionic conductivity at intermediate temperatures (< 800
C). The doped Bi2O3 oxides, especially BIMEVOX (where “ME” represents a metal ion, such as copper or cobalt, and “V” represents the vanadium), are suggested to be some of the highest ionic conductors at intermediate temperatures [8,76,77]. Even though ZrO2 and Bi2O3 have high oxygen ionic conductivity, their electronic conductivity is low (the electronic transference number is close to zero). These materials are thus commonly used as the electrolyte or anode in SOFC systems. To form oxygen-MIEC ceramic membranes, a dopant

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compound with a high electronic conductivity has to be introduced to ZrO2 or Bi2O3. If the added compound has similar thermal properties and can form a solid solution with ZrO2 or Bi2O3, the doped ZrO2 or Bi2O3 becomes a single-phase oxygen-MIEC material. CeO2 is a primary dopant because cerium has multivalent states (þ 4, þ 3) and the transition from Ce4þ and Ce3þ increases the electronic conductivity of the material [12,74,78,79]. However, even with the addition of cerium, doped fluorite compounds still show much lower electronic conductivity than ionic conductivity. For a relatively thick fluorite-based oxygen-MIEC membrane, the oxygen permeation is therefore determined by electron conduction [12,74,80–82]. Table 11.1 provides an overview of the more common fluorite-type oxygen-MIEC membrane materials reported in the literature to date.

Perovskites Any metal oxide with the general formula ABO3, where A is the larger cation with a 12-fold oxygen ion coordination and B is the smaller cation with a 6-fold coordination, is considered to be a perovskite oxide. Twelve oxygen anions are cuboctahedrally coordinated to the A-site cations and six oxygen ions are octahedrally coordinated to the B-site cations. After doping with other metal cations, the perovksite can be symbolized by the formula AxA10  xByB10  yO3  d [85]. Generally speaking, A-site ions are rare-earth metals, A0 -site ions are alkaline-earth metals such as Ca2þ, Sr2þ, and Ba2þ, and B- and B0 -site ions are transition metals such as Co3þ and Fe3þ. A single-phase perovskite compound with five or more constituent metal species is rare, while compounds with three and four metal species are quite common. Goldschmidt defined a tolerance factor t to evaluate the structural stability of perovskite materials: RA þ RO2 ; t ¼ pffiffiffi 2 RB þ RO2

ð11:1Þ

where RO2 is the radius of the oxygen ions, RA is the radius of the A-site ion, and RB is the radius of the B-site ion. In general, structural stability of a material increases as t increases, with t-values between 0.75 and 1.0 representing a stable material [86]. Oxygen transport in perovskites is largely determined by the activation energy for oxygen ion conduction. Reducing the activation energy will effectively increase oxygen ion diffusion and enhance net oxygen flux. Researchers believe that this activation energy depends on the average metal-oxygen bond energy, the lattice free volume, and the radius of opening between the two A-site cations and one B-site cation [14,15]. Low metal-oxygen bond energy, large lattice free volume, and large openings are associated with lower activation energy. Lattice free volume

TABLE 11.1 Comparison of Oxygen Fluxes for Various Fluorite Membranes Reported in the Literature Membrane Material Bi1.5Y0.3Sm0.2O3 [34] (Bi2O3)0.73(CaO)0.27 [75] (Bi2O3)0.75(Y2O3)0.25 [83]

Temperature (
C) 750–950 600–680 800–950

Oxygen Flux (mol s 1 m 2) 5

3.5  10

4

–4.3  10

Disk

1.27

5

Disk

1.2

5

4

Disk

1.4

Disk

2

Tube

1.5

Tube

1.5

Tube

2

Disk

2

Disk

3  10 3–16  10 3

7.840  10

BiY0.5Cu0.5O3 [81]

650–850

1.0  10

0.8ZrO2–0.1TiO2–0.1Y2O3 [71]

1305–1481

0.018–0.12 9  10

[(ZrO2)0.7(CeO2)0.3]0.9(MgO)0.1 [72]

1027–1477

0.12–3.8

(ZrO2)0.8(Y2O3)0.20 [84]

900 900–1000

–3.351  10 5

–7.6  10

1305–1481

(ZrO2)0.7(Tb2O3.5)0.3 [80]

–3.300  10

5

0.825ZrO2–0.075TiO2–0.1Y2O3 [71]

2

–4.2  10

7

2.6  10

5

2.1  10

Thickness (mm)

5

1.230  10

3

Membrane Configuration

5

–5.5  10

To convert to molar fluxes, all volumetric fluxes reported in the literature were assumed to be at T ¼ 25
C and P ¼ 1 atm unless otherwise stated in the reference.

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increases as the radii of the A- and B-site cations increase. Furthermore, the combination of large A-site cations and small B-site cations will increase the size of the opening between the two A-site cations and one B-site cation. Two principles should be noted when doping A- and B-site cations (1) the radii of dopants must match the radii of the cations that the dopants replace and (2) the concentration of heterovalent dopant within the perovskite lattice must less than the level at which the introduced lattice vacancies become ordered [15,44,45]. One example of the effect of lattice order on oxygen conductivity can be found in the orthorhombic brownmillerite structure (ABO2.5). A brownmillerite is essentially a highly ordered perovskite derivative that can occur when a variable oxygen content perovskite phase approaches an oxygen deficiency of 16.7% (i.e., one-sixth) [5]. It is generally believed that the highly ordered oxygen ion vacancies in brownmillerites lead to low oxygen flux because the brownmillerite phase is more stable than the related perovskite and does not support oxygen hole transport. However, at high enough temperature brownmillerite phases may disorder sufficiently to return the material to a high-vacancy perovskite structure with higher oxygen transport capability [11,39,49,87,88]. The difference in behavior observed between perovskites and their associated brownmillerite phases confirms that oxygen vacancy disorder can be a determining factor for oxygen conductivity. A- and B-site perovskite doping studies with heterovalent metal ions have produced many oxygen-MIEC compounds. Table 11.2 provides the compositions and properties of some of these perovskite-based membrane materials. Because a detailed review covering the entire progress of these perovskite oxygen-MIEC materials is beyond the scope of this book, the discussion will be limited to the most commonly studied perovskite or perovskite-like materials.

SCF-Based Materials Perovskites made from strontium, iron, and cobalt have attracted substantial attention since Teraoka [46] published his pioneering work with this family in 1985. Compared to other early oxygen-MIEC materials, Sr–Co–Fe perovskite materials, particularly SrCo0.8Fe0.2O3d, show a notably high oxygen flux. Teraoka et al. [46] reported an oxygen flux of about 0.02 mol m 2 s 1 in an air: helium gradient at 900
C with a 1-mm-thick disk-shaped SrCo0.8Fe0.2O3d membrane. It should be noted that SCF is not a perovskite-type compound at room temperature. However, it undergoes a phase transition from the brownmillerite phase to the defect perovskite phase at elevated temperatures [101]. This phase transition represents an increase in lattice oxygen vacancy disorder, which is responsible for the observed oxygen flux increase. Although SCF has exhibited high oxygen flux, additional work demonstrated that the SCF material suffers from poor chemical and structural stability,

TABLE 11.2 Comparison of Oxygen Fluxes for Various Perovskite Membranes Reported in the Literature Membrane Material

Temperature (
C)

Oxygen Flux (mol s 1 m 2)

Membrane Configuration

Thickness (mm)

(BiO1.5)0.86BaO [90]

650–950

3.7  10 5–9  10 3

Disk

0.55

Disk

1.1

650–950 650–950 Bi0.85Sr0.15FeO3 [90] Bi0.7Sr0.3FeO3 [90] Bi0.4Sr0.6FeO3 [90] Bi0.2Sr0.8FeO3 [90] BaCe0.4Fe0.6O3d [91] BaCo0.4Fe0.5Zr0.1O3d [9] Ba0.5Sr0.5Co0.8Fe0.2O3d [92] Ba0.5Sr0.5Zn0.2Fe0.8O3d [93] CaTi0.8Fe0.2O3d [17] La0.7Ca0.3CrO3d [20] LaCo0.8Fe0.2O3d [94] La0.6Sr0.4Co0.8Cr0.2O3d [94]

800 800 800 800 800–950 700–950 850–900 800–975 800–1000 800 860 860

5

3.7  10

5

3.7  10

3

–7.3  10

3

–6  10

Disk

1.65

4

Disk

1

4

Disk

1

4

Disk

1

Disk

1

4.2  10 7.7  10 2.2  10

3

1.1  10

4

7.4  10

3

2.0  10

4

9.0  10

2

1.1  10

5

8.0  10

3

–1.8  10

1

Disk

1

2

Disk

1.8

2

Disk

1.45

4

Disk

1

Disk

1.07

Disk

1

Disk

1.5

–6.8  10 –1.6  10 –2.6  10 –2.2  10

5

6  10

4

1.8  10

3

4.21  10

Disk

3

La0.6Sr0.4Co0.2Fe0.8O3d [95] La0.8Sr0.2Ga0.7Co0.3O3d [96] (2%wt)Y0.2Ce0.8O1.9–(98%wt)La0.7Ca0.3CrO3d [20] (4%wt)Y0.2Ce0.8O1.9–(96%wt)La0.7Ca0.3CrO3d [20] (6%wt)Y0.2Ce0.8O1.9–(94%wt)La0.7Ca0.3CrO3d [20] (8%wt)Y0.2Ce0.8O1.9–(92%wt)La0.7Ca0.3CrO3d [20] (10%wt)Y0.2Ce0.8O1.9–(90%wt)La0.7Ca0.3CrO3d [20] (La0.75Sr0.25)0.95Cr0.5Mn0.5O3d [16] SrSc0.5Co0.95O3d [97] La0.6Sr0.4Fe0.9Ga0.1O3d [98] (98%vol)La0.6Sr0.4Fe0.9Ga0.1O3d–(2%vol)MgO [98] (95%vol)La0.6Sr0.4Fe0.9Ga0.1O3sd–(5%vol)MgO [98] La0.4Sr0.6Co3d [46] SrCo0.9Nb0.1O3d [99]

850–900 700–1000 800 800 800 800 800 950–1000 675–900 875–975 825–925 825–925 870 700–900 700–900 700–900

SrFe0.6Cu0.3O3d [100]

750–950

4  10 4–1.1  10 3 3

2.3  10

2

–1.1  10

5

1.4  10

5

3.5  10

5

6  10

5

6.4  10

5

7  10

7

1  10

5

–1.3  10

3

1.8  10

5

2.2  10

5

4.4  10

4

1.2  10

2

–2.2  10

4

–1.2  10

4

–2.7  10

4

–4.2  10

3

3.8  10

3

2.5  10

2

–1.8  10

3

Tube

0.219

Disk

0.5

Disk

1.07

Disk

1.07

Disk

1.07

Disk

1.07

Disk

1.07

Disk

1

Disk

1

Disk

1

Disk

1

Disk Disk

1

Disk

1.5

2

–2.7  10

5  10

3

5.5  10

3

2.6  10

1

2

–3.2  10

3

–7.1  10


0.7 Disk

1.5

To convert to molar fluxes, all volumetric fluxes reported in the literature were assumed to be at T ¼ 25 C and P ¼ 1 atm unless otherwise stated in the reference.

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Inorganic, Polymeric and Composite Membranes

especially in a reducing atmosphere [54,88]. For example, severe cracks occurred in an SCF tubular membrane after only a few minutes of exposure to a POM environment [54]. Extensive work has been performed to improve the stability of SCF with the addition of promoters.

Promoted SCF Materials The addition of small amounts of silver to an SCF material led to an increase in oxygen flux, with the highest flux obtained at a silver content of 5 mol%. Upon further investigation of the Ag–SCF membrane, the added silver was determined to have increased the surface oxygen exchange rate [102]. However, the addition of silver did not impact the phase transition between the brownmillerite and perovskite phases, and the stability of the Ag–SCF membrane was not significantly improved. Wu et al. [103] attempted to utilize the thermal and chemical stability of Al2O3 to improve the stability of SCF. They were able to significantly improve the stability of the SFC in low oxygen partial pressure, high-temperature environments by introducing Al2O3. However, because Al2O3 reacts with cobalt to form a spinel phase (CoAl2O4) at elevated temperatures (1200
C), the doping of Al2O3 into SrCo0.8Fe0.2O3d resulted in a loss of cobalt from the perovskite phase and a dramatic decrease in membrane oxygen flux. Fan et al. [104] investigated the use of tin (Sn) to improve the stability of SCF. The introduction of tin as SrSnO3 not only decreased the thermal expansion coefficient of the material, but also reduced the onset temperature of the oxygen permeation to 560
C. Decreased thermal expansion is associated with greater mechanical stability, and a decrease in activity threshold temperature increases the potential utility of a membrane material. However, the oxygen permeability of the SrSnO3-doped SCF was lower than that of the parent SCF at temperatures higher than 850
C. Substituted SCF Materials Other researchers have modified the SrCo0.8Fe0.2O3d material by partially substituting for strontium in the A site of the perovskite structure. One example of this is La0.6Sr0.4Co0.2Fe0.8O3d [105,106]. Oxygen flux increased with temperature in a series of tests with a La0.6Sr0.4Co0.2Fe0.8O3d tubular membrane with 1.5-mm wall thickness. At 850
C, the oxygen flux was 9.67  10 4 mol m 2 s 1 under a helium:air gradient. However, the performance of the LSCF membrane deteriorated over time under this gradient, and the loss of flux was ascribed to the perovskite phase on the surface of the membrane decomposing to SrSO4, CoSO4, SrO, Co2O3, and La2O3 over the 110 h of testing [106]. Some researchers have substituted cobalt-doped materials with less reducible ions such as Ga4þ as a means of improving the chemical stability of the material. It was believed that materials with lower thermal expansion

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coefficients would have less lattice expansion and might be more stable in reducing environments [99,107]. Oxygen flux testing on a disk-shaped La0.7Sr0.3Ga0.6Fe0.4O3d membrane with 0.5-mm thickness showed that at 1000
C, oxygen fluxes of 0.012 mol m 2 s 1 under an air:helium gradient were achieved [107]. Although the study indicated that the high-density La0.7Sr0.3Ga0.6Fe0.4O3d ceramic had a low thermal expansion coefficient, it also showed that this material is susceptible to decomposition to LaSrGaO4 at temperatures higher than 700
C in a reducing atmosphere [108]. The most successful modification of SrCo0.8Fe0.2O3d (SCF) has been Ba0.5Sr0.5Co0.8Fe0.2O3d (BSCF). A 1.8-mm-thick disk-shaped BSCF membrane has exhibited an oxygen flux of 9.5  10 3 mol m 2 s 1 in an air:argon gradient at 850
C [92]. This partial substitution of the A-site strontium cation with barium has been shown to improve phase stability by preventing oxidation of the B-site cation, thus maintaining perovskite content in the membrane material [109]. However, some recent studies [60,110,111] have shown that the BSCF membranes are sensitive to CO2 because of the alkaline-earth elements in the structure. For example, Jiang et al. [111] investigated the performance of the BSCF membrane in the CO2 reforming reaction. The results showed that the presence of a large amount of CO2 leads to surface restructuring of the BSCF membrane, which decreases the oxygen flux. This restructuring is ascribed to the formation of carbonate on the reaction surface of the membrane.

Multiphase SrFeCo0.5Ox Balachandran et al. [112,113] adjusted the composition of Sr–Co–Fe perovskite to develop a new material, SrFeCo0.5Ox (SFC). Compared to SrCo0.8Fe0.2O3d, SFC showed good mechanical integrity under reaction conditions with a strongly reducing environment, and was reported to be quite stable even under the extreme oxygen potential gradients in a POM membrane reactor [115]. Initial reports of SFC’s high oxygen flux by Ma et al. and Maiya et al. led to numerous investigations of SFC, but the subsequent work demonstrated a large variation in oxygen flux with SFC membranes, with most results much lower those initially reported [39,40,50,51,55,115–117]. Stoichiometrically, SFC is not a perovskite material. However, SFC typically occurs as an adaptable three-phase solid solution (or “phase assemblage”) with small, highly intermixed phase moieties, including perovskite phases [48,50–53,55,114,116]. The two-phase categories observed in SFC samples in addition to the perovskite phases (SrFe1xCoxO3d) are spinel phases (Co3xFexO4) and the so-called “intergrowth” phases (SrFe1.5xCoxO3.25d). The name “intergrowth” refers to the structural appearance of this phase as interspersed perovskite and brownmillerite units [50]. All three phase types have adjustable stoichiometry, leading to a highly interdependent distribution of cations and phase shifts within phase categories as well as between phase categories. It has been shown that SFC’s initial phase

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composition can vary substantially with both material preparation and membrane fabrication methods [116]. The ratios and compositions of the three common phase types depend on the synthesis method, the preparation procedure, and the ultimate environment of the prepared material. One additional phase type has also been reported to appear on highly reduced membrane surfaces: the Co–Fe spinel phase can decompose to a cubic “rocksalt” phase (Co1xFexO) at temperatures greater than 900
C in air or at lower temperatures in reducing environments [52,53]. The overall body of work has determined conclusively that oxygen flux through SFC membranes correlates directly with its perovskite phase content [50,51]. Xia et al. [51] observed the creation of greater amounts of perovskite phase with annealing temperatures above 1150
C and also noted that the oxygen flux for SrFeCo0.5Ox slowly decreases over time due to a slow transition of perovskite phase to the intergrowth phase. Other researchers have noted increased perovskite content with exposure to high temperature/low pO2 environments (i.e., conditions that lead to low equilibrium oxygen content in the SFC material) [48,50,53]. Ikeguchi et al. [116] reported a 17% decrease in oxygen flux with a decrease in membrane sintering temperature from 1200 to 1150
C, with a decrease in perovskite content as the presumptive cause of the flux decrease. Armstrong et al. studied carefully prepared samples of known phase composition for each of the three phase types, and the oxygen flux results correlated emphatically with perovskite phase content. Oxygen flux through the intergrowth phase samples was two orders of magnitude lower than that with the representative perovskite material (SrFe0.75Co0.25O3d) and one order of magnitude lower than that of a typical three-phase SFC “phase assemblage” material [50]. Xia et al. [51] also reported slow conversion back to an intergrowth phase at temperatures below 1000
C. The relatively slow transition to the intergrowth phases at lower temperatures allows “meta-stable” perovskite phases to persist at lower temperatures than expected.

Dual-Phase Composite Materials It is clear from the literature that several key challenges still remain with singlephase oxygen-MIEC membranes. The oxygen flux of the MIEC membranes is highly correlated with the ionic and electronic conductivity of the material, with both high ionic and electronic conductivity necessary to achieve high oxygen flux. However, limited success has been found in identifying materials with both high ionic and electronic conductivity. Furthermore, materials that exhibit high oxygen flux often suffer from lower thermal or chemical stability when exposed to the large oxygen partial pressure gradients present for most membrane reactor applications. This is especially true when the membrane is used in a hydrocarbon conversion reactor, where one surface is exposed to a highly reducing reaction environment and the other surface is exposed to an oxygenrich environment.

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249

As work to overcome the challenges in single-phase oxygen-MIEC membrane development continues, researchers have been evaluating dual-phase composite materials as an alternative approach to fabricating membranes with high ionic and electronic conductivity while maintaining chemical and thermal stability [118–123]. A dual-phase composite membrane combines a material with high oxygen ionic conductivity with a second material with high electronic conductivity. In these materials, oxygen ion and electron transport occur in different phases. The ion- and electron-conducting phases must form a continuous and intimately intermixed network to allow simultaneous and balanced transport of oxygen ions and electrons under the electroneutrality constraint of zero net charge. The conductivity of the composite material therefore depends strongly on the volume ratio of two phases. Theoretically, it is easy to design high oxygen flux membranes by selecting appropriate materials and adjusting the ratio of two phases. As mentioned earlier, fluorite materials such as YSZ have high oxygen ionic conductivity. Mazanec et al. [120] were able to achieve high oxygen fluxes by incorporating noble metals such as platinum into YSZ. However, further study showed that at least 30 vol% of the metal has to be added to the ceramic material to form an effective electron-conducting phase network [121]. From a practical perspective, the high price of noble metals restricts the development and application of composite membranes that include a noble metal. Replacing expensive noble metals with perovskite materials that exhibit high electronic conductivity is a promising alternative. For example, Sirman et al. [122] produced a dual-phase composite with a very low thermal expansion coefficient by mixing Ce0.8Gd0.2O1.9 (a fluorite oxygen ion conductor) and La0.8Sr0.2Fe0.8Co0.2O3d (a perovskite electron conductor) in equal parts by volume and sintering in the temperature range of 1200–1400
C. This dualphase composite membrane exhibited an oxygen flux of 0.042 mol m 2 s 1 at 1000
C. Wang et al. [123] successfully prepared a La0.15Sr0.85Ga0.3 Fe0.7O3d–Ba0.5Sr0.5Co0.8Fe0.2O3d (LSGF–BSCF) composite that exhibited an oxygen flux of 3.1  10 3 mol m 2 s 1 at 915
C, which was about nine times greater than the oxygen flux through a single-phase LSGF membrane. Table 11.3 provides an overview of the dual-phase composite membrane work available in the literature and a more detailed discussion of recent work in dual-phase membranes is provided in Chapter 12.

Membrane Modifications to Improve Oxygen Flux Oxygen transport through dense oxygen-MIEC membranes requires multiple steps, as shown previously in Figure 11.1. For a given membrane, if the ratelimiting step of the oxygen permeation is determined, appropriate modification techniques can be employed to enhance the performance of the membrane. For example, if oxygen transport is limited by the surface exchange reactions, a surface modification such as a catalyst coating or an increase in surface area can increase the surface exchange of oxygen and thus increase membrane oxygen

TABLE 11.3 Comparison of Oxygen Fluxes for Various Dual-Phase Composite Membranes Reported in the Literature Membrane Material

Temperature (
C)

Oxygen Flux (mol s 1 m 2)

Membrane Configuration

Thickness (mm)

((ZrO2)0.94(Y2O3)0.06)0.5–Pd0.5 [120]

1100

1.5  10 2–1.6  10 2

Disk

0.8

((ZrO2)0.94(Y2O3)0.06)0.5–Pt0.5 [120]

1100

1.4  10 2

Disk

0.8

(ESB)0.6–(Ag)0.4 [75] (ESB¼(Bi2O3)0.75(Er2O3)0.25)

750–850

8.5  10 4–3.1  10 3

Disk

0.23–1.6

(ESB)0.6–(Au)0.4 [124] (ESB¼(Bi2O3)0.75(Er2O3)0.25)

650–850

2.8  10 4–5  10 4

Disk

1

Zr0.84Y0.16O1.92 [125]

850–1050

4  10 4–2  10 3

Tube

0.16

Bi1.6Y0.4O3–Ag [126]

650–950

2  10 3–6  10 3

Disk

1.44

(Y2O3)0.08(ZrO2)0.92–Ni [127]

950–1000

3  10 3–4.2  10 3

Disk

0.7

Disk

1.52

Disk

2.82

Tube

1.25

Disk

0.3

950–1000 950–1000 (SrFeO3s)0.7(SrAl2O4)0.3 [128] Ce0.8Sm0.2O2s–La0.8Sr0.2CrO3d [129]

800–900 850–950 850–1000 850–1000 850–1000

SrSc0.2Co0.8O3s–Sm0.5Sr0.5CoO3d [130]

700–900

3

3

–3  10

2.5  10

3

3

4

3

–1.5  10

1.1  10

1.6  10

–2.5  10

4

5  10

3

–1.6  10

4

–2  10

4  10

Disk

0.6

4

3

Disk

1

4

4

Disk

1.7

Disk

0.85

–1.3  10

3.2  10

–6.3  10

1.3  10

3

2  10

3

3

–7.5  10

Ce0.8Gd0.2O0.19–Gd0.2Sr0.8FeO3d [131]

800–1000

1.1  10 3–9  10 3 4

4

Disk

0.5

(50%wt)Ce0.8Gd0.2O2d–(50%wt)La0.8Sr0.2Fe0.8Co0.2O3d [121]

750–950

2.0  10

–6  10

Disk

1.0

(50%wt)Ce0.8Gd0.2O2s–(50%wt)La0.7Sr0.3FMnCo0.2O3d [121]

750–950

1  10 4–6  10 4

Disk

1.0

((ZrO2)0.94(Y2O3)0.06)0.5–Pd0.5 [120]

1100

1.5  10 2–1.6  10 2

Disk

0.8

To convert to molar fluxes, all volumetric fluxes reported in the literature were assumed to be at T ¼ 25
C and P ¼ 1 atm unless otherwise stated in the reference.

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flux. For membranes limited by bulk diffusion rates, reducing the thickness of the membrane can increase the oxygen flux significantly. If both surface exchange and bulk diffusion limit oxygen transport, surface modification and membrane thickness reduction can both be applied to increase oxygen flux. Recent studies have shown that surface modification and/or thinner membranes can both increase oxygen transport through MIEC ceramic membranes.

Surface Modifications When oxygen transport is limited by the rate of oxygen exchange on the surface of the membrane, increasing surface area can be one strategy to increase the oxygen flux. Polishing the membrane surface with different media can lead to the roughened surface which increases the surface-to-volume ratio. Experiments on a La0.1Sr0.9Co0.9Fe0.1O3d membrane have shown that an increase in the oxygen flux can be accomplished by using this simple method [132]. Another effective means to increase the surface area is coating a porous layer on the dense oxygen-MIEC membranes. The porous coating material can be the same as the membrane material or it can be another oxygen-MIEC material [133–135]. Compared to the simple roughening method, coating a porous layer on the dense membrane has been shown to be more effective and can lead to greater oxygen flux enhancement. Coating an oxygen dissociation catalyst on the membrane surface can facilitate the adsorption and dissociation of oxygen and increase the oxygen flux. Researchers showed that La0.6Sr0.4Co0.2Fe0.8O3d membranes [136] coated with a silver catalyst on the air side surface of the membrane improved the oxygen permeability of the membrane significantly. Compared to unmodified membranes, oxygen flux through silver-coated membranes at 1000
C in an air:helium gradient increased by a factor of 2.5. Similarly, studies of the effect of both a platinum pattern [115] and a very thin platinum film (< 1 nm) [137] showed that the deposited catalyst substantially enhanced membrane oxygen flux. As shown in Figure 11.2, an SFC membrane with a pattern of 20-nm-thick platinum disks exhibited an oxygen flux two times higher at 800
C than an identical membrane without the platinum pattern [115]. Studies with Sr0.97Ti0.6Fe0.4O3d membranes also demonstrated an increase in oxygen flux following the application of a platinum-based oxygen dissociation catalyst to the membrane surface [138]. In contrast to the studies on the SFC and Sr0.97Ti0.6Fe0.4O3d membranes, platinum and silver coatings on the air side surfaces of 1.8-mm-thick BSCF membranes did not result in an improvement in oxygen flux. Additional studies showed that oxygen transport through BSCF membranes greater than 1.8 mm in thickness is not limited by surface exchange on the air side of the membrane but is limited by bulk diffusion of oxygen ions through the membrane.

11

Membrane oxygen flux (mol m-2 s-1)

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1 ⫻ 10-3 8 ⫻ 10-4

Plain membrane Pt-patterned membrane

6 ⫻ 10-4 4 ⫻ 10-4 2 ⫻ 10-4 0 450

550

650 Temperature (°C)

750

850

FIGURE 11.2 Oxygen flux results for plain and Pt-patterned membranes (Pt pattern installed on oxygen supply side).

Membrane Thickness Reduction Under bulk diffusion control, oxygen flux across an oxygen-MIEC membrane can be correlated with the inverse of membrane thickness as described by the Wagner equation (Equation 11.2) [139]: ð lnP0 O2 RT si se d lnP; ð11:2Þ JO 2 ¼ 2 00 16F L lnPO si þ se 2

where R is the gas constant (J mol 1 K 1), T is the temperature (K), F is the Faraday constant (C mol 1), L is the thickness of membrane (m), si is the oxygen ionic conductivity (S m 1), se is the electronic conductivity (S m 1), and P0O2 and 00 PO2 are the oxygen supply side and oxygen permeate side partial pressures, respectively (Pa). In theory, then, reducing membrane thickness can increase oxygen flux when oxygen transport is limited by bulk diffusion. However, preparing a thin or ultra-thin membrane is not always feasible, and the thickness of the membrane can have a significant role in the mechanical stability of the membrane under operating conditions. For this reason, a substantial amount of research has focused on asymmetric membrane fabrication. Asymmetric membranes contain a thin film supported on a porous substrate. The initial studies in the literature concentrated on coating thin layers of oxygen-MIEC materials on conventional porous substrates such as Al2O3 because of the low cost and availability of the porous substrate [140]. However, to avoid the physical and chemical incompatibility between the thin layer and substrate, fabricating a thin layer and a porous substrate from the same materials is desirable. To date, many fabrication methods such as tape-casting [141], screen printing [142], slurry-coating [143], acid etching [144], and electrochemical vapor deposition (EVD) [84] have been used to prepare asymmetric oxygen-MIEC

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membranes. For example, recent work on a BSCF membrane demonstrated that the oxygen flux of a BSCF asymmetric membrane was effectively enhanced as the thickness of the dense BSCF layer in the asymmetric membrane decreased (Figure 11.3) [111]. This increase in oxygen flux was attributed to the reduction in the thickness of the dense BSCF layer in the asymmetric membrane, leading to a reduction in diffusion resistance for oxygen transport through the membrane. In these studies, the thickness of the porous substrates for each membrane was kept the same to isolate the effect of the dense layer thickness. The observed oxygen fluxes in Figure 11.3 are much lower than theoretical values predicted by the Wagner equation. Additionally, the deviation of the measured oxygen fluxes from theory increases as the thickness of the thin layer decreases. Two factors could be responsible for the lower-thanpredicted oxygen fluxes. The first is that the porous support could still exert some resistance on oxygen transport. The second is that, as the thickness of the thin layer decreases, the influence of surface exchange could increase, causing a transition to mixed control scenario in which both bulk diffusion and surface exchange limit oxygen transport. In the mixed control scenario, oxygen flux is not inversely proportional to thickness and, thus, deviations from the theoretical oxygen flux values predicted by Wagner’s equation would be expected. New techniques for fabricating ultra-thin dense layers on thin yet robust porous supports continue to be explored. Novel hollow fiber oxygen-permeable ceramic membranes have been reported which have several advantages over membranes prepared using traditional fabrication techniques [144–148]. Compared to conventional tubular or disk-shaped membranes, hollow fiber membranes have a higher surface-to-volume ratio. The thinner membrane wall

Oxygen flux (mol m-2 s-1)

8 ⫻ 10-3

6 ⫻ 10-3

4 ⫻ 10-3

Asymmetric membrane Dense membrane

2 ⫻ 10-3 Theorectical oxygen fluxes 0 0

2 4 Inverse of thin layer thickness (mm−1)

6

FIGURE 11.3 The effect of the thicknesses of the dense layers of BSCF asymmetric membranes to the oxygen flux of BSCF membranes (T ¼ 800
C).

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255

and asymmetric structure of hollow fiber membranes are favorable to oxygen permeation, especially for materials limited by bulk diffusion. Finally, the hollow fiber membrane configuration has been suggested to be more suitable for large-scale fabrication and for utilization in membrane reactors. Hollow fiber membranes can be prepared using a phase inversion spinning and sintering method [136,149]. The typical morphology of a hollow fiber membrane has finger-like structures near both the inner and outer membrane walls and a sponge-like thin layer at the center of the membrane wall. This asymmetric structure can be attributed to rapid precipitation at both the inner and outer walls close to the coagulant, which results in short finger-like pores, while slow precipitation at the center of the wall leads to the sponge-like structure. Compared to conventional tubular or disk-shaped membranes, a hollow fiber membrane with the same composition shows higher oxygen fluxes. For example, Tan et al. [145] reported that at 900
C, the oxygen flux of a La0.6Sr0.4Co0.2Fe0.8O3d hollow fiber membrane fiber is about 5.5  10 3 mol m 2 s 1 in an air:argon gradient, which is about 3.5 times greater than a conventional tubular La0.6Sr0.4Co0.2Fe0.8O3d membrane. Oxygen flux increases with hollow fiber membranes have been attributed to the significant reductions in wall thickness with the hollow fiber membranes and the attendant reduction in bulk diffusion limitations [136,146].

MIEC MEMBRANES FOR SYNTHESIS GAS PRODUCTION Reforming and partial oxidation of hydrocarbons to produce mixtures of hydrogen and carbon monoxide known as synthesis gas (or “syngas”) have been widely studied in MIEC membrane reactors [29,35,43,110,150–160] because of the potential for an oxygen-MIEC membrane to provide cheap, safe, and distributed oxygen. From a membrane performance perspective, motivations for incorporating MIEC membranes into hydrocarbon conversion reactors include (1) the high reaction temperature of the production of the synthesis gas allows effective membrane oxygen transport without a requirement for additional energy and (2) the strongly reducing atmosphere created by the reaction products provides a large oxygen potential gradient to facilitate oxygen transport through a membrane. Figure 11.4 illustrates a simple scheme for syngas production from methane in an oxygen-permeable membrane reactor. This figure shows the membrane in a disk configuration, but it should be noted that tubular configurations are also widely studied in the literature and are much more likely to be utilized in an industrial process.

Synthesis Gas Production Overview Syngas is an important industrial feedstock for hydroformylation and Fischer–Tropsch processes to make chemicals and fuels. Generally speaking,

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Inorganic, Polymeric and Composite Membranes

syngas is produced by either partial oxidation or reforming of hydrocarbons. Among the hydrocarbons typically studied, methane (CH4) has received the most attention because it is both the simplest alkane and the main component of natural gas [43,150,155,159,161]. The three primary reactions used to produce synthesis gas from methane are shown in Table 11.4 [27,152]. POM yields a H2:CO product ratio of 2:1 that is ideal for Fischer–Tropsch and direct methanol syntheses of liquid fuels, which are also referred to as gasto-liquids (GTL) processes. One disadvantage of partial oxidation in general is its large oxygen requirement. Pure gas-phase oxygen is an expensive feedstock and the nitrogen in air precludes its direct use. About one-third of a POM synthesis gas facility’s operating costs and up to 45% of its capital costs can come from the air separation unit alone [39,162]. Any reduction in the cost of supplying oxygen to such a system would directly lower its production costs and thus increase its competitiveness. For example, Dyer et al. [10] estimated in 2000 that a 25% reduction in the cost of current GTL technology including oxygen separation would make GTL products competitive with oil at $20 per barrel, which could substantially increase GTL production. FIGURE 11.4 Membrane reactor scheme for the methane conversion to syngas using an MIEC ceramic membrane.

CH4

H2, CO, CO2, H2O

Catalyst

Ceramic membrane

e−

O2−

Air

TABLE 11.4 Comparison of Oxygen Fluxes for Various Dual-Phase Composite Membranes Reported in the Literature Reaction

Chemical Equation

DH298 K∘(kJ mol 1)

Product H2: CO Ratio

Partial oxidation

CH4 þ 12 O2 ! CO þ 2H2

36

2

Steam reforming

CH4 þ H2O ! CO þ 3H2

þ206

3

CO2 reforming

CH4 þ CO2 ! 2CO þ 2H2

þ247

1

To convert to molar fluxes, all volumetric fluxes reported in the literature were assumed to be at T ¼ 25
C and P ¼ 1 atm unless otherwise stated in the reference.

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257

The most common method for producing synthesis gas in industry is steam reforming [163]. The high H2:CO ratio of the products is ideal for ammonia and urea production and also makes this the most desirable reaction for hydrogen production from hydrocarbons. However, steam reforming is highly endothermic, which is an expensive attribute with target reaction temperatures of at least 750
C for steam reforming of methane. In addition to the energy required to maintain the reaction temperature, a high steam to hydrocarbon feed ratio is typically required to minimize catalyst deactivation [27]. The CO2 reforming reaction utilizes methane and carbon dioxide, both of which are key greenhouse gases, to produce synthesis gas, and has been suggested as an environmentally benign strategy for reusing CO2 produced by other processes. The lower H2:CO ratio (1:1 or less) of the CO2 reforming reaction products also provides an opportunity to adjust the composition of synthesis gas to meet the needs of downstream processes such as hydroformylation that require lower H2:CO ratios than those produced by steam reforming and partial oxidation. Moreover, CO2 reforming can utilize the CO2 that naturally exists in natural gas. This removes the need to separate the CO2 from natural gas, which is particularly beneficial for oilfield gas sources that can contain large amounts of CO2. The major disadvantages for the practical application of CO2 reforming include the potential for rapid catalyst deactivation from carbon deposition, high reaction temperatures (beyond 1073 K), and unfavorable energetics. Like steam reforming, CO2 reforming is a highly endothermic reaction. A more recent approach to producing H2:CO ratios between one and two is a combination of partial oxidation and CO2 reforming known as combined reforming. Combined reforming has been explored with the intention of decreasing the endothermicity and poor catalyst life associated with straight CO2 reforming [164–176]. It also offers the potential for “tuning” the H2:CO ratio via manipulation of the ratio of O2 and CO2 in the reactor feed. Adding oxygen to promote catalyst performance and decrease the energy requirements of CO2 reforming could bring this potentially valuable syngas production reaction closer to mainstream industrial application. However, combined reforming requires pure oxygen as a feedstock, and as previously stated, an energy-intensive air separation unit represents a substantial capital and operating cost. Oxygen-MIEC membranes offer the potential to replace separated oxygen gas in future partial oxidation and combined reforming processes, and thereby reduce synthesis gas process operating costs.

Benefits of MIEC Membranes for Synthesis Gas Production In addition to the potential economic benefits, membrane-supplied oxygen incorporates several significant environmental and safety benefits over gasphase oxygen supplies for synthesis gas operations (1) a substantial reduction

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in process energy consumption; (2) safer oxygen (i.e., no hotspot or flammability issues); (3) inherently distributed oxygen introduction, which can increase the selectivity of oxidation reactions, produce a more uniform and predictable reactor temperature profile, and possibly even reduce overall oxygen consumption; and (4) a reduction in homogeneous thermochemical reactions involving gas-phase oxygen which can produce soot-forming precursors. The potential advantages of membrane oxygen have led to numerous investigations of oxygen-MIEC membrane reactors for synthesis gas production. Membrane material composition and the effect of composition and structure on oxygen flux have a significant impact on the activity and stability of oxygenMIEC membranes in synthesis gas production environments. Several requirements must be met by industrially feasible membranes, including (1) high and steady oxygen flux during reaction, (2) considerable long-term mechanical and thermal stability under a reducing environment, and (3) inexpensive starting materials and fabrication methods [31].

Overview of Work to Date The majority of the work published to date on the use of oxygen-MIEC membranes for the production of synthesis gas has focused on the POM reaction. While a comprehensive review of all POM studies on oxygen-MIEC membranes is beyond the scope of this chapter, a brief review of selected membranes with high fluxes and/or high stability is included in this section. Table 11.5 provides a summary of the oxygen fluxes observed for the membrane materials discussed under air:inert and air:POM gradients. Early work utilizing MIEC membranes in POM reactions sought to take advantage of the high oxygen fluxes exhibited by SCF membranes [39,89]. However, when these membranes were exposed to reaction environments at high temperature, they fractured quickly. Changing the composition to the mixed-phase SFC resulted in membranes that were more stable in reducing environments. However, the oxygen flux of the SFC material is significantly less than that of the perovskite SCF, and SFC is thus not a viable candidate for POM applications [39,50,51,55,116]. Other studies with La0.6Sr0.4Co0.8Fe0.2O3d [35], La0.2Ba0.8Fe0.8Co0.2O3d [42], and Ba0.5Sr0.5Co0.8Fe0.2O3d [43,177] have found that doping higher valence metal ions such as La2þ and Ba2þ into the A site of SCF perovskites can achieve oxygen fluxes in air:inert gradients that are similar to the SCF membranes, but with significantly improved stability under reaction conditions. For example, Ba0.5Sr0.5Co0.8Fe0.2O3d maintained an oxygen flux of 7.8  10 2 mol m 2 s 1 at 850
C for over 1000 h in an air:helium environment and achieved > 97% methane conversion at 875
C for over 500 h [43]. The suggested reaction mechanism for this work was combustion of a portion of the CH4 to CO2 and H2O by membrane oxygen followed by CO2 and steam reforming of the remaining CH4 to form synthesis gas.

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TABLE 11.5 Oxygen Flux Results Under Reaction Conditions for Various Membrane Materials

Membrane Material

Oxygen Flux Air: Inert Gradient (mol s 1 m 2)

Oxygen Flux Air:POM Gradient (mol s 1 m 2)

Thickness (mm)

SrCo0.8Fe0.2O3d [25,39,88]

7.4  10 4–6.7  10 3 Fractured quickly [25,39]

1

SrFeCo0.5O3d [39,50,51,55,116]

7.4  10 6–9.7  10 4 2.2  10 3 [39]

1

La0.6Sr0.4Co0.8Fe0.2O3d [35]

7.4  10 4 (900
C)

7.4  10 4

1

La0.2Ba0.8Fe0.8Co0.2O3d [42]

5.9  10 3

3.0  10 2

1

Ba0.5Sr0.5Co0.8Fe0.2O3d [43,177]

7.8  10 3–1.0  10 2 7.1  10 2 [43]

2.2  10 2 [39]

1.8

5.8  10 2 [177]

La0.8Sr0.2Co0.1Fe0.8Cr0.1O3d [5] Not reported

9.9  10 2

Not reported

BaCo0.4Fe0.4Zr0.2O3d [25]

4.1  10 3

3.8  10 2

1

BaCo0:7Fe0:2Nb0:1O3d [22]

Not reported

>0.14

1

BaCo0.7Fe0.2Ta0.1O3d [23] BaCe0.1Co0.4Fe0.5O3d [156]

2

1.4  10

2

< 1.0  10

0.11–0.12 2

6.5  10

0.6–0.7 1

To convert to molar fluxes, all volumetric fluxes reported in the literature were assumed to be at T ¼ 25
C and P ¼ 1 atm unless otherwise stated in the reference.

More recent investigations have focused on doping higher valance metal ions into both the A and B sites. The majority of the studies have used La2þ and Ba2þ in the A site with Cr7þ, Nb5þ, Ta5þ, and Zr4þ in the B site. La0.8Sr0.2Co0.1Fe0.8Cr0.1O3d membranes exhibited oxygen fluxes of 9.9  10 2 mol m 2 s 1 in an air:syngas gradient and were stable for the POM reaction at 900
C for 340 h [5]. The addition of Zr4þ was observed to increase stability further, with a BaCo0.4Fe0.4Zr0.2O3d membrane operating for over 2200 h with 96–98% methane conversion and 98–99% selectivity of carbon to CO. The flux during reaction was reported to be 3.7  10 2–4.0  10 2 mol m 2 s 1 at 850
C [25]. Substituting Nb5þ [22] and Ta5þ [23] in the B site resulted in membranes that were stable under POM reaction conditions for 300 h and 400 h, respectively, with unusually high oxygen fluxes. At 900
C, the BaCo0.7Fe0.2Ta0.1O3d membrane achieved > 99% CH4 conversion with 94% selectivity to CO and a flux of 0.11 mol m 2 s 1 during the reaction, while fluxes > 0.14 mol m 2 s 1 were

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Inorganic, Polymeric and Composite Membranes

observed under reaction conditions for the Nb-doped membrane. These fluxes are some of the highest reported in the literature to date.

Effect of Reaction Temperature on Membrane Performance Reactor temperature is a key factor for syngas production with oxygen-MIEC membranes because both catalyst activity and membrane oxygen flux are strongly dependent on temperature. Differences in reaction conditions among studies prevent a direct comparison of the effect of reaction temperature on oxygen flux for specific membrane compositions. However, general correlations between oxygen flux, methane conversion, and temperature are evident for the majority of the studies. Figure 11.5 depicts the effect of temperature on methane conversion and oxygen flux for the POM reaction using a nickel-based catalyst and a BaCo0.7Fe0.2Ta0.1O3d membrane [23]. Increasing the reaction temperature resulted in a significant increase in both oxygen flux and methane conversion, with conversion ultimately leveling off near 90% at 950
C. These same trends of increasing oxygen flux and methane conversion with increasing temperature were observed with other catalysts and membranes. Examples include the use of a LiLaNiOx/g-Al2O3 catalyst in conjunction with BaCo0.4Fe0.4Zr0.2O3d [25] and BaCe0.1Co0.4Fe0.5O3d [156] membranes. Tong et al. ascribed the increased oxygen flux to an enhancement of both the diffusion rate of lattice oxygen vacancies in the bulk membrane material and the surface exchange rate. One interesting point in these studies is the trend in CO selectivity. As the temperature, flux, and conversion increase, product selectivity to CO decreases.

O2 flux (ml min-1 cm-2)

19

Oxygen flux Methane conversion

90

17 15

80

13

70

11 60

CH4 conversion (%)

100

21

9 7 750

800

900 850 Temperature (°C)

950

50 1000

FIGURE 11.5 O2 flux and CH4 conversion as a function of temperature for BaCo0.7Fe0.2Ta0.1O3d during partial oxidation of methane. Experimental conditions: air: 200 ml min 1; 48.0% helium diluted methane: 50 ml min 1; and membrane thickness: 0.7 mm [23]. (To convert to molar fluxes, all volumetric fluxes were assumed to be at T ¼ 25
C and P ¼ 1 atm.)

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This drop in CO selectivity is due to the ratio of methane and oxygen in the reactor. When CH4 is in excess (i.e., at relatively low temperatures), any increase in oxygen flux results in a direct increase in methane conversion and CO selectivity remains high. As the temperature is increased and additional membrane oxygen becomes available, the methane excess diminishes, which can lead to deeper oxidation of the reaction products and a reduction in CO selectivity. The exact temperature at which CO selectivity begins to decline and the rate at which CO selectivity decreases depend on the membrane material and also on other factors such as the catalyst used and the composition of the reactor feed.

Effect of Reaction Environment on Membrane Oxygen Flux In general, regardless of membrane composition, oxygen fluxes for MIEC membranes under POM reaction conditions are higher than fluxes observed under an air:inert gradient. The increase in oxygen flux during the POM reaction is attributed to the rapid consumption of oxygen on the reaction side of the membrane, which leads to a vanishingly low oxygen partial pressure in the gasphase environment and a low oxygen potential in the membrane surface exposed to the reaction environment. The lower oxygen potential in the reaction surface than in the low oxygen surface under an inert atmosphere produces a larger oxygen potential gradient across the membrane, which ultimately leads to higher flux. CH4, CO, and H2 are the most prominent reducing agents in the synthesis gas production process. Studies have been performed in different gas environments to investigate the effect of a reducing gas on membrane oxygen flux. The order of increase in oxygen flux in the presence of different reducing gases was reported to be H2 > CO > CH4 [154] and this order has been suggested to correlate with the tendency of the gases to react with surface oxygen ions. Previous studies have shown that carbon monoxide and hydrogen readily react with lattice oxygen from the membrane to produce CO2 and water while CH4, even at high temperatures, exhibits very little activity with the membrane surface in the absence of a reforming catalyst [153]. These studies have also shown that the Ba0.5Sr0.5Co0.8Fe0.2O3d membrane material reacts more readily with hydrogen than with CO, so the surface reaction results correlate with the oxygen flux increase results. It is interesting to note that some researchers have seen a shift in the limiting factor for oxygen transport by changing the permeate side environment. By varying membrane thickness, Luo et al. [23] found that oxygen flux through BaCo0.7Fe0.2Ta0.1O3d membranes with thicknesses between 0.6 and 1.2 mm was limited by bulk diffusion in an air:helium gradient at 900
C. However, when membrane thickness was varied with an air:POM gradient, no significant change in oxygen flux was observed. This observation suggests that the BaCo0.7Fe0.2Ta0.1O3d membrane switched from a bulk diffusion-limited mode

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to a surface exchange-limited mode when the permeate side environment was switched from inert gas to POM/synthesis gas. The effect of changing the permeate side environment during the POM reaction by varying the flow rate of a pure methane feed has been investigated using a BaCe0.1Co0.4Fe0.5O3d membrane. Figure 11.6 shows the effect of feed flow rate on oxygen flux, methane conversion, and CO selectivity, where a higher methane feed rate is assumed to correspond to a higher methane concentration in the reactor. Similar to the temperature effect results discussed in the earlier section, oxygen flux increases significantly and then plateaus at approximately 8.2  10 2 mol m 2 s 1 as methane flow rate continues to increase. However, unlike in the temperature results, as the methane flow rate is increased, methane conversion decreases while CO selectivity increases. The decrease in methane conversion and corresponding increase in selectivity have been ascribed to the change in the methane:oxygen ratio in the reactor [25,156]. When the methane flow rate is low, the amount of available oxygen is sufficient to react with almost all of the methane. However, a decrease in the methane:oxygen ratio favors deeper oxidation of methane and thus produces lower CO selectivity. As the feed flow rate of CH4 is increased, so does the methane: oxygen ratio in the reactor. Eventually, the reaction becomes limited by the oxygen supplied through the membrane. While oxygen flux and CO selectivity remain constant as the flow rate is increased further, methane conversion continues to decrease. These trends have also been observed with other membrane materials and with diluted methane feeds, suggesting that a general relationship between reactant feed and oxygen flux, hydrocarbon conversion, and reaction selectivity may exist regardless of the membrane material. 120

20

100

16 14

80

12 10

60

8 40

6

Oxygen flux Conversion CO selectivity

4 2 0 5

10 15 20 25 Methane flow rate (mI min-1)

20

Conversion/selectivity (%)

O2 flux (mI min-1 cm-2)

18

0 30

FIGURE 11.6 O2 flux, CH4 conversion, and CO selectivity as a function of CH4 flow rate for BaCe0.1Co0.4Fe0.5O3d during partial oxidation of methane. Experimental conditions: air flow rate: 250 ml min 1; thickness of the membrane: 1.0 mm; temperature: 875
C [156]. (To convert to molar fluxes, all volumetric fluxes were assumed to be at T ¼ 25
C and P ¼ 1 atm.)

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The effect of varying the air flow rate on the oxygen source side has also been investigated. Studies showed that at low air flow rates, oxygen flux increases with increasing air flow rate. These results indicate that oxygen transport from the gas phase to the membrane surface is a rate-controlling factor when air flow rate is low. However, as the air flow rate continues to increase, oxygen flux again reaches a plateau. The flow rate at which flux levels off depends on the membrane material, but in general was found to be in the range of 200–250 ml min 1 [23,156]. Methane conversion and CO selectivity were also found to be affected by air flow rate. However, the change in these values was on the order of 1–2% whereas a much larger change was observed when methane flow rate was varied, as discussed earlier. Thus, air flow rate has a much smaller impact on oxygen flux, methane conversion, and selectivity than the composition of the reactor contents. An isothermal membrane exposed simultaneously to two different oxygen partial pressures experiences nonuniform chemical expansion (or contraction) to a degree that depends on the material phases present and the ratio of the oxygen partial pressures [54,59]. The oxygen potential gradient across a partial oxidation reactor membrane is much more extreme than the gradient during a typical oxygen flux test. In a POM membrane reactor, the membrane’s permeate side can be exposed to an oxygen partial pressure of between 10 17 and 10 30 atm at 850
C [21,24,40], compared to a permeate side oxygen partial pressure of 10 3–10 5 atm in an oxygen flux test under an air:inert gradient [24,55,88]. Considering the much larger chemical expansion gradient within a membrane in a POM reactor, among other differences, it is to be expected that fewer viable membranes have been proposed for membrane reactor than for oxygen separation applications. The very properties that give some materials high oxygen diffusivity can make them unfit for use in methane conversion applications [6].

CONCLUSIONS Oxygen-MIEC membranes have been extensively studied for over three decades and significant advances have been, and continue to be, made. Reported oxygen fluxes are approaching levels that have commercial application, and the new oxygen-MIEC materials being tested have also been reported to exhibit greatly improved mechanical stability compared with early membrane material candidates. However, substantial work is still needed to develop and thoroughly evaluate these more robust, higher flux membrane materials. The largest practical challenge facing designers of oxygen-conducting ceramic membranes is clearly the need to simultaneously maximize oxygen conduction—which is expected to correlate positively with bulk oxygen absorption/desorption capacity and the resulting lattice contractions/expansions—and membrane mechanical stability, which is expected to correlate negatively with these characteristics. The commercial potential of new membrane materials,

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based on both flux and mechanical stability, can only be evaluated properly under reaction conditions. Membrane stability must also be assessed under relevant industrial constraints, both economic and practical. Catalyst and membrane systems must be inexpensive to fabricate and must demonstrate long-term, low maintenance operation. Also, test reactor feed composition must represent actual commercial supplies. For example, a majority of the reaction studies reported in the literature use diluted methane as a reactor feed which is not consistent with industrial operations. Industrial hydrocarbon feedstocks will also contain impurities such as CO2 that can have a significant impact on oxygen flux, which means that the true commercial viability of new MIEC membrane reactor candidate materials must be evaluated in the presence of these impurities. Finally, reactor system design and engineering will become even more important as membranes become thinner in the quest for increased oxygen flux. Not only does a reduction in membrane thickness have an impact on the mechanical strength and possible configuration of a membrane reactor, but it may also amplify the effect of any catalyst–membrane interactions, as thinner membranes bring a greater proportion of the total membrane material in contact with the catalyst. Understanding and predicting long-term catalyst–membrane interactions is a critical research area with any membrane material or configuration, but thinner membranes make it even more urgent. Future work will benefit from a combined knowledge of material science, reaction engineering, and catalyst design.

ACKNOWLEDGMENTS Financial support for this work was provided by the Office of Naval Research (N00014-03-10601) and the US Department of Transportation Research Innovative Technology Administration (DTOS59-06-G-0047).

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