Doctoral Dissertation

Tailoring the microporous structure of a ceramic composite
for development of molecular sieving membranes(セラミック複合物のマイクロポーラス制御による分子ふるい膜の開発)

Sulaiman Oladipo Lawal

Department of Chemical Engineering,
Graduate School of Engineering, Hiroshima University(広島大学 大学院工学研究科 化学工学専攻)

September 2021

Hiroshima University Doctoral Dissertation

Abstract of this dissertation
In this thesis, organic coordination/complexation ligands have been applied to themodification of a SiO2-ZrO2 ceramic composite. The various sections seek to deepentheunderstandingandaccessfurther

opportunities

offeredbyorganic

coordination/complexation ligands modification of ceramic composite membranes.
First, the chemical modification of SiO2-ZrO2 via the sol-gel method was
successfully accomplished by chelating acetylacetonate (acac-) with zirconium
tetrabutoxide (ZrTB) prior to co-hydrolysis and condensation with tetraethoxysilane(TEOS), which was then used to fabricate a gas separation layer. The acac--modified
SiO2-ZrO2-derived membrane showed a high H2 permeance of 9.9 x 10 -7 mol m-2 s-1
Pa-1 with a H2/SF6 permeance ratio of 7,600, which was better than the results whenusing a pure SiO2-ZrO2-derived membrane (H2 permeance: 1.4 x 10 -6 mol m-2 s-1 Pa-1,
H2/SF6 permeance ratio: 11).
In the next phase, the microstructural control of ligand-modified SiO2-ZrO2
composite membranes was attempted. Hybrid materials with crosslinking between theorganic moieties were made by utilizing network forming Si precursor and chelating
ligands. The effects of using a network-forming and network-modifying ligand werestudied. Gas permeation experiments showed that the organic crosslink-derived
membranes possess molecular sieving properties (H2 permeance: ≈10-7 mol m-2 s-1 Pa1; H2/N2 selectivity: ≈20-30) that are superior to those of non-crosslink-derived

membranes (H2 permeance: ≈10-7 mol m-2 s-1 Pa-1; H2/N2 selectivity: ≈7-11).
Furthermore, the preparation of a carbonized SiO2-ZrO2 composite membrane wasstudied. The sol synthesis of the precursor SiO2-ZrO2-acetylacetonate composite was
optimized by studying the effect of water/alkoxide molar ratio and relative Si/Zr
atomic ratio on the final carbon-SiO2-ZrO2 properties. A combination of SiO2 andZrO2 in a ratio of 9/1 in a carbon-SiO2-ZrO2 membrane showed an H2 permeance of
16 x 10-7 mol m-2 s-1 Pa-1 over a similarly pore-distributed unmodified SiO2-ZrO2
membrane with an H2 permeance of 3 x 10-7 mol m-2 s-1 Pa-1. More interestingly, acombination of SiO2 and ZrO2 in a ratio of 5/5 displayed a unique pressure-induced
CO2 flow switching behaviour.

Hiroshima University Doctoral Dissertation

The various modification techniques offered more insights for molecularseparations.

Hiroshima University Doctoral Dissertation

TABLE OF CONTENTS
Page
Chapter 1
General introduction
1.1

A brief background on membranes

1

1.1.1 Synthetic membranes

1

1.1.2 Gas separation membranes

4

1.1.3 Gas separation mechanisms through membranes

4

1.1.3.1

Bulk fluid flow5

1.1.3.2

Knudsen flow5

1.1.3.3

Surface diffusion6

1.1.3.4

Gas translational diffusion6

1.1.3.5

Solid-state diffusion6

1.1.4 Classification of gas separation membranes

7

1.2

1.1.4.1

Classification by membrane material

8

1.1.4.2

Classification by physical configuration10

1.1.4.3

Classification by material composition

11

Silica and composite matrix ceramic membranes

15

1.2.1 State of the art of ceramic composite materials and membranes

18

1.2.2 Organic coordination ligands: a broad field of opportunities

20

1.2.2.1

Coordination ligands20

1.2.2.2

Organic coordination ligands as molecular engineering agents
21

1.2.3 Ligand-modified ceramic composite membranes

23

1.3

The scope and outline of this thesis

25

Cited literature

29

Chapter 2
Development of a molecular sieving membrane from an acetylacetonate-modified
silica-zirconia ceramic composite
2.1

Introduction

43

2.2

Experimental

46iHiroshima University Doctoral Dissertation

2.3.1 Preparation of SiO2-ZrO2-acac2 sol and membrane fabrication

46

2.3.2 Characterizations of the SiO2-ZrO2-acac2, the SiO2-ZrO2 sol, the gel,and the membranes

48

2.3.3 Single-gas permeation measurements

48

2.3

Results and discussion

49

2.3.1 Sol-gel characterization

49

2.3.1.1

Chemical modification of a ZrTB precursor

49

2.3.1.2

Thermal analysis of acac- ligands in SiO2-ZrO2-acac2 gels51

2.3.2 Effect of acetylacetonate ligands on the pore size controllability of
SiO2-ZrO2 (1:1) membranes

54

2.3.3 Heat-treatment effect on pore size and gas permeation properties of
SiO2-ZrO2-acac2 membranes

58

2.4

Conclusions

66

Cited literature

67

Chapter 3
Microstructure engineering of ligand-modified silica-zirconia membranes
derived from network-forming precursors
3.1

Introduction

72

3.2

Experimental

75

3.2.1 Preparation of MAPTMS-SZ-LIGAND sols, powders, and films
3.2.1.1

76

Fabrication of MAPTMS-SZ-AAA and MAPTMS-SZ-ACAC

membranes

77

3.2.2 Evaluation of gas permeation through membranes

78

3.3

78

Results and discussion

3.3.1 Evolution of organic cross-linking units via thermal radical
polymerization
3.3.1.1

78

Facile control of particle size distribution and heterogeneity in

sols
3.3.1.2

78
Microstructural evolution of organically cross-linked materials
80

3.3.2 Thermal stabilities and decomposition mechanisms: effect of organic
crosslinking

85
ii

Hiroshima University Doctoral Dissertation

3.3.3 Membranes fabricated from cross-linked hybrids and gas permeation
analysis

89

3.3.3.1

Morphologies of supported membranes

89

3.3.3.2

Gas sieving properties at 200 oC and membrane microstructures
91

3.3.3.3
3.4

Analysis of microporous structures via the gas permeance ratio

and activation energy

96

Conclusions

101

Cited literature

102

Chapter 4
Enhanced H2-permselectivity in a carbon-silica-zirconia membrane derived froman acetylacetonate-ceramic composite pyrolysis
4.1

Introduction

108

4.2

Experimental

111

4.2.1 Materials preparation by sol-gel

111

4.2.2 Preparation of membrane separation layers

112

4.2.3 Material characterizations

112

4.3

Results and discussion

113

4.3.1 Control and optimization of ligand chelation degree via hydrolysis 113
4.3.1.1

Characterization

of

sols,

films,

andpowders

beforecarbonization

113

4.3.1.2

119

Characterization after carbonization

4.3.2 Microstructuraltransformation

between

SiO2-ZrO2-ACAC

Carbon-SiO2-ZrO2: effect of Si/Zr ratio

and
121

4.3.3 Carbonization-enhanced H2-permselectivity in a SiO2-ZrO2 composite
membrane

128

4.4Conclusions

132

Cited literature

133

Chapter 5CO2 flow-switching behaviour in a pyrolysis-derived carbon-silica-zirconia
membrane promises a versatile high-temperature H2/CO2 separation
5.1

Introduction

138
iii

Hiroshima University Doctoral Dissertation

5.2

Experimental

141

5.2.1 Preparation and pyrolysis of SiO2-ZrO2-acac powder and C-SiO2-ZrO2
membrane fabrication

141

5.2.2 Characterization of materials

142

5.2.3 Gas permeation measurement

143

5.3

144

Results and discussion

5.3.1 Carbonization of SiO2-ZrO2-acac into C-SiO2-ZrO2 and identificationof the carbon state144

5.3.2 Unique permeation properties of CO2 in C-SiO2-ZrO2 membranes 148
5.3.2.1

Understanding the interaction between CO2 and C-SiO2-ZrO2
148

5.3.2.2

Pressure-induced transition of CO2 flow in C-SiO2-ZrO2

membranes

151

5.3.3 Binary H2/CO2 mixture separation performance of a C-SiO2-ZrO2derived membrane
5.3.3.1

156

Effects of temperature, pressure, and CO2 feed mole fraction on

the binary mixture separation performance156

5.3.4 Outlook on the practical application of a C-SiO2-ZrO2 membrane for
H2/CO2 separation

159

5.4

Conclusions

161

Cited literature

162

Chapter 6
Conclusions and considerations for further studies
6.1

Summary and conclusions

168

6.2

Considerations for further studies

170

List of publications

173

Acknowledgements

175

iv

Hiroshima University Doctoral Dissertation

Chapter 1
General introduction
1.1 A brief background on membranes
The separation, concentration and purification of material mixtures is a major
challenge facing industries such as chemical, petrochemical, petroleum, textile, food,and pharmaceutical industries [1]. For this reason, several separation techniques havebeen developed over the years for different applications. Such processes include
distillation, adsorption, absorption, extraction, crystallization, cryogenic separation
etc. In recent decades however, a growing attention and shift towards a new
separation technique, membrane processes and membrane separation, has occurred[2]. This is because membranes provide a simple, cost effective and energy efficient
alternative to the state-of the-art techniques [3, 4].
A membrane in the general sense refers to a ‘thin’, selectively permeable, or semipermeable barrier or boundary that allows some material species such as molecules,
ions, or other similarly sized particles to pass through it and prevents the passage ofothers. Membranes are abundantly found in nature as part of our daily lives and arereferred to as biological membranes [2]. The basic function of a biological membranein biological systems is to define boundaries between and within cellular components.
This ability to separate between components and allowing communication between
them by a tightly regulated transport of materials make many cellular processes
possible [5]. These biological membranes served as model systems that inspired theresearch into synthetic membranes [5]. Perhaps the first recorded study of a synthetic
membrane inspired by pig’s bladder was in an ethanol-water mixture by A. Nollet in
1752 [6]. Thereafter, T. Graham in 1866 studied the diffusion of gases in a synthetic
rubber membrane [7].
1.1.1 Synthetic membranes
A synthetic membrane, therefore, is a permeable or semi-permeable phase such asa polymer, inorganic material or metal fabricated as a selective or contacting barrier tocontrol the relative transport of matter (such as gases, vapors, liquids, or a mixture ofthese) and energy through itself [2, 8]. Separation membranes serve as a third phase1

Hiroshima University Doctoral Dissertation

creating a barrier between two bulk phases. The aim of membrane processes is toconcentrate a component into a phase by mass transport into another phase throughthe permeable barrier as illustrated by Fig. 1-1 [9]. Mass transport is achieved byseveral driving forces of which pressure driving force is the most easily andeconomically applied to achieve global impact [10]. This broad definition of amembrane and the membrane process gives rise to a very wide application of
membranes as shown in Table 1-1.

Fig. 1-1 Mass transport through a synthetic membrane [9]From all the membrane separation processes listed in Table 1-1, perhaps one of
the most important application of membranes is in the separation of gas mixtures [9,11]. Table 1-2 lists the different membrane processes against their respective yearly
growth in industrial sales. From these processes, gas separation shows the highest
yearly growth of 15% [9].

2Hiroshima University Doctoral Dissertation

Table 1-1 Types of synthetic membrane processes and applications [8, 10]
Membrane process

Driving force

Application(s)Microfiltration (MF)Transmembrane pressure

Clarification, sterile

difference

filtration

Transmembrane pressure

Separation ofdifference

macromolecular solutions

Transmembrane conc.

Separation ofdifference

macromolecular solutions

Ultrafiltration (UF)
Dialysis (D)and micro-solutes
Nanofiltration (NF)

Transmembrane pressure

Separation of small

difference

organic compounds andsalt solutions

Reverse Osmosis

Transmembrane pressure

Separation of micro-(RO)difference

solutes and salt solutions

Pervaporation (PV)Transmembrane fugacity

Separation of volatile

difference

liquid mixtures

Gas or vapor

Transmembrane pressure

Separation of

separation (GS/VP)

difference

gaseous/vapor mixtures

Electrodialysis (ED)Transmembrane voltage

Dewatering of suspended

difference

solids in solution

Table 1-2 Membrane processes and yearly growth in sales
Membrane process

Growth[%/yr]

Dialysis10

Microfiltration

8

Ultrafiltration

10

Reverse osmosis

10

Gas separation

15

Electrodialysis

5Electrolysis

5

Miscellaneous

103Hiroshima University Doctoral Dissertation

1.1.2 Gas separation membranes
As pointed out previously, gas separation processes are one of the most important
applications of synthetic membranes. Sholl and Lively very importantly noted an‘inexhaustive’ list of seven separation that could potentially change the world [12]. 3
among these separation processes include gas separation- light hydrocarbons fromheavy ones, alkenes from alkanes and CO2 capture from greenhouse gases. Apartfrom these, other very important gas separations in no order of importance include
natural gas processing (acidic gas separation, C3+ hydrocarbons removal, water
separation, nitrogen separation, helium separation) [13-15], separation of hydrogenfrom light hydrocarbons, nitrogen/argon separation in ammonia plants [14-16].
Emerging important applications also include CO2/N2 separation, hydrogen/CO2
separation, oxygen separation from air etc. [14, 15].
Gas separation membranes are usually tailored to specific applications as no one
membrane type can be applied to all the gas separation scenarios. Therefore,
processes are thoroughly understood before any useful membrane is to be designed[14]. Furthermore, to choose the best membrane, the interactive mechanism of thegases involved with the membrane material and its structure are important. Several
gas/membrane interactive mechanisms have been proposed.
1.1.3 Gas permeation mechanisms through membranes
There are several mechanisms that have been proposed for the permeation ofgases through thin-layer membranes. Some membrane materials exhibit mostly one
type of gas permeation mechanism while others exhibit different kinds ofmechanisms. Also, a single gaseous specie can be found to permeate a membrane bydifferent mechanisms at the same time or at different temperature ranges as iscommon for CO2 permeation through silica-based membranes [17-19]. The mostencountered gas permeation mechanisms are shown in Fig. 1-2 [20].

4Hiroshima University Doctoral Dissertation
(a)

(b)

(c)

(d)

Bulk flow

Knudsen diffusion

Surface diffusion

Gas translational
diffusion(e)

Solid-state
diffusion

Fig. 1-2 Different mechanisms of gas permeation through membranes [20]Generally, gas transport properties through porous media can be determined by
the ratio of the number of molecule-molecule collisions to molecule-pore wall
collisions [21]. This gave rise to the Knudsen number expressed as the ratio of themean free path to the mean pore width as shown in Equation 1-1.
𝜆

(1-1)

𝐾𝑛 = 𝑑𝑝

Where 𝐾𝑛 is the Knudsen number, 𝜆 is the mean free path and 𝑑𝑝 is the mean pore
diameter. The mean free path is defined as the means distance travelled by a gas
molecule between successive collisions. When 𝐾𝑛 << 1, this indicates that the pore
width is much larger than the mean free path and therefore, molecules tend to collide
more frequently with each other than with the walls of the pore. In this situation,
viscous flow results. On the other hand, when 𝐾𝑛 >> 1, the mean free path is muchgrater than the mean pore with and the gas molecules collide with the walls more
frequently than they collide with themselves. Specific mechanisms are discussed in
the following.
1.1.3.1 Bulk-fluid flow
Bulk-fluid or viscous flow through a membrane pore occurs by the HagenPoiseuille mechanism whereby the pore diameter of the membrane is very large
compared to the mean free path of travel of the permeating gas molecules. In thissituation, the Knudsen number, 𝐾𝑛 << 1 and the gas molecules collide with eachother more than they do with the pore walls [22]. Fig. 1-2a illustrates such a scenario.
1.1.3.2 Knudsen diffusion
5

Hiroshima University Doctoral Dissertation

A Knudsen permeation mechanism dominates when the Knudsen number, 𝐾𝑛 >>1 and the gas molecules collide with the pore wall more than with themselves as
illustrated in Fig. 1-2b. Knudsen mechanism is strictly dependent on the molecularweight of the permeating gas species and lighter weighing gas permeate faster thangases with higher molecular weights [21]. Here, the gas diffusivity is determined by
the combined consideration of the gas kinetic velocity and the geometric parametersof the associated pores [20].
1.1.3.3 Surface diffusion
In the preceding Knudsen mechanisms discussed, it was assumed that thecollisions occurring between the gaseous molecules and the pore walls are strictly andcompletely elastic without any interactions. However, this situation is not always thecase. At low temperatures when the kinetic energy of the permeating gas molecules is
not enough to escape the potential energy barrier of the pore surface, the gas
molecules tend to adsorb onto the pore wall surface at the pore mouth and diffuse
through the membrane via the surface of the pore wall (Fig. 1-2c). Gases like CO2,
CH4, N2 and CO show significant surface diffusion contribution to the overall
transport mechanism at temperatures below 100 oC [23, 24].
1.1.3.4 Gas translational diffusion
When the mean pore diameter of a membrane is very small in the region of
micropores, mechanisms such as Knudsen diffusion, surface diffusion and molecularsieving combine to dictate the permeation of gas molecules [20, 25]. In microporous
membranes, the diffusion of gaseous molecules is heavily affected by the moleculelattice and molecule-molecule interactions. A unified mechanism was proposed by
Xiao and Wei referred to as gas-translational diffusion mechanism [26, 27].
1.1.3.5 Solid-state diffusion
A solid-state diffusion also known as solution diffusion occurs with a further
reduction in the mean pore size of the permeation medium. The pores become sosmall that the interaction of the permeating molecules with the membrane material
becomes very significant to the transport of molecules. This phenomenon occurs indense non-porous membranes. In solution-diffusion the permeating molecules
6Hiroshima University Doctoral Dissertation

dissolve into the membrane material and transport through the material via aconcentration gradient. Different molecules of similar sizes permeate at different rates
because of their different relative solubilities [28, 29]. In this instance, thepermeability is a combination of the solubility and diffusivity effects expressed in
Equation 1-2.
(1-2)

𝑃=𝑆𝑥𝐷

Where 𝑃 is the observed permeability of the permeating specie, 𝑆 is the solubility and𝐷 is the diffusivity. A solution-diffusion mechanism occurs in three steps [28]: (1) the
adsorption of the permeating molecule onto the surface of the membrane in the
upstream side (2) diffusion of the molecules through the membrane by a chemical
potential gradient and (3) the desorption of the molecules from the membrane surface
at the downstream end. This phenomenon is common in polymeric membranes, glassy
membranes, and metallic membranes [20].
1.1.4 Classification of gas separation membranes
Today there exists hundreds of different membrane types used in different
applications which makes it difficult to understand the trend in membranedevelopment. To evaluate the state of the art in membrane development and drive
further improvements in membrane science, technology and processes it is important
to classify membranes into various categories. Many efforts have been directed
towards this objective e.g., membranes for hydrogen separation by Ockwig and
Nenoff [30]. They identified and categorized membranes for hydrogen separation by
membrane material. However, this may still be limited in helping to get an overview
of the whole membrane science landscape. In this section, it is proposed that
membranes be categorized by material, configuration, and membrane composition.
An overview of this categorization is displayed in Fig. 1-3.

7

Hiroshima University Doctoral Dissertation
Classification of membranes

By material make-up

By physical configuration

By material composition➢ Polymeric membranes➢ Flat-sheet membranes➢ Single matrix membranes➢ Inorganic membranes➢ Hollow-fiber membranes➢ Mixed-matrix membranes✓ Microporous➢ Symmetric membranes➢ Doped matrix membranes✓ Dense➢ Asymmetric membranes➢ Composite matrix membranes

Fig. 1-3 An overview of the categorization of membranes by different parametersproposed here.
1.1.4.1 Classification by membrane material
The material make-up of a membrane refers to the chemical nature of its
constituting matter and can provide a reasonable basis for categorization. A solid
matter can be either organic or inorganic. Therefore, membranes can be generallyclassified by their material make-up into two large categories: polymeric andinorganic membranes. Several works have reported membranes from these twocategories [6, 7, 31-36].
(a) Polymeric membranes
As pointed out earlier, a solution-diffusion mechanism of mass transport ofpermeating molecules occurs in polymer membranes. Polymers are dense
macromolecular organic compounds composed of long sequences of one or more
species of atoms or group of atoms linked to each other by usually covalent bonds[37]. Polymers can be formed into membranes as flat sheets or hollow fiber
configurations [33] and are categorized majorly as glassy and rubbery polymers.
The performance and suitability of a polymer membrane to a specific application
may depend on its structure. Rubbery polymers operate above the glass transition
temperature. Above this temperature, the polymer chains have sufficient thermalenergy to allow segmental motions [38]. These include silicone rubber, polyisoprene8Hiroshima University Doctoral Dissertation[14], amide copolymers [11], neoprene, polyether, polyvinylchloride, polyurethane,
polybutadiene [38, 39] etc. In glassy polymers, operating below the glass transition
temperature means that the chains are rigid and possess lower free volume because
the polymer chains are more restricted in their motions [38, 39]. Glassy polymers
include cellulose acetate, polycarbonates, polyimides, polysulfone, polyethersulfone,
polyethylene oxide [11, 14, 38, 40] etc.
Challenges: the challenges associated with polymer membranes include
1. Trade-off relationship. In an ideal membrane, a combination of highselectivity and permeability is desired. However, this is difficult to achievewith polymer membranes [14, 32, 41, 42]
2. Physical ageing and plasticization. Plasticization occurs when the polymer
membrane is exposed to highly soluble gas species that result in considerable
swelling of the polymer matrix overtime. This increases free volume whichleads to higher flux and reduced selectivity. [43, 44](b) Inorganic membranes
Inorganic membranes can comprise of oxides (ceramic), metals or elemental
carbon [45, 46]. Going by their nomenclature inorganic membranes comprise of
inorganic matter which are arranged in either amorphous or crystalline from. They
can be broadly divided into two categories: (i) microporous inorganic membrane and(ii) dense inorganic membranes.
(i)Microporous inorganic membranes
Microporous inorganic membranes comprise of amorphous and crystalline

ceramic membranes with pore sizes less than 2 nm according to the IUPAC definition[47]. These membranes are capable of separating gas molecules by the molecularsieving mechanism because of the inherent range of pore sizes. Microporous
inorganic membranes are prepared as thin layers on porous inorganic supports that
present little permeation resistance but considerable mechanical strength [47].
Examples include amorphous silica [34, 48-51], zeolites (which contain silicon,
aluminum, and oxygen atoms arranged into a lattice structure) [52-55] and alumina[56, 57].
9

Hiroshima University Doctoral Dissertation(ii)

Dense inorganic membranes
Dense inorganic membranes on the other hand contain pores whose size is lessthat 0.5 nm [45]. Dense inorganic membranes comprise majorly of metals that allow
preferential permeation of certain molecules through their crystal lattice. Dense metal
membranes can achieve exceptionally high H2 selectivities and they typically include
special metals and their alloys such as palladium [35], niobium, tantalum, vanadium,nickel, platinum [58], Pd-Au, Pd-Ag, Pd-Cu [35], Pd-Ta [59], Pd-Ru, Pd-Nb, Pd-Mo,
Pd-Cu-Ag [60] etc.
Challenges: some challenges of inorganic membranes include [45]1. Hydrothermal instability in silica membranes.
2. High cost of inorganic processing in zeolites and reproducibility issues.
3. High cost of fabricating thin metal films (>$2000/m2). Instability of metal
membranes against phase transition and contaminants such as H2S.
1.1.4.2 Classification by physical configuration
Membranes can also be classified according to the manner in which the membrane
module has been physically configured for application. Polymeric and inorganic
membranes can be configured in various ways that are not particular to a membranematerial, although a configuration may be more suited to a membrane material thananother. The common configurations are discussed below and illustrated in Fig. 1-4 ad.
(a) Flat-sheet membranes
Flat-sheet membranes are formed as plate-and-frame modules (Fig. 1-4a) [33] andvarious examples exist in applications such as ultrafiltration [61], forward osmosis[62], membrane distillation [63] etc.
(b) Hollow-fiber membranes
Hollow-fiber membranes are tubular in nature (Fig. 1-4b) with very high packing
density [33]. Many polymeric membranes exist as hollow-fiber membranes [64, 65].
Many carbon molecular sieve membranes formed from hollow-fiber polymer modules
10Hiroshima University Doctoral Dissertation

also exist in this configuration [66]. Li et al also reported a successful fabrication of
an inorganic hollow-fiber membrane based on alumina [67].

(b) Hollow-fiber membrane(a) Flat-sheet membrane(c) Symmetric membrane(d) Asymmetric membrane

Fig. 1-4 The physical configurations of synthetic membranes(c) Symmetric membranes
A symmetric membrane is also referred to as an isotropic membrane and isdefined as one which is comprised of a uniform structure throughout the thickness ofthe membrane [33], as shown in Fig 1-4c. Membranes possessing symmetricconfiguration are usually free-standing unsupported membranes [68, 69].
(d) Asymmetric membranes
Asymmetric membranes also referred to as anisotropic membranes are configured
as multilayer modules. The structure of these membranes is such that the layers are
graded in porosity and particle size (Fig. 1-4c). The active separation layer is usuallysituated as the topmost densest layer which determines the gas separation ability of
the whole membrane [33] while the underlying layers serve as supports. Many
11

Hiroshima University Doctoral Dissertation

supported membranes are asymmetric in nature [70, 71]. For example, Fig. 1-5a and b
show a ZIF-8 [71] membrane supported on porous titania and a SiO2-ZrO2 supportedon porous alumina, respectively.

(a)

(b)
SiO2-ZrO2 layer

Porous alumina support

3 µm

Fig. 1-5 SEM images of supported membranes showing asymmetric structure. (a) a
ZIF-8 membrane supported on porous titania [71] (b) a SiO2-ZrO2 supported on
porous alumina.
1.1.4.3 Classification by material composition
In the preceding sections, it has been established that membranes are made up
differently in material and these membranes can be physically configured in different
ways depending on the applications. A new class of materials (and thus newclassification of membranes) is however necessary, borne out of the need to evolve
membrane separation performances to a new generation [38, 72, 73]. This new
generation of membranes are developed to give better performance metrics(selectivity, permeability, cost, handling etc.) over the previous generation
membranes. Such evolutionary methodologies involve the combination of different
matrices towards harnessing the strengths of each matrix for improved performance.
Some of the possible classifications that can result are discussed further and illustrated
in Fig. 1-6.

12

Hiroshima University Doctoral Dissertation

(b)

(a)

Inorganic filler
Pure matrix

Polymer matrix

(c)

(d)
i±

i±

DopantParent matrix

i±
i±

i±

i±

i±
i±

i±

i±
i±

i±

i±
i±

i±

Fig. 1-6 Illustration of the categorization of membranes by material composition. (a)single matrix membranes (b) mixed matrix membranes (c) doped matrix membranes(d) composite matrix membranes.
(a) Single matrix membranes
This is the most basic form of a membrane matrix- a membrane comprising of
only one type of material i.e., a single-component membrane. Many of the popularmembrane types fall under this category as they serve as the bases for further
development. Typically, single matrix membranes will include all various types ofpurepolymer

membranes(polydimethoxysilane,

polysulfone,polyimide,

polyetherimide, cellulose acetate etc.), pure ceramic membranes and organosilica
membranes (silica, zirconia, titania, alumina, bis(triethoxysilyl)ethane (BTESE) [74,75], bis(triethoxysilyl)methane BTESM [76], bis(triethoxysilyl)propane BTESP [77],bis(triethoxysilyl)acetylene BTESA [78] etc.), pure metals (palladium, nickel,
tantalum, molybdenum, niobium etc.), zeolites (ZIF, SAPO, AlPO4, MeAPO
structures etc.).
(b) Mixed-matrix membranes
Mixed-matrix membranes (MMM) are a combination of polymer and inorganic
matrices. A pioneering work on MMMs was published by Koros and co-workers [79].
13Hiroshima University Doctoral Dissertation

The concept of mixed-matrix membranes was introduced to overcome the limitationsof both polymer membranes (vis trade-off, physical ageing, and plasticization) andinorganic membranes (vis brittleness and cost of inorganic processing) as enumerated
in section 1.1.4.1 (a) and (b) [38, 39, 80, 81]. Typically, an MMM consists of a bulk
polymer phase into which the inorganic phase is dispersed [80, 81]. The inorganic
phase dispersed into the polymer bulk phase as shown in Fig. 1-6 (b) are referred to as
fillers. The inorganic fillers traditionally incorporated can be categorized as zeolitic
and non-zeolitic fillers [82]. The zeolitic fillers include conventional zeolites, AlPO
and SAPO molecular sieves, while non-zeolitic fillers include carbon molecularsieves (CMS), porous and dense silica nanoparticles, metal oxide nanoparticles [82],activated carbon, metal-organic frameworks (MOFs), carbon nanotubes (CNTs),
mesoporous materials, lamellar inorganic materials [39, 81].
(c) Doped matrix membranes
A membrane can acquire new characteristics and properties that improve its
performance by introducing into it a metal or non-metal ion. Such an addition isreferred to as doping. This strategy has been especially used in silica and silica-based
membranes [83-86]. Doping a material with an anion or cation is used to direct amembranes microstructure or functionality towards enhanced performance. For
example, fluorine as an anion is used to add a functionality to an organosilica
membrane for improved propylene/iso-butane separation [85] or for improvedhydrothermal stability of a silica membrane [87, 88], or for microstructural control ofthe membrane pore size [86]. On the other hand, Pd doping of a silica membrane has
been used to improve the hydrogen-permselectivity [89]. Fig. 1-6 (c) shows the effectof a dopant interacting with the permeating molecule.
(d) Composite matrix membranes
In mixed-matrix membranes and doped matrix membranes, fillers or dopants arepresent as separate entities in the bulk of the primary matrix [82, 90]. In mixed-matrix
membranes, the challenge of inorganic particles agglomeration in the polymer matrix
especially in glassy polymers exists as a major concern [81]. In most doped matrix
membranes (except for Pd-doped membranes for H2-permselectivity), the dopants are
not present as separate matrices but as functionality or microstructural change agents.
14

Hiroshima University Doctoral Dissertation

For composite matrix membranes, two separate structures are integrated into a single
matrix and act as one matrix for active gas separation. Several examples of composite
matrix membranes have been reported in literature. Such include ceramic composites(SiO2-ZrO2 [91-93], TiO2-ZrO2 [94-96], SiO2-TiO2 [97]). Furthermore, the different
structures composing a composite matrix can act as structural, functional, or
molecular sieving agents or a combination of these. These composites can beclassified into two major categories: inorganic-organic composites and inorganicinorganic composites.
(i)

Inorganic-organic composites
Inorganic-organic composites are diphasic hybrids that can be grouped into twolarge categories: class I and class II inorganic-organic hybrids. A class I hybrid issuch that the organic and inorganic matrices are linked by weak interactive forces
such as hydrogen and van der Waals bonding [98, 99]. Examples of class I hybrids isinterpenetrating polymer and inorganic networks (IPNs) [100]. In class II hybrids, on
the other hand, the interaction between the inorganic and organic parts, which
includes strong bonds like covalent and iono-covalent bonds, serve to improve theinorganic matrix qualities such as improved hydrothermal stability [101], pore size
controllability, [102, 103] targeted functionalization, [104] etc. Examples include
organically modified transition metal oxides, metal-organic frameworks, etc. [102,
104]
(ii)

Inorganic-inorganic composites
Inorganic-inorganic composites include several ceramic composites such as SiO2-

ZrO2, TiO2-ZrO2, SiO2-TiO2. Composite SiO2-ZrO2 membranes have been used forpervaporation and nanofiltration applications [105-107]. Hydrogen separation
performance of SiO2-ZrO2 hybrid membranes fabricated through chemical vapor
deposition (CVD) has also been investigated [93].

1.2 Silica and composite matrix ceramic membranes
Amorphous silica membranes and their derivatives have drawn a lot of interest in
the past few decades of gas separation membrane research due to their potential
application to a wide variety of gas separation systems in general [34, 51, 108, 109],
15

Hiroshima University Doctoral Dissertationand in particular to hydrogen separation due to their ease of fabrication, very highpermselectivity and scalability [30]. Factors required for a permselective membrane
are high flux, high selectivity, pore size controllability, and good physical andchemical stability. Silica membranes exhibit moderately high flux and excellent
selectivity [34, 110, 111], but membranes based on silica still suffer from someproblems [110]:(i)

Pore size control for other gas separation applications beyond hydrogen
separation(ii)

Instability under hydrothermal conditions

Fig. 1-7 shows the typical performance of a silica membrane under ahydrothermal condition. Fig. 1-8a shows the single gas permeance at 200 oC as afunction of kinetic diameter for a tetraethoxysilane (TEOS)-derived silica membranemeasured for various gases before and after exposure to boiling water for 2 hours as
carried out by ten Hove and co-workers [112]. Before exposure to boiling water, themembrane showed an impressive molecular sieving ability with high H2permselectivity. After an exposure of 72 hours to boiling water, this molecular sieving
property all but disappeared with the performance now similar to that of the
supporting alumina membrane. In a separate experiment by Castricum and coworkers,
a 5% water by weight impure n-butanol was dehydrated by pure and hybrid silica at
150 oC by PV process and the separation performance observed for several days as
shown in Fig. 1-8b [101]. Here, it can be seen that the dehydration performance of the
inorganic silica membrane dropped off after just a few days because of the highhydrophilicity [113].
(a)

(b)

16

Hiroshima University Doctoral Dissertation

Fig. 1-7 Performance of silica membranes in hydrothermal conditions (a) single gas
permeance at 200 oC as a function of kinetic diameter for a TEOS-derived silica
membrane measured for various gases before and after exposure to boiling water for
72 hours [112] (b) long-term separation performance: water content of permeate of a
hybrid silica membrane towards dehydration of n-butanol (5 wt% water) bypervaporation at 150 oC [113].
To develop silica, silica-based and other ceramic membranes with high permselectivity, pore size tunability, and hydrothermal stability, many researchers havestudied various techniques such as:(i)Spacer methods (-Si-R-Si-) [108, 114],(ii)

Templating methods (-Si-R’) [102, 115],(iii)

Anion and cation doping [83-90], and(iv)

Hybridization of silica with other metal oxides such as ZrO2 [91-93], TiO2[97], and Al2O3 [116, 117].

The hydrothermal, acidic, and basic stability of SiO2-ZrO2 ceramic composites
have been particularly proven [91, 93, 107]. Fig. 1-8 shows the gas permeance for a
SiO2-ZrO2 ceramic composite membrane as a function of kinetic diameter measured at
300 oC before and after exposure to steam at 40 and 100 kPa partial pressure and 500o

C for 20 hours, respectively, carried out by Yoshida and co-workers [91]. The

membrane was made with a 9:1 mixture of SiO2 to ZrO2. When compared with a pure
SiO2 membrane exposed to boiling water for 72 hours as shown in Fig. 1-7a, it is seen
here that the membrane quality of a composite SiO2-ZrO2 membrane withstands ahydrothermal atmosphere to a better degree.

17

Hiroshima University Doctoral Dissertation

(a)
(c)
(d)

(b)

Fig. 1-8 Gas permeance for a SiO2-ZrO2 ceramic composite membrane as a functionof kinetic diameter measured at 300 oC before and after exposure to steam: (a) fresh
membrane (b) membrane after exposure to 500 oC steam at 40 kPa partial pressure for
10 hours (c) membrane after exposure to 500 oC steam at 100 kPa partial pressure for
10 hours (d) Knudsen diffusion (based on SF6 permeance) [91].
1.2.1 State of the art of ceramic composite materials and membranes
Hybridization of silica with proven hydrothermally stable metal oxides such as
Al2O3, TiO2 and ZrO2 has been studied extensively [30, 91, 105, 117]. Initial specificstudies into the effect of ZrO2 and TiO2 on SiO2 nanoparticles were carried out by
catalysis researchers [120-124]. It was extensively observed that the acidic propertiesof the composite oxides were considerably better than the individual component
oxides. In another consideration, the presence of a second oxide containing a metal
ion with a larger ionic radius tends to reduce the surface energy and hence thesintering rate in severe environmental conditions [125]. Therefore, there seems to be asynergistic effect between the contributing oxides as on one hand, an oxide may bepresent to improve the acidic conditions of the composite, on the other hand, thesecond oxide may be present to improve reduce the surface energy of the composite.
As mentioned before, the various combinations of these silica and metal oxides inmembranes that have seen very extensive study include SiO2-ZrO2, TiO2-ZrO2, SiO2TiO2 and SiO2-Al2O3 [118] to a lesser degree. A ternary combination such as Al2O3SiO2-TiO2 has also been investigated [119]. These ceramic composites have been18

Hiroshima University Doctoral Dissertation

studied in various applications. Table 1-3 shows the various applications of ceramiccomposite membranes.
Table 1-3 Various applications of ceramic composite membranes (with preference forgas separation applications).
Ceramic composite

Composition Application

Ideal

Reference(ratio)systemselectivity [-]SiO2-ZrO2 (sol-gel)

9:1

GSa-H2, N2

28 (H2/N2)

SiO2-ZrO2 (CVD)

5:5

GSa -He, H2, N2

17

[91]

(He/N2), [92]25 (H2/N2)
SiO2-ZrO2 (CVD)

9:1

GSa -H2, N2

1400 (H2/N2)[93]

ZrO2-Y2O3-SiO2

9:0.5:0.5

GSa -H2, N2

3 (H2/N2)[126]

SiO2-ZrO2 (sol-gel)

5:5

NF-

-

[107, 107]

SiO2-TiO2 (sol-gel)

5:5

NF- Hep/Tol.

-

[97]

TiO2-ZrO2 (sol-gel)

5:5

GS- H2, CO2, N2

12 (H2/CO2), [127](sol-gel)

5 (H2/N2)
TiO2-ZrO2 (sol-gel)

5:5

NF-

-

[94, 95]

TiO2-ZrO2 (sol-gel)

5:5

NF-

-

[128]

Al2O3-SiO2-TiO2

4:1:1

-

-

[119]

(sol-gel)a

Operated in a steam atmosphere, GS- gas separation, NF- nanofiltration, Hep.–

Heptane, Tol.- Toluene
A few observations can be made from Table 1-3. The first is that only thecomposite of SiO2 and ZrO2 have been applied to gas separation. This is very likely
due to the highly molecular sieving ability that SiO2 can imbue the composite with.
Consequently, it does seem that a SiO2-ZrO2 composite with a higher proportion of
the SiO2 component will achieve better selectivities as observed for references 91 and93. Yoshida and co-workers [91] as well as Choi and co-workers [92], havedemonstrated that although a composite of 5:5 of SiO2 and ZrO2 can achieve a very
impressive hydrothermal stability, the selectivities can be low and sometimes
19Hiroshima University Doctoral Dissertation

approaching Knudsen levels [91]. Furthermore, it also appears that other ceramiccomposites not SiO2-ZrO2 are limited to nanofiltration (NF) in their applications.
Therefore, the challenges of improving the molecular sieving properties of anequimolar (5:5) ceramic composite and broadening their applicability are identified.
Researchers have found a few opportunities in organic coordination ligands [129].
1.2.2 Organic coordination ligands: a broad field of opportunities
The versatility of sol-gel chemistry has been instrumental to the formation of new
functional materials. One major aspect of this versatility is the ability to obtainproducts with predetermined characteristics based on the experimental conditions, andanother aspect is that it allows for the synthesis of inorganic-organic hybrid materials
through the introduction of organic groups in solution [130]. One way to add anorganic group is to chemically modify a precursor prior to synthesis [120]. A good
example of chemically modifying precursors is found in the use of coordination
ligands to form coordinate complexes where the precursor to be modified is atransition metal ion compound. One of the pioneering works on the chemical
modification and sol-gel chemistry of transition metal alkoxides were done by
Sanchez, Livage and their co-workers [131-133]. They studied topics that included
the partial charge model, gelation, coordination ligands, alcohol interchange etc.
1.2.2.1 Coordination ligands
As most transition metal alkoxides are very reactive in hydrolysis andcondensation reactions, there arise the need to stabilize or slow them down in order toavoid precipitation or agglomeration especially when forming composite oxides [130,132]. This is achieved by chemically modifying the alkoxides with complexing agentsthat are usually hydroxylated nucleophilic ligands such as carboxylic acids and βdiketones. For example, acetylacetone ACAC (CH3-CO-CH2-CO-CH3) in its enolic
form contains hydroxyl groups which can react with metal alkoxides. According to
the partial charge theory, the acetylacetonate group ACAC- (CH3-CO-CH2-CO-CH2)
reacts with a titanium tetrapropoxide for example to give the oligomeric specie Ti(OPr)3 (CH3-CO-CH2-CO-CH2)-structurally depicted in Fig 1-9- and favors therelease of the alcohol PrOH (propanol).

20

Hiroshima University Doctoral Dissertation

OPr

H3C
HC

C

O

C

O

OPr

Ti

H3C

OPr

Fig. 1-9 Structural formula of an oligomer formed between a titanium tetrapropoxide
and acetylacetone.
Such nucleophilic reactions between hydroxylated ligands and transition metal
alkoxides have been studied extensively: effect of acetylacetone on the particle sizesof TiO2 colloids [134], x-ray structures of titanium alkoxide β-diketones derivatives[135], ethyl acetoacetate distribution in the course of titanium alkoxide chelation[136], the influence of chemical modification on zirconium alkoxides as precursors in
sol-gel process studied by EXAFS spectroscopy [137], hydrolytic stability of organic
ligands in different alkoxide complexes [138] etc.
Since Hoebbel and co-workers [138] found that some of the ligands chelated totransition metals are hydrolytically stable to a very good extent in the range of a week
period under continuous stirring in the presence of water, the opportunities for
adoption of chelating ligands into composite ceramic membranes has drawn lots ofattention beyond catalysis applications.
1.2.2.2 Organic coordination ligands as molecular engineering agents
Complexation of alkoxides with nucleophilic complexing ligands allows for thedesign of molecular precursors and the sol-gel synthesis of tailor-made materials[132]. Consequently, the role of a hydrolytically stable complexing ligand can be predesigned according to two major functions: as a network modifier or network former[139] (Fig. 1-10).

21

Hiroshima University Doctoral Dissertation

(a)

(b)

Fig. 1-10 Illustration of network modification and network formation organicallymodified ceramics [139](a) Network modifiers
When the complexing ligand is a simple no hydrolysable group, it will have a
network modifying effect [140] (Fig. 1-10a). In this case, the ligand can create aphysically or chemically different site [139]. Organic coordination ligands that fall inthis category include acetylacetone, acetic acid, isoeugenol, ethyl acetoacetate etc. andthey are depicted in Fig. 1-11a.
(b) Network formers
A complexing ligand working as a network former is when the ligand bears areactive group that can, for example, polymerize or copolymerize [139, 140] (Fig. 110b). Complexing ligands that can serve as network formers include allyl
acetoacetate, methacryloxyethyl acetoacetate, methacrylic acid etc. as shown in Fig.
1-11b.

22

Hiroshima University Doctoral Dissertation

Acetylacetone

Isoeugenol

Acetic acid

Ethyl acetoacetate(a)Allyl acetoacetate

Methacryloxyethyl acetoacetate

Methacrylic acid(b)
Fig. 1-11 Some organic coordination/ complexation ligands: (a) network modifiers (b)network formers.
1.2.3 Ligand-modified ceramic composite membranes
Embracing the unique chemistry between transition metal ions and organic
coordination ligands, a number of researchers have been able to develop a new classof membranes that have shown promise in gas separation and filtration applications.
The ligand-modified ceramic composites have shown better molecular sieving
performance compared to the corresponding unmodified composites with the same
mixed oxide composition. For example, Fukumoto et al., prepared an isoeugenolmodified TiO2-ZrO2 composite membrane that showed superior molecular sieving
selectivities compared to a an unmodified TiO2-ZrO2 composite membrane [141]. The
theory of the improved molecular sieving is that isoeugenol coordinated to thetitanium and zirconium transition metal centers modify the intra-particle network
pores by physically occupying the networks and serving as shield against the transportof large molecules. Fig. 1-12 illustrates this theory. Fukumoto et al. were able toprove this theory of intra-network occupation by changing the ligand/alkoxide molar
ratio and they found that by increasing the ligand content, the molecular sieving wasable to be improved. Elsewhere, Spijksma and co-workers also showed adiethanolamine-modified TiO2-ZrO2 composite membrane with H2 permeance andH2/butane selectivity of 3.0 x 10-7 mol m-2 s-1 Pa-1 and 54, respectively [142].

23

Hiroshima University Doctoral Dissertation

This development has thus led to new applications of ligand-modified composite
ceramic membranes to gas separation and filtration. Table 1-4 shows some of thesemembranes and their applications. Since one of the first research into the chemistry oftransition metals and chelating ligands occurred between the 1940s and 1960s [143,144] and the applications into composite ceramic membranes are relatively new (oneof the earliest appearing in 2001 by Benfer and co-workers [145]), there seems to be a
lot of gap in the scope of possibilities that can be achieved with organic
coordination/complexation ligands and composite ceramic membranes.

M

O

O

M

O

M

M

O

O

O
M

M

M
O

M

O
O

O
M

Intra-network occupying ligands

Fig. 1-12 Mechanism of molecular sieving in organic coordination ligand modified
ceramic network24

Hiroshima University Doctoral Dissertation

Table 1-4 Different ligand-modified ceramic composite membranes and theirapplications
Membrane

Modifying ligand

Appl.

Selectivity

Reference[-]
TiO2-ZrO2

IsoeugenolGS

65 (He/N2)[141]

46 (CO2/N2)
TiO2-ZrO2

Diethanolamine (DEA)GS

12 (He/N2)[141]

1.1 (CO2/N2)
TiO2-ZrO2

Diethanolamine (DEA)GS

54 (H2/butane)[142]

TiO2-ZrO2

Methyl gallateGS

62 (He/N2)[146]

111 (He/CH4)
TiO2-ZrO2Ethyl ferrulateGS10.6 (He/N2)[146]

13.6 (He/CH4)
SiO2-TiO2

Acetylacetone

GS2.43 (He/N2)[147]

2.16 (N2/CO2)
SiO2-TiO2

IsoeugenolGS

2.39 (He/N2)[147]

2.22 (N2/CO2)
TiO2-ZrO2Ethyl acetoacetate

NF

[129]

TiO2-ZrO2

2,3-dihydroxynaphthalene

NF

[129]

Appl.- Applications, GS- gas separation, NF-nanofiltration

1.3 The scope and outline of this thesis
The overall objective of this dissertation is to engineer the microstructure of SiO2ZrO2 ceramic composite towards the development of improved molecular separation
membranesbyapplying

diversemodification

techniques.

Organic

coordination/complexation ligands have shown that the networks of ceramiccomposites can be tailored towards fabricating molecular sieving membranes, but a
few concerns remain. First, most of the reported works on ligand-modified ceramiccomposite membranes have been limited to TiO2-ZrO2 when opportunities may lie
elsewhere. Also, the engineered modifications of ceramic composite membrane
networks have stopped at the ligand modification phase with no further investigation
into more modification opportunities by the ligand transformation. Fig. 1-13 shows
the possible modification directions considered during this research period. Therefore,25

Hiroshima University Doctoral Dissertation

this work seeks to deepen the understanding and access further opportunities inorganic coordination/complexation ligands modification of ceramic composite
membranes.

Utilizing
network
forming ligands

Pyrolysis ofligands

Ligand modified network

Zr O Zr

Si

Si

O

O Zr

O Si

Zr O SiO Si

Organically crosslinked network

Carbonized ceramic network

Fig. 1-13 Alternative modification opportunities offered by ligand-modified ceramic
networks
The scope of this work is comprised of the following contents:
1. Preparation of organic ligand-modified SiO2-ZrO2 films, gels and membranes andevaluation of gas permeation characteristics
2. Evaluation of the effect of organic ligand types for microstructural engineering of
SiO2-ZrO2 membranes
3. Pyrolysis transformation of organic coordination/complexation ligands chelated to
SiO2-ZrO2 membranes and evaluation of gas permeation characteristics
This dissertation is outlined in 6 chapters thus:26

Hiroshima University Doctoral Dissertation

Chapter 1 is the “General introduction”, which presents a brief overview of
separation membranes and a detailed look at silica and composite ceramic membranes
for gas separation.
Chapter 2 is titled “Development of a molecular sieving membrane from anacetylacetonate-modified silica-zirconia ceramic composite”. In this work, anacetylacetonate-modified equimolar SiO2-ZrO2 composite-derived membrane withmolecular sieving properties was successfully fabricated. The sol-gel method was
successfully employed to chemically modify zirconium tetrabutoxide prior to cohydrolysis and -condensation with tetraethoxysilane, and it was then used to fabricate
a gas separation layer. An acetylacetonate-modified SiO2-ZrO2-derived membrane
prepared at 300 oC showed H2 permeance of 9.9 x 10-7 mol m-2 s-1 Pa-1 with a H2/SF6
permeance ratio of 7,600, which was a significant improvement over pure SiO2-ZrO2derived membranes (H2 permeance: 1.4 x 10-6 mol m-2 s-1 Pa-1, H2/SF6 permeanceratio: 11). Heat-treatment of an acac--modified SiO2-ZrO2-derived membrane
prepared at temperatures that ranged from 250 oC (H2 permeance: 4.5 x 10-8 mol m-2 s1

Pa-1, H2/CH4: 100, CO2/CH4: 60, H2/SF6: > 18,000 at 50 oC) to 550 oC resulted in an

improved H2 permeance of 3.4 x 10-6 mol m-2 s-1 Pa-1 with reduced permeance ratios(H2/CH4: 3, H2/SF6: 9) at 50 oC. A membrane prepared by heat-treating a 250 oC-fired
membrane at 300 oC showed the best trade-off with H2 permeance-H2/SF6 permeanceratios above the trade-off line compared with membranes prepared at othertemperatures.
Chapter 3 is titled “Microstructure engineering of ligand-modified silicazirconia membranes derived from network-forming precursors”. In this study aneasy template was devised for the fabrication and microstructural control of ligandchelated, silica-zirconia-based composite membranes with a thickness of less than 400
nm. Hybrid materials with crosslinking between the organic moieties weresuccessfully fabricated via the sol-gel co-condensation of 3-(trimethoxysilyl)propyl
methacrylate and zirconium (IV) butoxide chelated via acetylacetone and allyl
acetoacetate ligands using an azobisisobutyronitrile radical initiator as a crosslinking
agent. Characterization of the films and gels via FTIR, XRD and thermogravimetry
indicated that both the polymerization and hybridization of the networks wasachieved. Gas permeation experiments showed that the organic crosslink-derived
27

Hiroshima University Doctoral Dissertation

membranes possess molecular sieving properties (H2 permeance: ≈10-7 mol m-2 s-1 Pa1; H2/N2 selectivity: ≈20-30) that are superior to those of non-crosslink-derived

membranes (H2 permeance: ≈10-7 mol m-2 s-1 Pa-1; H2/N2 selectivity: ≈7-11).
Quantitative analysis of the microstructures of the membranes based on thepermeation behaviors of H2 and N2 revealed that organic crosslinking led tosignificant control of the microstructural characteristics of the membranes.
Chapter 4 is titled “Enhanced H2-permselectivity in a carbon-silica-zirconia
membrane derived from an acetylacetonate-ceramic composite pyrolysis”. One
viable technique for the retardation of high-temperature densification of H2permselective inorganic silica-derived membranes is the combination with ZrO2.
However, while ZrO2 successfully inhibits SiO2 phase densification, H2 permeance issacrificed. This work reports the preparation of a carbonized SiO2-ZrO2 composite
membrane with enhanced H2 permselectivity via a sol-gel process. The sol synthesisof the precursor SiO2-ZrO2-acetylacetonate composite was optimized by studying theeffect of water/alkoxide molar ratio on the final carbon-SiO2-ZrO2 properties.
Characterization by thermogravimetry, XRD and N2 adsorption revealed that utilizing
a water/alkoxide molar ratio of 60 was optimal for the retention of carbon,
amorphousness and microporosity. Furthermore, by combining SiO2 and ZrO2 in aratio of 9/1 in a carbon-SiO2-ZrO2 membrane, a high H2-permselectivity could beachieved. This membrane showed an H2 permeance of 16 x 10-7 mol m-2 s-1 Pa-1 andH2/N2 and H2/CH4 ideal selectivities of 75 and 148, respectively. More importantly,
the carbon-SiO2-ZrO2 membrane’s H2-permselectivity exceeded that of a similarly
pore-distributed unmodified SiO2-ZrO2 membrane with an H2 permeance of 3 x 10-7
mol m-2 s-1 Pa-1. This enhanced H2-permselectivity was attributed to the increase inpore volume contributed by ultra-micropore-bearing carbon nanoparticles.
Chapter 5 is titled “CO2 flow-switching behaviour in a pyrolysis-derived
carbon-silica-zirconia membrane promises a versatile high-temperature H2/CO2
separation”. Many researchers regard silica, silica-based and zeolite membranes as
the agents that will accomplish H2 separation. These membranes are expected to be
productive in various mixture systems and under very high temperatures. This work
reports the successful fabrication of a composite carbon-SiO2-ZrO2 ceramic
membrane with a unique pressure-induced switching of CO2 flows that allows
28

Hiroshima University Doctoral Dissertation

versatile H2/CO2 separation at elevated temperatures. TG-MS, DTG-TGA, FT-IR,
CP-MAS-13C-NMR, and TEM provide corroborative evidence of the carbonization ofstarting material SiO2-ZrO2-acetylcetonate into C-SiO2-ZrO2. The resultant C-SiO2ZrO2 displayed significant hysteresis in CO2 adsorption isotherm at a temperature
well above the critical temperature of CO2 (31 oC), which indicates structuralconformation. Furthermore, single-gas permeation measurements showing upstreampressures of 200 and 500 kPa reveal different permeation values for CO2 at 300 oC. In
separating a H2/CO2 mixture at 50 and 300 oC under upstream pressures of 200 and
500 kPa, respectively, the flow of H2 permeance reduces as the concentration of CO2
increases in the feed side at 50 oC (1.14 x 10-8 down to 3.9 x 10-9 mol m-2 s-1 Pa-1 at
200 kPa). The pressure-induced surface flow of CO2 at 300 oC and 500 kPa, however,reduces the hindrance to H2 flow and results in H2/CO2 selectivity of ~20-30 for all
CO2 concentrations, which is on a par with molecular sieving membranes. This novel
C-SiO2-ZrO2 material shows promise for many interesting applications.
Chapter 6 is the “Conclusions and considerations for further studies”. In thischapter the major conclusions from chpaters 2 through 5 are summarised and closedwith itemisation of the aspects that may be studied in the future.

Cited literature
1.

R. W. Rousseau, Handbook of separation process technology, John Wiley andSons, New York, 1987.

2.

H. Strathmann, L. Giorno, E. Drioli, An introduction to membrane science andtechnology, Consiglio Nazionale Delle Ricerche, Rome, 2006.

3.

R. Bounaceur, N. Lape, D. Roizard, C. Vallieres, E. Favre, Membraneprocesses for post-combustion carbon dioxide capture: a parametric study,
Energy, 31 (2006) 2556–2570.

4.

C. Charcosset, A review of membrane processes and renewable energies fordesalination, Desalination, 245 (2009) 214–231

5.

Y.-H. M. Chan, S. G. Boxer, Model membrane systems and their applications,Curr. Opin. Chem. Biol., 2007, 11:581–587

29

Hiroshima University Doctoral Dissertation

6.

J. A. Nollet, Extract of the observations made in Italy, by the Abbé Nollet, F.
R. S. on the Grotta de Cani, Philos. T. R. Soc., 47 (1752) 48-61

7.

T. Graham, On the absorption and dialytic separation of gases by colloid
septa, Philos. T. R. Soc., 156 (1866) 399-439

8.

K. Scott, Handbook of industrial membranes, Elsevier Advanced Technology,
Oxford, United Kingdom, 2006.

9.

H. Strathmann, Membrane separation processes: current relevance and future
opportunities, AIChE J., 47 (2001) 1077-1087

10.

W. J. Koros, Evolving beyond the thermal age of separation processes:
membranes can lead the way, AIChE J., 50 (2004) 2326-2334

11.

P. Bernardo, E. Drioli, G. Golemme, Membrane gas separation: a review/stateof the art, Ind. Eng. Chem. Res., 48 (2009) 4638-4663

12.

D. S. Sholl, R. P. Lively, Seven chemical separations to change the world,
Nature, 532 (2016) 435-437

13.

C. A. Scholes, G. W. Stevens, S. E. Kentish, Membrane gas separation innatural gas processing, Fuel, 96 (2012) 15-28

14.

R. W. Baker, B. T. Low, Gas separation membrane materials: a perspective,
Macromolecules, 47 (2014) 6999−7013

15.

R. W. Baker, Future directions of membrane gas separation technology, Ind.
Eng. Chem. Res. 41 (2002) 1393-1411

16.

D. F. Sanders, Z. P. Smith, R. Guo, L. M. Robeson, J. E. McGrath, D. R. Paul,
B. D. Freeman, Energy-efficient polymeric gas separation membranes for asustainable future: A review, Polymer 54 (2013) 4729-4761

17.

M. Kanezashi, J. O’Brien-Abraham, and Y. S. Lin, Gas permeation through
DDR-type zeolite membranes at high temperatures, AIChE J. 54 (2008) 14781486

18.

L. Yu, M. Kanezashi, H. Nagasawa, J. Ohshita, A. Naka, T. Tsuru, Fabricationand microstructure tuning of a pyrimidine-bridged organoalkoxysilane
membrane for CO2 separation, Ind. Eng. Chem. Res., 56 (2017) 1316-1326
30

Hiroshima University Doctoral Dissertation

19.

D. Lee, S. T. Oyama, Gas permeation characteristics of a hydrogen selective
supported silica membrane. J. Membr. Sci., 210 (2002) 291–306

20.

S. T. Oyama, M. Yamada, T. Sugawara, A. Takagaki, R. Kikuchi, Review onmechanisms of gas permeation through inorganic membranes, J. Jpn. Petrol.
Inst., 54 (2011) 298-309

21.

Y. Wall, Ompe-Aime-Mudimu, G. Braun, G. Brunner, Gas transport throughceramic membranes under super-critical conditions, Desalination 250 (2010)
1056–1059

22.

D. W. Stops, The mean free path of gas molecules in the transition régime, J.
Phys. D: Appl. Phys., 3 (1970) 685

23.

P. F. Zito, A. Caravella, A. Brunetti, E. Drioli, G. Barbieri, Knudsen andsurface diffusion competing for gas permeation inside silicalite membranes, J.
Membr. Sci., 523 (2017) 456–469

24.

P. F. Zito, A. Caravella, A. Brunetti, E. Drioli, G. Barbieri, Discrimination
among gas translation, surface, and Knudsen diffusion in permeation throughzeolite membranes, J. Membr. Sci., 564 (2018) 166–173

25.

H. Nagasawa, T. Niimi, M. Kanezashi, T. Yoshioka and T. Tsuru, Modified
gas-translation model for prediction of gas permeation through microporous
organosilica membranes, AIChE J., 60 (2014) 4199-4210

26.

J. Xiao, J. Wei, Diffusion mechanisms of hydrocarbons in zeolites- I. Theory,
Chem. Eng. Sci., 47 (1992) 1123-1141

27.

J. Xiao, J. Wei, Diffusion mechanisms of hydrocarbons in zeolites- II.
Analysis of experimental observations, Chem. Eng. Sci., 47 (1992) 1143-1159

28.

P. Pandey, R. S. Chauhan, Membranes for gas separation, Prog. Polym. Sci.,
26 (2001) 853-893

29.

J. G. Wijmans, R. W. Baker, The solution-diffusion model: a review, J.
Membr. Sci. 107 (1995) 1-21

30.

N. W. Ockwig, T. M. Nenoff, Membranes for hydrogen separation, Chem.
Rev. 107 (2007) 4078-4110
31

Hiroshima University Doctoral Dissertation

31.

J. K. Mitchel, On the penetrativeness of fluids, J. Membr. Sci., 100 (1995) 1116

32.

L. M. Robeson, Correlation of separation factor versus permeability forpolymeric membranes, J. Membr. Sci., 62 (1991) 165-185

33.

I. Pinnau, B. D. Freeman, Formation and modification of polymeric
membranes,

ACS

Symposium

Series;AmericanChemicalSociety:

Washington, DC (1999) 1-22
34.

R. M. de Vos, H. Verweij, Improved performance of silica membranes for gas
separation, J. Membr. Sci., 143 (1998) 37-51

35.

Ø. Hatlevik, S. K. Gade, M. K. Keeling, P. M. Thoen, A. P. Davidson, J. D.
Way, Palladium and palladium alloy membranes for hydrogen separation andproduction: History, fabrication strategies, and current performance, Sep.
Purif. Technol. 73 (2010) 59–64

36.

Z. Tang, J. Dong, T. M. Nenoff, Internal Surface Modification of MFI-Type
Zeolite Membranes for High Selectivity and High Flux for Hydrogen,Langmuir 25 (2009) 4848–4852

37.

R. J. Young, P. A. Lovell, Introduction to polymers, CRC Press, Taylor andFrancis Group, Boca Raton, 2011.

38.

M. R. Abdul Hamid, H.-K. Jeong, Recent advances on mixed-matrix
membranes fro gas separation: opportunities and engineering challenges,
Korean J. Chem. Eng., 35 (2018) 1577-1600

39.

G. Dong, H. Li, V. Chen, Challenges and opportunities for mixed-matrix
membranes for gas separation, J. Mater. Chem. A, 1 (2013) 4610-4630

40.

X. Feng, P. Shao, R. Y. M. Huang, G. Jiang, and R. -X. Xu, A study ofsilicone rubber/polysulfone composite membranes: correlating H2/N2 andO2/N2 permselectivities, Sep. Purif. Technol. 27 (2002) 211-223

41.

L. M. Robeson, Polymer membranes for gas separation, Curr. Opin. Solid St.
M. 4 (1999) 549-552

32

Hiroshima University Doctoral Dissertation

42.

L. M. Robeson, The upper bound revisited, J. Membr. Sci., 320 (2008) 390400

43.

M. Galizia, W. S. Chi, Z. P. Smith, T. C. Merkel, R. W. Baker, B. D. Freeman,
50th anniversary perspective: polymers and mixed matrix membranes for gas
and

vapor

separation:areview

andprospective

opportunities,

Macromolecules, 50 (2017) 7809-7843
44.

D. F. Sanders, Z. P. Smith, R. Guo, L. M. Robeson, J. E. McGrath, D. R. Paul,
B. D. Freeman, Energy-efficient polymeric gas separation membranes for asustainable future: a review, Polymer 54 (2013) 4729-4761

45.

H. Verweij, Inorganic membranes, Curr. Opin. Chem. Eng. 1 (2012) 156–162

46.

G. Q. Lu, J. C. Diniz da Costa, M. Duke, S. Giessler, R. Socolow, R. H.
Williams, T. Kreutz, Inorganic membranes for hydrogen production andpurification: a critical review and perspective, J. Colloid Interf. Sci. 314(2007) 589–603

47.

Y. S. Lin, Microporous and dense inorganic membranes: current status andprospective, Sep. Purif. Technol. 25 (2001) 39–55

48.

G. Cao, Y. Lu, L. Delattre, C. J. Brinker, G. P. Ldpez, Amorphous silica
molecular sieving membranes by sol-gel processing, Adv. Mater. 8 (1996)
588-591

49.

Y. Iwamoto, K. Sato, T. Kato, T. Inada, Y. Kubo, A hydrogen-permselective
amorphous silica membrane derived from polysilazane, J. Eur. Ceram. Soc. 25(2005) 257–264

50.

T. Tsuru, T. Morita, H. Shintani, T. Yoshioka, M. Asaeda, Membrane reactore
performance of steam reforming of methane using hydrogen-permselective
catalytic SiO2 membranes, J. Membr. Sci., 316 (2007) 53-62

51.

G. Li, M. Kanezashi, T. Yoshioka, T. Tsuru, Ammonia decomposition in
catalytic membrane reactors: simulation and experimental studies, AIChE J.
59 (2013) 168-179

52.

A. Tavolaro, E. Drioli, Zeolite membranes, Adv. Mater., 11 (1999)975-996

33

Hiroshima University Doctoral Dissertation

53.

J. Caro, M. Noack, P. Kolsch, R. Schafer, Zeolite membranes - state of theirdevelopment and perspective, Micropor. Mesopor. Mat. 38 (2000) 3-24

54.

J. Caro, M. Noack, Zeolite membranes – recent developments and progress,
Micropor. Mesopor. Mat. 115 (2008) 215–233

55.

N. Rangnekar, N. Mittal, B. Elyassi, J. Caro and M. Tsapatsis, Zeolite
membranes – a review and comparison with MOFs, Chem. Soc. Rev., 44(2015) 7128—7154

56.

H. P. Hsieh, R. R. Bhave, H. L. Fleming, Microporous alumina membranes, J.
Membr. Sci., 39 (1988) 221-241

57.

A. F. M. Leenaars, K. Keizer, A. J. Burggraaf, The preparation andcharacterization of alumina membranes with ultra-fine pores. Part 1Microstructural investigations on non-supported membranes, J. Mater. Sci. 19(1984) 1077-1088

58.

N. A. Al-Mufachi, N. V. Rees, R. Steinberger-Wilkens, Hydrogen selective
membranes: a review of palladium-based dense metal membranes, Renew.
Sust. Energ. Rev. 47 (2015) 540–551

59.

N. M. Peachey, R. C. Snow, R. C. Dye, Composite Pd/Ta metal membranes
for hydrogen separation, J. Membr. Sci., 111 (1996) 123-133

60.

T. A. Peters, T. Kaleta, M. Stange, R. Bredesen, Development of thin binary
and ternary Pd-based alloy membranes for use in hydrogen production, J.
Membr. Sci., 383 (2011) 124– 134

61.

N. Pezeshk, D. Rana, R.M. Narbaitz, T. Matsuura, Novel modified PVDF
ultrafiltration flat-sheet membranes, J. Membr. Sci., 389 (2012) 280– 286

62.

J. Wei, C. Qiu, C. Y. Tang, R. Wang, A. G. Fane, Synthesis andcharacterization of flat-sheet thin film composite forward osmosis membranes,
J. Membr. Sci., 372 (2011) 292–302

63.

M. Tomaszewska, Preparation and properties of flat-sheet membranes frompoly(vinylidene fluoride) for membrane distillation, Desalination 104 (1996)
1-11

34

Hiroshima University Doctoral Dissertation

64.

I. U. Khan, M. H. D. Othman, A. F. Ismail, T. Matsuura, H. Hashim, N. Abdul
Hadi Md Nordin, M. A. Rahman, J. Jaafar, A. Jilania, Status and improvementof dual-layer hollow fiber membranes via coextrusion process for gas
separation: a review, J. Nat. Gas Sci. Eng., 52 (2018) 215–234

65.

X. Y. Chen, S. Kaliaguine, D. Rodrigue, A Comparison between Several
Commercial Polymer Hollow Fiber Membranes for Gas Separation, Journal of
Membrane and Separation Technology, 6 (2017) 1-15

66.

L. Lei, F. Pan, A. Lindbrathen, X. Zhang, M. Hillestad, Y. Nie, L. Bai, X. He,
M. D. Guiver, Carbon hollow fiber membranes for a molecular sieve with
precise-cutoff ultramicropores for superior hydrogen separation, Nat. Comm.,
12 (2021) 268

67.

X. Tan, S. Liu, K. Li, Preparation and characterization of inorganic hollow
fiber membranes, J. Membr. Sci., 188 (2001) 87–95

68.

S. Kumbar, T. Selvam, C. Gellermann, W. Storch, T. Ballweg, J. Breu, G.
Sextl, ORMOCERs (organic-inorganic hybrid copolymers)-zeolite beta (BEA)
nanocomposite membranes for gas separation applications, J. Membr. Sci.,
347 (2010) 132-140

69.

H. Lu, S. Zhu, Interfacial synthesis of free-standing metal–organic framework
membranes, Eur. J. Inorg. Chem. (2013) 1294–1300

70.

J. A. van ‘t Hof, A. J. Reuvers, R. M. Boom, H. H. M. Rolevink, C. A.
Smolders, Preparation of asymmetric gas separation membranes with highselectivity by a dual-bath coagulation method, J. Membr. Sci., 70 (1992) 17-30

71.

J. Caro, Hierarchy in inorganic membranes, Chem. Soc. Rev., 45 (2016) 3468

72.

W. J. Koros, Evolving beyond the thermal age of separation processes:
membranes can lead the way. AIChE J., 50 (2004) 2326-2334

73.

W. J. Koros, A more aggressive approach is needed for gas separation
membrane development. Membrane, 31 (2006) 157-160

74.

N.

Moriyama,

H.

Nagasawa,

M.

Kanezashi,

K.

Ito,

T.

Tsuru,

Bis(triethoxysilyl)ethane (BTESE)-derived silica membranes: pore formation

35

Hiroshima University Doctoral Dissertation

mechanism and gas permeation properties, J. Sol-Gel Sci. Technol., 86 (2018)
63-72
75.

H. F. Qureshi, R. Besselink, J. E. ten Elshof, A. Nijmeijer, L. Winnubst,
Doped microporous hybrid silica membranes for gas separation, J Sol-Gel Sci
Technol 75 (2015) 180–188

76.

M. Kanezashi, W. N. Shazwani, T. Yoshioka, T. Tsuru, Separation of
propylene/propane binary mixtures by bis(triethoxysilyl)methane (BTESM)derived silica membranes fabricated at different calcination temperatures, J.
Membr. Sci., 415–416 (2012) 478–485

77.

R. Inoue, M. Kanezashi, H. Nagasawa, K. Yamamoto, T. Gunji, T. Tsuru,
Pore size tuning of bis(triethoxysilyl)propane (BTESP)-derived membrane forgas separation: Effects of the acid molar ratio in the sol and of the calcination
temperature, Sep. Purif. Technol., 242 (2020) 116742

78.

M. Guo, M. Kanezashi, H. Nagasawa, L. Yu, K. Yamamoto, T. Gunji, J.
Ohshita, T. Tsuru, Tailoring the microstructure and permeation properties ofbridged organosilica membranes via control of the bond angles, J. Membr. Sci.,
584 (2019) 56–65

79.

C. M. Zimmerman, A. Singh, W. J. Koros, Tailoring mixed matrix composite
membranes for gas separations, J. Membr. Sci., 137 (1997) 145-154

80.

M. A. Aroon, A. F. Ismail, T. Matsuura, M. M. Montazer-Rahmati,
Performance studies of mixed matrix membranes for gas separation: a review,Sep. Purif. Technol., 75 (2010) 229-242

81.

P. S. Goh, A. F. Ismail, S. M. Sanip, B. C. Ng, M. Aziz, Recent advances ofinorganic fillers in mixed matrix membrane for gas separation, Sep. Purif.
Technol., 81 (2011) 243–264

82.

H. Vinh-Thang, S. Kaliaguine, Predictive models for mixed-matrix membrane
performance: a review, Chem. Rev. 113 (2013) 4980−5028

83.

M. Kanezashi, M. Asaeda, Hydrogen permeation characteristics and stabilityof Ni-doped silica membranes in steam at high temperature, J. Membr. Sci.,
271 (2006) 86–93

84.

D. Uhlmann, S. Liu, B. P. Ladewig, J. C. Diniz da Costa, Cobalt-doped silica
membranes for gas separation, J. Membr. Sci., 326 (2009) 316–321
36

Hiroshima University Doctoral Dissertation

85.

M. Takenaka, H. Nagasawa, T. Tsuru, M. Kanezashi, Hydrocarbon
permeation properties through microporous fluorine-doped organosilica
membranes with controlled pore sizes, J. Membr. Sci., 619 (2021) 118787

86.

M. Kanezashi, M. Murata, H. Nagasawa, T. Tsuru, Fluorine doping of
microporous organosilica membranes for pore size control and enhancedhydrophobic properties, ACS Omega 3 (2018) 8612−8620

87.

M. Kanezashi, T. Matsutani, H. Nagasawa, T. Tsuru, Fluorine-induced
microporous silica membranes: Dramatic improvement in hydrothermal
stability and pore size controllability for highly permeable propylene/propane
separation, J. Membr. Sci., 549 (2018) 111–119

88.

M. Kanezashi, N. Hataoka, R. Ikram, H. Nagasawa, T. Tsuru, Hydrothermal
stability of fluorine-induced microporous silica membranes: effect of steam
treatment conditions, AIChE J. (2021) e17292.

89.

J. Yanga, W. Fana, C.-M. Bell, Effect of calcination atmosphere on
microstructure and H2/CO2 separation of palladium-doped silica membranes,
Sep. Purif. Technol., 210 (2019) 659–669

90.

M. Kanezashi, M. Sano, T. Yoshioka, T. Tsuru, Extremely thin Pd–silica
mixed-matrix membranes with nano-dispersion for improved hydrogen
permeability, Chem. Commun., 46 (2010) 6171–6173

91.

K. Yoshida, Y. Hirano, H. Fujii, T. Tsuru, M. Asaeda, Hydrothermal stability
and performance of silica-zirconia membranes for hydrogen separation inhydrothermal conditions, J. Chem. Eng. Jpn., 34 (2001) 523-530

92.

H.-S. Choi, C.-H. Ryu, G.-B. Hwang, Obtention of ZrO2-SiO2 hydrogenpermselective membrane by chemical vapor deposition method, Chem. Eng.
J., 232 (2013) 302-309

93.

S.-J. Ahn, A. Takagaki, T. Sugawara, R. Kikuchi, S. T. Oyama, Permeation
properties of silica-zirconia composite membranes supported on porousalumina substrates, J. Membr. Sci., 526 (2017) 409-416

94.

S. Anisah, W. Puthai, M. Kanezashi, H. Nagasawa, T. Tsuru, Preparation,
characterization, and evaluation of TiO2-ZrO2 nanofiltration membranes fired37

Hiroshima University Doctoral Dissertation

at different temperatures, J. Membr. Sci., 564 (2018) 691-699.
95.

S. Anisah, M. Kanezashi, H. Nagasawa, T. Tsuru, Hydrothermal stability andpermeation of TiO2-ZrO2 (5/5) nanofiltration membranes at high temperatures,
Sep. Purif. Technol., 212 (2019) 1001-1012.

96.

H. Guo, S. Zhao, X. Wu, H. Qi, Fabrication and characterization of TiO2/ZrO2
ceramic membranes for nanofiltration, Micropor. Mesopor. Mater. 260 (2018)
125–131.

97.

C. Guizard, A. Ayral, A. Julbe, Potentiality of organic solvents filtration withceramic membranes. A comparison with polymer membranes, Desalination
147 (2002) 275-280

98.

P. Judeinstein, C. Sanchez, Hybrid organic-inorganic materials: a land of
multidisciplinarity, J. Mater. Chem., 6 (1996) 511-525

99.

C. Sanchez, J. de A. A. Soler-Illia, F. Ribot, T. Lalot, C. R. Mayer, V. Cabuil,
Designed

Hybridorganic-inorganic

nanocomposites

fromfunctional

nanobuilding blocks, Chem. Mater., 13 (2001) 3061-3083
100.

K. Nakahiro, L. Yu, H. Nagasawa, T. Tsuru, M. Kanezashi, Pore structure
controllability and CO2 permeation properties of silica-derived membranes
with a dual-network structure, Ind. Eng. Chem. Res., 60 (2021) 8527−8537

101.

H. L. Castricum, A. Sah, R. Kreiter, D. H. A. Blank, J. F. Vente, J. E. ten
Elshof, Hydrothermally stable molecular separation membranes fromorganically linked silica, J. Mater.Chem., 18 (2008) 2150-2158

102.

N. K. Raman, C. J. Brinker, Organic “template” approach to molecular sieving
silica membranes, J. Membr. Sci., 105 (1995) 273-279

103.

M. Kanezashi, M. Kawano, T. Yoshioka, T. Tsuru, Organic-inorganic hybrid
silica membranes with controlled silica network size for propylene/propane
separation, Ind. Eng. Chem. Res. 51 (2012) 944-953

104.

L. Yu, M. Kanezashi, H. Nagasawa, J. Oshita, A. Naka, T. Tsuru, Pyrimidinebridged organoalkoxysilane membrane for high-efficiency CO2 transport via
mild affinity, Sep. Purif. Technol., 178 (2017) 232-241

105.

M. Asaeda, Y. Sakou, J. Yang, K. Shimasaki, Stability and performance ofporous silica-zirconia composite membranes for pervaporation of aqueousorganic solutions, J. Membr. Sci. 209 (2002) 163-175.

38

Hiroshima University Doctoral Dissertation

106.

W. Puthai, M. Kanezashi, H. Nagasawa, K. Wakamura, H. Ohnishi, T. Tsuru,
Effect of firing temperature on water permeability of SiO2-ZrO2 membranes
for nanofiltration, J. Membr. Sci. 497 (2016) 348-356.

107.

W. Puthai, M. Kanezashi, H. Nagasawa, T. Tsuru, SiO2-ZrO2 nanofiltration
membranes of different Si/Zr molar ratios: stability in hot water andacid/alkaline solutions, J. Membr. Sci. 524 (2017) 700-711.

108.

A. P. Dral, J. E. ten Elshof, Organic groups influencing microporosity inorganosilicas, Micropor. Mesopor. Mater. 267 (2018) 267–273

109.

H. L. Castricum, G. G. Paradis, M. C. Mittelmeijer-Hazeleger, R. Kreiter, J. F.
Vente, J. E. ten Elshof, Tailoring the separation behavior of hybridorganosilica membranes by adjusting the structure of the organic bridging
group, Adv. Funct. Mater. 21 (2011) 2319–2329

110.

M. Kanezashi, Sol-Gel Derived Silica Membranes, Encyclopedia of
Membrane Science and Technology, John Wiley and Sons, 2013

111.

R. M. de Vos, H. Verweij, High-selectivity, high-flux silica membranes forgas separation, Science. 279 (1998) 1710-1711

112.

M. ten Hove, M. W. J. Luiten-Olieman, C. Huiskes, A. Nijmeijer, L.
Winnubst, Hydrothermal stability of silica, hybrid silica and Zr-doped hybrid
silica membranes, Sep. Purif. Technol., 189 (2017) 48–53

113.

Q. Wei, F. Wang, Z.-R. Nie, C.-L. Song, Y.-L. Wang, Q.-Y. Li, Highlyhydrothermally stable microporous silica membranes for hydrogen separation,
J. Phys. Chem. B, 112 (2008) 9354–9359

114.

M. Kanezashi, Y. Yoneda, H. Nagasawa, T. Tsuru, K. Yamamoto, J. Ohshita,
Gas permeation properties for organosilica membranes with different Si/C
ratios and evaluation of microporous structures, AIChE J. 63 (2017) 44914498.

115.

Y. E. Lee, B. S. Kang, S. H. Hyun, C. H. Lee, Organic‐templating approach tosynthesis of nanoporous silica composite membranes (ii): MTES‐templating
and CO2 separation, Sep. Sci. Technol., 39 (2004) 3541–3557

116.

B. N. Nair, K. Keizar, W. J. Elferink, M. J. Gilde, H. Verweij, A. J. Burggraaf,
Synthesis, characterization and gas permeation studies on microporous silica
39

Hiroshima University Doctoral Dissertation

and alumina-silica membranes for separation of propane and propylene, J.
Membr. Sci., 116 (1996) 161-169
117.

Y. Gu, P. Hacarlioglu, S. T. Oyama, Hydrothermally stable silica-alumina
composite membranes for hydrogen separation, J. Membr. Sci., 310 (2008)
28-37

118.

M. Amanipour, E. G. Babakhani, A. Safekordi, A. Zamaniyan, M. Heidari,Effect of CVD parameters on hydrogen permeation properties in ananocomposite SiO2–Al2O3 membrane, J. Membr. Sci., 423–424 (2012)530–535

119.

Z.-Q. Zeng, X.-Y. Xiao, Z.-L. Gui, L.-T. Li, Al2O3 –SiO2 –TiO2 composite
ceramic membranes from sol–gel processing, Mater. Lett., 35 (1998) 67–71

120.

D. A. Ward, E. I. Ko, Preparing catalytic materials by the sol-gel method, Ind.
Eng. Chem. Res. 34 (1995) 421-433

121.

J. Navarrete, T. Lopez, R. Gomez, Surface acidity of sulfated TiO2-SiO2 solgels, Langmuir, 12 (1996) 4385-4390

122.

F.

Garbassi,

L.

Balducci,

R.

Ungarelli,

Sol–gel

preparation

andcharacterization of spherical ZrO2–SiO2 particles, J. Non-Cryst. Solids, 223(1998) 190–199
123.

T. Lopez, F. Tzompantzi, J. Navarrete, R. Gomez, J. L. Boldu´, E. Mun˜oz, O.
Novaro, Free radical formation in ZrO2–SiO2 sol–gel derived catalysts, J.
Catal., 181 (1999) 285–293

124.

Z.-G. Wu, Y.-X. Zhao, D.-S. Liu, The synthesis and characterization ofmesoporous silica–zirconia aerogels, Micropor. Mesopor. Mater. 68 (2004)
127–132

125.

Y.-S. Lin, C.-H. Chang, R. Gopalan, Improvement of thermal stability ofporous nanostrucutured ceramic membranes, Ind. Eng. Chem. Res. 33 (1994)
860-870

126.

J. Fan, H. Ohya, T. Suga, H. Ohashi, K. Yamashita, S. Tsuchiya, M. Aihara,
T. Takeuchi, Y. Negishi, High flux zirconia composite membrane forhydrogen separation at elevated temperature, J. Membr. Sci., 170 (2000) 113–
125
40

Hiroshima University Doctoral Dissertation

127.

T. Van Gestel, D. Sebold, F. Hauler, W. A. Meulenberg, H.-P. Buchkremer,
Potentialities of microporous membranes for H2/CO2 separation in future
fossil fuel power plants: evaluation of SiO2, ZrO2, Y2O3–ZrO2 and TiO2–ZrO2
sol–gel membranes, J. Membr. Sci., 359 (2010) 64–79

128.

H. Guo, S. Zhao, X. Wu, H. Qi, Fabrication and characterization of TiO2/ZrO2
ceramic membranes for nanofiltration, Micropor. Mesopor. Mater., 260 (2018)125-131

129.

Y. Sada, T. Yoshioka, K. Nakagawa, T. Shintani, R. Iesako, E. Kamio, H.
Matsuyama, Preparation and characterization of organic chelate ligand (OCL)templated TiO2–ZrO2 nanofiltration membranes, J. Membr. Sci., 591 (2019)117304130.

C. A. Milea, C. Bogatu, A. Duta, The influence of parameters in silica sol-gel
process, in: Bulletin of the Transilvania University of Brasov, Series I:
Engineering sciences, 4 (2011) 59-66

131.

C. Sanchez, J. Livage, M. Henry, F. Babonneau, Chemical modification of
alkoxide precursors, J. Non-Cryst. Solids, 100 (1988) 65-76

132.

J. Livage, C. Sanchez, Sol-gel chemistry, J. Non-Cryst. Solids., 145 (1992) 1119

133.

J. Livage, M. Henry, C. Sanchez, Sol-gel chemistry of transition metal oxides,
Prog. Solid St. Chem., 18 (1988) 25-341

134.

H.-J. Chen, L. Wang, W.-Y. Chiu, Chelation and solvent effect on the
preparation of titania colloids, Mater. Chem. Phys., 101 (2007) 12–19

135.

R. J. Errington, J. Ridland, W. Clegg, R. A. Coxall, J. M. Sherwood, βdiketonate derivatives of titanium alkoxides: x-ray crystal structures andsolution dynamics of the binuclear complexes [{Ti(OR)3(dik)}2], Polvhedron,
17 (1998) 659-674

136.

S. Kurajica, I. Skoric, I. Lozic, V. Mandic, Ethyl acetoacetate ligand
distribution in the course of titanium n-butoxide chelation, Mater. Chem.
Phys., 147 (2014) 1058-1067

41

Hiroshima University Doctoral Dissertation

137.

D. Peter, T. S. Ertel, H. Bertagnolli, EXAFS study of zirconium alkoxides asprecursors in sol-gel process: II. The influence of the chemical modification, J.
Sol-Gel Sci. Technol., 5 (1995) 5-14

138.

D. Hoebbel, T. Reinert, H. Schmidt, E. Arpac, On the hydrolytic stability oforganic ligands in Al-, Ti- and Zr-alkoxide complexes, J. Sol-Gel Sci.
Technol., 10 (1997) 115-126

139.

H. Schmidt, Preparation, application and potential of ORMOCERs, Sol-gel:science and technology; proceedings of the winter school on glass andceramics from gels, São Carlos (SP), Brazil, 14 - 19 August 1989. Ed. by M.
A. Aegerter. Singapore: World Scientific, (1989) 432-469

140.

C. Sanchez, J. de A. A. Soler-Illia, F. Ribot, T. Lalot, C. R. Mayer, V. Cabuil,
Designed

Hybridorganic-inorganic

nanocomposites

fromfunctional

nanobuilding blocks. Chem. Mater., 13 (2001) 3061-3083
141.

T. Fukumoto, T. Yoshioka, H. Nagasawa, M. Kanezashi, T. Tsuru,
Development and gas permeation properties of microporous amorphous TiO2ZrO2-organic composite membranes using chelating ligands, J. Membr. Sci.,
461 (2014) 96-105.

142.

G. I. Spijksma, C. Huiskes, N. E. Benes, H. Kruidhof, D. H. A. Blank, V. G.
Kessler, H. J. M. Bouwmeester, Microporous zirconia–titania composite
membranes derived from diethanolamine-modified precursors, Adv. Mater.,18 (2006) 2165–2168

143.

G. Klopman, A semiempirical treatment of molecular structures. I.
electronegativity and atomic terms, J. Am. Chem. Soc., 86 (1964) 1463-1469

144.

W. H. Stockmayer, Theory of molecular size distribution and gel formation in
branched-chain polymers, J. Chem. Phys., 11 (1943) 45-55

145.

S. Benfer, U. Popp, H. Richter, C. Siewert, G. Tomandl, Development andcharacterization of ceramic nanofiltration membranes, Sep. Purif. Technol.,22-23 (2001) 231–237

146.

T. Tachibana, T. Yoshioka, K. Nakagawa, T. Shintani, E. Kamio, H.
Matsuyama, Gas permeation characteristics of TiO2-ZrO2-aromatic organic
chelating ligand (aOCL) composite membranes, Membranes, 10 (2020) 388

147.

J. Mendez-Vivar, R. Mendoza-Serna, SiO2-TiO2 membranes by the sol-gel
process, Silicon Chemistry 3 (2006) 59–64
42

Hiroshima University Doctoral Dissertation

Chapter 2
Development of a molecular sieving membrane from an acetylacetonatemodified silica-zirconia ceramic composite
2.1 Introduction
The quest for a highly efficient and cost-effective means for separating important
gases and liquids has gained prominence in the field of separation technology in thepast century. Such research efforts have successfully yielded a variety of techniquessuch as membrane separation technology, which has attracted the most interest [1].
Inorganic membranes are currently in great demand for use in H2 [2-5, 6], CO2/CH4[7, 8], and CO2/N2 [9, 10] separation systems, as well as in olefin/paraffin systems[11, 12] and in the purification of natural gas [11]. Inorganic membranes can bebroadly divided into porous and dense versions. Whereas dense membranes such as
those fabricated from palladium and its alloys offer separation ability via a solutiondiffusion mechanism, porous inorganic membranes such as silica, alumina, andzirconia separate gases by selective passage of molecules through a network of poresaccording to their size. Palladium membranes have shown high H2 permeance, butthey are difficult to fabricate and are plagued by high module costs [13, 22]. A mixedmatrix membrane consisting of palladium or palladium-cobalt particles dispersed in a
network of amorphous inorganic silica is also possible as already reported by
Kanezashi [14] and Ballinger et al. [15].
Much interest has been shown to inorganic amorphous membranes and theirderivatives concerning their potential application to a wide variety of gas separation
systems in general [2-11], and in particular to hydrogen separation due to their ease of
fabrication, low cost of production and scalability [1]. Factors required for a permselective membrane are high flux, high selectivity, pore size controllability, and good
physical and chemical stability. Silica membranes exhibit moderately high flux and
excellent selectivity [14, 16, 17], but size control and instability under hydrothermal
conditions are known problems [14]. To develop silica and silica-based membranes
with high perm-selectivity, pore size tunability, and hydrothermal stability, many
researchers have studied various techniques such as spacer methods [18, 19],43

Hiroshima University Doctoral Dissertation

templating methods [20], anion and cation doping [6, 16, 23], and hybridization ofsilica with other metal oxides such as Al2O3 [12, 24], TiO2 [25, 26], and ZrO2 [5, 21].
Hybridization of silica with proven hydrothermally stable metal oxides such as
TiO2 and ZrO2 has been studied extensively [1, 5, 21, 26, 27]. Hybrid SiO2-ZrO2
membranes have been used for pervaporation and nanofiltration applications [27, 28].
Hydrogen separation performance of SiO2-ZrO2 hybrid membranes fabricated throughchemical vapor deposition (CVD) has also been investigated [5]. Zirconium oxide
added to silica in an equi-molar ratio has shown the best hydrothermal stability tomoderately high temperatures, but perm-selectivity remains very poor due to the large
pore size created by adding a high zirconia content [21]. The large pores permit bothlarge and small molecules to permeate as in Knudsen flow [29, 30]. However, afurther modification to this network structure offers the possibility to change themolecular sieving characteristics of silica-zirconia (1:1) composite membranes.
In the fields of catalysis and separation technology, the versatility of sol-gel
chemistry has been instrumental to the formation of new functional materials. One
major aspect of this versatility is the ability to obtain products with predeterminedcharacteristics based on the experimental conditions, and another aspect is that itallows for the synthesis of inorganic-organic hybrid materials through the introductionof organic groups in solution [31]. One way to add an organic group is to chemicallymodify a precursor prior to synthesis [32]. A good example of chemically modifying
precursors is found in the use of chelating ligands to form coordinate complexes
where the precursor to be modified is a transition metal ion compound. Coordinate
polymer chemistry has been widely studied and found to be useful in synthesizing
metal-organic frameworks (MOFs) [33-35]. Coordinate covalent bonds are formed incoordinate complexes by the occupation of d-orbitals belonging to a transition metal
by free electron pairs donated by chelating ligands. For polydentate ligands, MetalOrganic Frameworks (MOFs) usually result [33-35]. In the case of mono- or bidentate
ligands, MOFs are rarely formed [36, 37] and the formation of the coordinate
complex chains will require other polymerization sites.
The formation of coordination complexes between zirconium ions and various
chelating ligands has been studied [36, 38] with the most hydrolytically stable being
the use of acetylacetone (Hacac), a bidentate ligand, to form an acetylacetonate44

Hiroshima University Doctoral Dissertation

zirconium coordination complex [38], which implies the retention of acetylacetonate
anions in the polymer sol even after hydrolysis and condensation have occurred forseveral hours. As such, a hydrolytically stable modified zirconium precursor can be
co-polymerized and co-condensed with silica to achieve an acac--modified equi-molar
SiO2-ZrO2 polymer sol. The acetylacetone ligand remaining in the modified SiO2ZrO2 (1:1) is expected to translate into a modified pore structure applicable to gas
separation. This possibility has been reported by both Fukumoto et al. [50] andSpijksma et al. [51], where the pore sizes of TiO2-ZrO2 were controlled using
isoeugenol (isoH) and diethanolamine (DEA), respectively.
To study the possibility of a chemically modified silica-zirconia (1:1) membrane
for gas separation applications, acetylacetone-modified silica-zirconia (1:1) was
prepared by the sol-gel method and used to fabricate a membrane. Single-gas
permeation measurements were subsequently performed. Fig. 2-1 schematicallyshows the theoretical modification of zirconium tetra butoxide with acetylacetonate
and the subsequent formation of an acac--modified Si-O-Zr network. Theoretically,
acetylacetone is a bidentate ligand with two sites through which it can bond totransition metal zirconium ions in order to form a normal covalent bond withzirconium through one of the oxygen atoms and a coordinate covalent bond by
donating a free pair of electrons from the other oxygen atoms that occupy the empty
d-orbital. This theory assumes that acetylacetone approaches the zirconium ions in its
enol tautomer [36]. Furthermore, the thermal stability of the acetylacetonate
modifying ligand was examined, and the pore modification effect that the addition ofacetylacetonate ligands and heat-treatment exerted on the pore structure and on thegas permeation characteristics were explored. Activation energies were used tocompare the SiO2-ZrO2-acac2 (SZa2) membrane performance with that of both apolymer membrane and a ligand-modified TiO2-ZrO2 membrane.

45

Hiroshima University Doctoral Dissertation

H9C4O

Zr

Tautomerism

O

O

O

OC4H9

H

O
O

OC4H9

- H9C4OH

OC4H9Keto Tautomer

H9C4O

Enol TautomerO

ZTIB Modification

Zr

OC4H9

OC4H9

-Si-O-Zr-acac Formation
OEt
EtO

Si

acac
OEt

+

OEt

BuO

Zr

OBu

OH

+ H2O

OBu

O

Zr

Zr

acac

O

O

Si

Si

O Zr

OBu

O

Si

O

OEt

SiO2-ZrO2-acac

Fig. 2-1 Theoretical formation of the -Si-O-Zr-acac linkages in a polymer sol

2.2 Experimental
2.2.1 Preparation of SiO2-ZrO2-acac2 sol and membrane fabrication
A 5 wt% modified SiO2-ZrO2 (1:1) polymeric sol (referred to as SZa2 sol in later
sections) and membranes were prepared and fabricated according to the following
procedures. Modified SiO2-ZrO2 (1:1) sol was prepared in two stages. In the firststage, an 80% pure zirconium (iv) tert-butoxide (Aldrich) dissolved in ethanol(Aldrich) was modified by reacting with 99% pure acetylacetone (Aldrich)(ligand/ZrTB molar ratio 2) for one hour in the proportions listed in Table 2-1. In thesecond stage, a solution of 98% pure tetraethoxysilane (Aldrich) in ethanol was cohydrolyzed with acetylacetone-modified zirconium (iv) tert-butoxide in a molar ratioof 1:1 using deionized water (water/ZrTB molar ratio 4) and hydrochloric acid (37%,Nacalai Tesque) as a catalyst (acid/alkoxide molar ratio 0.25). Hydrolysis and polycondensation were carried out by stirring the mixture at 500 rpm for more than 12hours at room temperature.
In a similar manner, without pre-modification of zirconium (iv) tert-butoxide, a 5
wt% unmodified SiO2-ZrO2 (1:1) polymeric sol (referred to as SZ sol in later
sections) was prepared in a process that required only one stage. The chemical
reagents and their respective proportions are also summarized in Table 2-1. Dried gelswere prepared by drying the sols slowly at a temperature of about 50 oC for a few
days.
46

Hiroshima University Doctoral Dissertation

Table 2-1. Composition of chemical reagents and precursors for 5 wt% polymeric sol
preparation
Amount of Sol Precursors and Reagents [g]Sol Type

Water

HCl

(WR=4)(AR=0.25)

94

0.61

0.207

0

94

0.61

0.207

1.7

TEOS

ZrTB

EtOH

0:1

1.77

4.06

2:1

1.77

4.06

(Ligand/Alkoxide)

Hacac

SiO2-ZrO2-acac2 membranes were fabricated by coating 3 graded layers ofdifferent suspensions/sols onto cylindrical α-alumina porous supports (60% porosity;
1.2 µm pore size; outer diameter, 1 cm; and, length, 10 cm; Nikkato Corporation,
Japan) connected to non-porous supports that were closed and open at either end. The
first layer was prepared by coating a 10 wt% suspension of α-alumina particles(diameter 2-3 µm and 0.2 µm) onto a colloidal SiO2-ZrO2 sol for use as a binder
prepared by ultrasonification. The preparation of colloidal SiO2-ZrO2 sol has been
described elsewhere [28]. Each coating was repeated 6 times, and each coating was
fired to a temperature of 550 oC for 15 – 20 min to solidify and bind the particle layer
to the support and to cover the larger pores of the porous alumina support.
Next, pure SiO2-ZrO2 (1:1) sol diluted to 0.5 wt% with deionized water wascoated onto the α-alumina particle layers using a hot coating method to form anintermediate layer, where the membrane was pre-heated to about 200 oC, and thencoated almost immediately with the SiO2-ZrO2 sol followed by firing at 550 oC for 15– 20 min to bind the intermediate layer to the particle layer. This process was repeatedseveral times until large pores and defects from the particle layer were covered and anarrow pore size distribution suitable for the separation layer coating was achieved.
The pore size distribution was obtained by nanopermporometry [39]. The top
separation layer was comprised of the acetylacetonate-modified SiO2-ZrO2 (1:1)
polymeric sol (derived from diluting the as-prepared sol to 0.5 wt% in ethanol) coatedonto the SiO2-ZrO2 intermediate layer followed by firing under a 600 ml min-1
nitrogen gas flow at 250 oC for 15 min in a tube furnace. After this process was
repeated 6 – 8 times, the membrane fabrication process was concluded, and themembranes were ready for single-gas permeation measurements. The same procedure47

Hiroshima University Doctoral Dissertation

was followed for the fabrication of an unmodified SiO2-ZrO2 (1:1) membrane thatwas used for comparison.
2.2.2 Characterizations of the SiO2-ZrO2-acac2, the SiO2-ZrO2 sol, the gel, andthe membranes
The particle sizes of the as-prepared sols were analyzed at 25 oC via dynamic lightscattering (DLS) using a Malvern Zetasizer Nano ZS (Malvern Instruments Ltd.). The
presence and intensities of different functional groups were analyzed via AttenuatedTotal Reflection-Fourier Transform Infrared Spectroscopy (ATR-FTIR, FTIR-4100,
JASCO, Japan). This analytical method was also used to track the thermal stability ofligands. Thermal stability of the acetylacetonate ligands in modified SiO2-ZrO2 wasanalyzed using thermogravimetry (TG, TG-50 Shimadzu Co., Japan) under a nitrogengas flow of 80 ml min-1 and a heating rate of 10 oC min-1. The cross-section andsurface morphologies were examined by Field Emission-Scanning Electron
Microscopy (FE-SEM, Hitachi S-4800, Japan). Prior to examination, shared samplesof the membranes were attached to sample holders via carbon tape and vacuum-driedat 50 oC for 48 hours.
2.2.3 Single-gas permeation measurements
A schematic representation of the gas permeation measurement setup is shown in
Fig. 2-2. Prior to the gas permeation measurement, a membrane was fitted into its
module immediately after fabrication and placed inside the furnace in the gas
permeation measurement rig at a temperature of 150 or 200 oC under helium flow of
100 ml min-1 for about 12 hours. This was done to ensure the removal of water
adsorbed into the membrane pores after fabrication. To test for the molecular sieving
properties of the membranes, high-purity gases (H2, He, CO2, N2, CH4, CF4, SF6 inthat order) were flowed through molecular sieves to remove humidity and were thenfed to the feed side of the membrane module at 100 kPa gauge pressure attemperatures ranging from 50 – 200 oC. Permeate side pressure was kept atmosphericand permeate gas flow was measured using a bubble film flow meter (HORIBASTEC, Japan). Temperature dependence of gas permeance was examined by varying
the temperature during the flow of a single gas and measuring the permeance flow at48

Hiroshima University Doctoral Dissertationeach desired temperature. It should also be noted that permeate gas flow was only
recorded after a steady state was attained.
Membrane heat-treatment was done by heating the membrane fired at 250 oC in
the membrane module to a higher temperature (300, 400 or 550 oC) while flowing N2
gas through the membrane. The flow of N2 was maintained until a steady state wasreached and the single-gas permeance was subsequently measured at 200 oC. This
procedure was repeated for each heat-treatment temperature.
4
2

5

6

1. Gas cylinder
2. Drier
3. Pressure controller
4. Pressure gauge
5. Mass flow controller
6. Stop valve

3

1

4
2

5

6

3

7. Membrane module
8. Membrane
9. Furnace
10. Temperature controller
11. Back pressure valve
12. Bubble film meter

1
4

11

12

8

7

12

9
10

Fig. 2-2 Schematic diagram of single-gas permeation measurement.

2.3 Results and discussion
2.3.1 Sol-gel characterization
2.3.1.1 Chemical modification of a ZrTB precursor
Fig. 2-3 (a) shows the ATR-FTIR spectra of freshly prepared modified and purezirconium tetrabutoxide precursors, measured directly on an ATR prism window that
ranged from 4,000 to 500 cm-1 compared with that of acetylacetone (Hacac). Peaks
ranging from 1,700 to 1,500 cm-1 are magnified for clarity in Fig. 2-3 (b). The two
peaks that appear around 1,600 cm-1 and 1,525 cm-1 were assigned to the enol C=O
49

Hiroshima University Doctoral Dissertation

and C=C bonds that are characteristic of the acetone species [36, 38]. The presence of
the C=C bond supports the assumption that Hacac approaches ZrTB in its enol
tautomer. These peaks were not detected for the unmodified precursor and Hacac,
which presented only the C=O (1615 cm-1) peak of the keto tautomer. This clearly
indicates the introduction of the acetylacetonate group to the chemical structure of
zirconium tetrabutoxide. Furthermore, peaks assigned to the butoxide groups (-CH2, CH3) in the range of 3,000 to 2,750 cm-1 were observed for both modified and purezirconium tetrabutoxide precursors, although different peak heights were exhibited.
The -CH2, -CH3 peaks observed for the modified zirconium tetrabutoxide precursor
were shorter than those observed for the unmodified precursor with representative
absorbance peak ratios, A2960/A1400, of 0.29 and 0.56 for modified ZrTB and pure
ZrTB, respectively. A2960 represents the absorbance peak height of a -CH2 bondnormalized by a common Zr-O-C peak observed at around A1400 (1,400 cm-1). This
indicates the replacement of some butoxy groups by the acetylacetonate groups thatwould have been otherwise present without modification.
-CH2

C=O (Keto)-CH3

C=O (Enol) C=C

Absorbance [a. u.]Zr(OBu)3acac

Zr(OBu)4

Absorbance [a. u.]

Zr-O-C
Zr(OBu)3acac
Zr(OBu)4

Hacac
Hacac
4000 3500 3000 2500 2000 1500 1000 500-1Wavenumber [cm ]

1700

(a)

1650
1600
1550
-1
Wavenumber [cm ]
(b)

1500

Fig. 2-3 ATR FT-IR spectra of ZrTB (zirconium tetrabutoxide) modification byacetylacetonate ligands ranging (a) from 4,000 to 500 cm-1 and (b) from 1,700 to
1,500 cm-1.
Fig. 2-4 also shows highly magnified ATR-FTIR spectra ranging from 1,700 to1500 cm-1 that compare modified zirconium tetrabutoxide, modified silica-zirconia,
50

Hiroshima University Doctoral Dissertationand pure unmodified silica-zirconia sols. The C=O and C=C peaks characteristic ofacetylacetonate ligands observed for the modified zirconium tetrabutoxide precursor
were also detected for the modified silica-zirconia sol. This is an indication that theacetylacetonate-zirconium chelation structure was unaffected by the introduction ofsilica. Furthermore, the acetylacetonate characteristic peaks (1,600 cm-1 and 1,525 cm1) in the modified silica-zirconia sol were absent from the unmodified sol, whichsuggests that acac- ligands are hydrolytically stable, and were not hydrolyzed by theaddition of water under acidic conditions. This was previously observed duringhydrolyzation of the acac-ZrTB coordinate complex polymer for several days [36].
The rate of ZrTB hydrolysis is much faster than that of TEOS, and one advantage ofmodifying the zirconium tetrabutoxide with a hydrolytically stable ligand prior to
mixing with TEOS was that the hydrolysis rate of ZrTB was slowed considerably sothat during co-hydrolysis of the ZrTB and TEOS mixtures, the rates of hydrolysis
were similar, which improved the -Si-O-Zr- composition degree [32].

Absorbance [a. u.]

C=O

C=C

Zr(OBu)3acac
SiO2-ZrO2-acac
Unmodified SiO2-ZrO2

1700

1650
1600
1550
-1
Wavenumber [cm ]1500

Fig. 2-4 ATR-FTIR spectra ranging from 1,700 to 1,500 cm-1 for modified zirconium
tetrabutoxide, modified silica-zirconia, and unmodified silica-zirconia sols.
2.3.1.2 Thermal analysis of acac- ligands in SiO2-ZrO2-acac2 gels
The hydrolytic stability of acac in the sol phase was expected, having been
established in section 2.3.1.1, and this section discusses the thermal stability of the51

Hiroshima University Doctoral Dissertationacac- ligands in the SZa2 powder. The thermogravimetric (TG) curves of pure SiO2ZrO2 and SiO2-ZrO2-acac2 gels fired under a N2 flow at temperatures as high as 1,000oC (ramping rate: 10 oC min-1) and the ATR-FTIR spectra of SZa2 powders fired

under N2 at different temperatures are shown in Fig. 2-5. In Fig. 2-5 (a), the TG
profile of SZa2 powder can be divided into 3 sections according to their different
slopes. Section i (100-300 oC) can be assigned to the decomposition of unhydrolyzed
butoxy groups and to the partial decomposition of acetylacetonate linkages based onweight loss. This is supported by the ATR-FTIR profile of powders fired at different
temperatures (Fig. 2-5 (b)).
In section ii (300 and 500 oC), pure SZ powder exhibited no further weight loss
beyond 300 oC, indicating that weight loss was due only to the decomposition ofunhydrolyzed butoxy groups. For SZa2 powder, further weight loss was observed to aresidual weight ratio that reached as high as 0.7, indicating the complete
decomposition of acac ligands. This also appears in Fig. 2-5 (b) as the characteristic
C=O and C=C bonds were decreased with temperature. Beyond 500 oC, as shown insection iii, free carbon from the decomposed acac- ligands was released into the pore
network, which resulted in no observed weight loss in this region. The dark
appearance of the residual powders after TG measurement supports the theory of theencapsulation of free carbon in the SiO2-ZrO2 network, which is similar to carbonencapsulated nickel nanoparticles prepared by the instant pyrolysis of nickel
acetylacetone reported by Chen et al. [40], as well as the C/C-ZrC composites
fabricated by zirconium acetylacetone pyrolysis [41]. The continued weight loss of
SZa2 powder beyond 300 oC indicates that an acac--modified SiO2-ZrO2 layer can be
fabricated even when fired to a temperature as high as 300 oC.

52

Hiroshima University Doctoral Dissertation

1.0
0.9

-CH3

C=C

Si-O-SioUnmodified SiO2-ZrO2

0.8

Modified SiO2-ZrO2

0.7
i

0.6

ii

iii

Absorbance [a. u.]

Residual weight [-]

-CH2

500 C
o

400 C
o

300 C
o

250 C
Unfired0.5

0200 400 600 800 1000 3500 3000 2500 2000 1500 1000 500-1oWavenumber [cm ]Temperature [ C]

(b)

(a)

Fig. 2-5 TG curves of pure SiO2-ZrO2 and SiO2-ZrO2-acac2 gels under N2 attemperatures as high as 1,000 oC (a) and Powder ATR-FTIR spectra of SiO2-ZrO2acac gels fired under a N2 atmosphere at different temperatures (b)Fig. 2-5 (b) shows the ATR-FT-IR spectra of SZa2 powders fired under N2 at
different temperatures. The peak at around 1,500 cm-1 was assigned to the C=C bondcharacteristics of an acetylacetonate ligand, as mentioned in a previous section [36,38]. The presence of this peak is indicative of the fact that the acac-Zr chelation
structure was unaffected by the change of state from sol to gel. The peak stretching
between 1,100 and 900 cm-1 was assigned to the -SiOSi- ring while the -SiOH group
was assigned to the peak at about 800 cm-1. The peaks between 2,800 and 3,000 cm-1
were assigned to the unhydrolyzed ethoxy and butoxy groups [42-44]. The -SiOSipeaks were detected irrespective of the firing temperature, which is indicative of theformation of silica networks by hydrolysis and by the condensation of the ethoxy
groups of TEOS. The C=C peaks began to disappear as the firing temperature
increased the AC=C/A-SiOSi- absorbance peak ratio decreased as the firing temperature
increased.
The simultaneous reduction and disappearance of -CH2, -CH3 and -SiOH peaks attemperatures higher than 300 oC was confirmed along with the subsequent -SiOSipeak shift and broadening. This can be attributed to the condensation of ethoxy,
butoxy and silanol groups and the formation of more -SiOZr- connections. The
broadening and red shift of the -SiOSi- peak in a silica-zirconia composite FT-IR
spectrum is indicative of more -SiOZr- connections, and the inclusion of Zr4+ ions53

Hiroshima University Doctoral Dissertation

into the -SiOSi- ring caused a ring deformation and a resultant peak shift andbroadening [45]. Fig. 2-6 shows the XRD patterns of pure SZ-derived and SZa2derived powders after firing at 700 oC under N2. After firing at 700 oC, the tetragonal
ZrO2 peaks at 2θ = 30o and 50o indicated that the pure SZ-derived powder had begun
to show an aggregation of ZrO2 crystallites, whereas SZa2-derived powder remained
amorphous without aggregation of ZrO2 crystallites indicating a better -Si-O-Zrcomposition degree.

 Tetragonal ZrO2o

700 C, N2

Intensity [a. u.]

Unmodified SiO2-ZrO2



SiO2-ZrO2-(acac)2

10

20

30

o

40

50

60

2 [ ]

Fig. 2-6 XRD patterns of pure SZ- and SZa2-derived powders before and after firing
at 700 oC under N2
2.3.2 Effect of acetylacetonate ligands on the pore size controllability of SiO 2ZrO2 (1:1) membranes
Fig. 2-7 shows the kinetic diameter dependence of single-gas permeance at 200 oC(a) and dimensionless permeance based on He (b) for SZ and SZa2-derived
membranes prepared at 300 oC under a N2 atmosphere. Pure SZ-derived membranes
showed a higher gas permeance compared with that of SZa2-derived membranes, but
He and H2 selectivity over other gases was approximately the same as the Knudsen
ratio. For example, a pure SZ-derived membrane showed a H2 permeance of 1.4 x 10 6

mol m-2 s-1 Pa-1 with a H2/SF6 permeance ratio of 11.
The SZa2-derived membrane showed a trend similar to that of the SZ-derived

membrane; He (0.26 nm) permeance was lower than that of H2 (0.289 nm), and N2
54

Hiroshima University Doctoral Dissertation(0.364 nm) permeance was lower than that of CH4 (0.38 nm). This suggests a network
pore size large enough for the separation of He-H2 and N2-CH4 and Knudsen diffusionthat is dominated by the permeation properties [29, 30]. However, the SZa2-derived
membrane showed a low SF6 permeance of 1.3 x 10-10 mol m-2 s-1 Pa-1, with a high
H2/SF6 permeance ratio of 7,600. This shows that the chelation of Zr4+ with acacligands uniformly modified the pores of the SZa2-derived membrane for molecularsize-dependent permeation resistance. SZa2-derived membranes, therefore, showpromise for the separation of H2/large organic gases, because the molecular size ofSF6 (0.55 nm) is similar to that of organic gases in benzene-cyclohexane and tolueneMCH systems (0.585 and 0.6 nm respectively) [46].
-5

1

-6

-2

-1

10

10
He

H2

Dimensionless permeance [-]-1

Permeance [mol m s Pa ]

10

CO2
N2CH4
CF4

-7

10

SF6

-8

10

-9

10

Modified SiO2-ZrO2
Unmodified SiO2-ZrO2

-10

10

0.2

0.3
0.4
0.5
Kinetic diameter [nm]
(a)

0.6

0

10

He

H2

CO2
CH
N2 4

CF4

SF6

-1

10

-2

10

-3

10

Unmodified SiO2-ZrO2-4

10

Modified SiO2-ZrO2

0.2

Knudsen

0.3
0.4
0.5
Kinetic diameter [nm]
(b)

0.6

Fig. 2-7 Molecular size dependence of gas permeance (a) and dimensionlesspermeance based on He (b) at 200 oC of SiO2-ZrO2-acac2 and pure SiO2-ZrO2
membranes all prepared at 300 oC under a N2 atmosphere.
Fig. 2-8 shows the FE-SEM cross-sectional morphologies (x 50,000) of SZa2 (a)and pure SZ (b)-derived membranes. A thin separation layer of acac- modified silicazirconia was successfully fabricated, although it was indistinguishable from theintermediate layer at this magnification. The α-Al2O3 particle layer (average 1 µm)
was clearly distinguishable as a compact impenetrable support for the intermediate
and top layers (< 400 nm) of both membranes. It should be noted that the FE-SEM
images were taken after gas permeation experiments were performed. The laminationof the SZ and SZa2 that topped the SiO2-ZrO2 intermediate layers was retained after55

Hiroshima University Doctoral Dissertation

gas permeation experiments, which suggests a favorable interaction between silicabased top layers containing organic groups and the SiO2-ZrO2 layer. Therefore, theaddition of the acac- ligands to SiO2-ZrO2 had no structural effect on the top layer
during long-term use.
(b)

(a)

SiO2-ZrO2 Top Layer / SiO2ZrO2 Intermediate Layer

SiO2-ZrO2-acac 2 Top Layer /SiO2-ZrO2 Intermediate Layer

α-Al2O3 Particle Layer

α-Al2O3 Particle Layer

(a)

(b)

Fig. 2-8 FE-SEM micrographs of the cross-section for SiO2-ZrO2-acac2 (a) and pure
SiO2-ZrO2 (b) membranes.
The schematic diagram in Fig. 2-9 (a) summarizes the effect of the acac- chelating
ligands on the Si-O-Zr network. The chelation of acac- ligands to Zr+ sites creatednetwork spaces small enough to prevent permeation by large gases yet large enoughto allow small gases to permeate, which created better molecular sieving properties.
Hypothetically, as Fig. 2-9 (b) illustrates, interlayer condensation may occur between
two successive layers when the free chelation sites in acac- ligands on one layer
chelate with the Zr+ sites on another layer. Such interlayer interactions like the selfassembly of thiourea-crosslinked graphene oxide [47], topotactic dehydrationcondensation of layered silicate into a zeolite [48], and hydrogen-covalent bondinteraction between hybridized layers of PVA/TEOS [49] have been studied andproven. Characterization to prove these types of interaction is currently beyond thescope of this study and will require further inquiry. This phenomenon was notapparent in the SEM micrographs even at very high magnification.

56

Hiroshima University Doctoral Dissertation

SF6
H2

Zr

O

Si

SF6

H2

O
Si

O
Zr
O
Si
O

Zr

Si

O

O
O

O
Si

O

O

Zr

Zr

O

O
Si

O

O

O

O

O

O

Si
O

O

Zr

O

O

O

Zr

Zr

O

Si

O
O

Zr

O O
Zr
O
O

Zr

Si

O

Si
O
O

O

O

O
Zr

O
O
Zr
O

Zr

(a)

O
Zr O
O
O

Zr
O O
Si

O

O
O
Si
O O O
Zr

O

O
Zr

O
O
O
O
O

Zr

O
Zr O
O
O

Zr
O O
Si

O

O
O
O

Zr

O
Si
O
O O O
Zr O
O
O
Zr
O

Si

O
Si

O
Si

O
Si

O

O
O

O

Zr
O

O

Zr
O
O
Zr

O

O
Si

Zr
O

O

Interlayer
coordination
between Zr+
sites

O
Si

H2

SF6

(b)

Fig. 2-9 Schematic image of the intra-particle pore modification effect ofacetylacetone chelating ligands forming coordination complexes with Zr atoms andthe influence on gas permeation ability; (a) network pore modification theory; (b)interlayer coordination theory.

57

Hiroshima University Doctoral Dissertation

2.3.3 Heat-treatment effect on pore size and gas permeation properties of SiO 2ZrO2-acac2 membranes
The thermogravimetric results indicate different sections on the TG profile asdiscussed in section 2.3.1.2. Therefore, it was interesting to evaluate the gas
permeation properties of SZa2-derived membranes fired at different temperatures andthe effect that firing temperature exerted on the membrane pore sizes. Fig. 2-10 shows
the molecular size dependence of single-gas permeance at 200 oC (a) and thedimensionless permeance based on He (b) for SZa2-derived membranes heat-treated attemperatures ranging from 250-550 oC under a N2 flow, as mentioned in section 2.2.3.
Fig. 2-11 shows the values for single-gas permeance and permeance ratios at 200 oC
as a function of heat-treatment temperatures. Gas permeation properties changed asmembranes were heat-treated at higher temperatures. A membrane fired at 250 oC
showed a low H2 permeance of 1.8 x 10-7 mol m-2 s-1 Pa-1, but its H2/CH4 selectivity of
50 was higher than with heat-treatment ranging from 300-550 oC. It should be notedthat the permeance for large gases such as CF4 and SF6 was below the detection levelsof 10-11 mol m-2 s-1 Pa-1, and the H2/SF6 permeance ratio was assumed to be within theacceptable range of 104-105, which suggests that these membranes had very dense
structures.

1
1

-5

-6

10

H2
CO2
N2 CH4
CF4-2

Dimensionless permeance [-]

He

10

-1

-1

Permeance [mol m s Pa ]

10

SF6

-7

10

-8

10

-9

10

-10

10

0.2

o

Fired at 250 Co
HT at 300 Co
HT at 400 Co
HT at 550 C

0.3

0

10

0.5

CO2
N2 CH4

CF4

10
-2

10

10

o

-3

10

0.2

Fired at 250 Co
HT at 300 Co
HT at 400 Co
HT at 550 C
Knudsen

0.3

0.4

10
0.5

Kinetic diameter [nm]Kinetic diameter [nm]
(a)

(b)

58

10

SF6

10

10

0.6

1

H2

-1

-4

0.4

He

0.6

Hiroshima University Doctoral Dissertation

Fig. 2-10 Comparison of the kinetic diameter dependence of single-gas permeance (a)and dimensionless permeance based on He (b) at 200 oC for SiO2-ZrO2-acac2
membranes fired at 250 oC and heat-treated at different temperatures (300-550 oC).
5

10
Selectivity [-]

4

10

2

10

H2/CH4

1

10

0

10-5
10
-1

Permeance [mol m s Pa ]H2/SF6

3

10

CO2/CH4

-6

10

-2

-1

H2
-7

10

-8

CO2

-9

CH4

10
10

SF6
-10

10

250

300 350 400 450 500 550oHeat treatment temperature [ C]

Fig. 2-11 Heat-treatment temperature (300-550 oC) dependence of gas permeance andselectivity (H2/CH4, H2/SF6, CO2/CH4) at 200 oC for SiO2-ZrO2-acac2 membranes
fired at 250 oC under a N2 atmosphere.
After heat-treatment at 300 oC, developments such as a drastic increase in gas
permeance and a decrease in gas selectivity were observed (H2 permeance: ~10 -6 mol
m-2 s-1 Pa-1, H2/SF6 permeance ratio: ~8,000). At 300 oC and above, the changes to gas
permeance and permeance ratios were more subdued. It should be noted that H2
selectivity over large gases such as CF4 and SF6 remained above the Knudsen-basedvalues until the membranes were heat-treated at 400 oC. However, a membrane heattreated at 550 oC showed gas permeance on the order of values like that of theintermediate layer, and H2 selectivity over large gases such as CF4 and SF6
approached Knudsen ratios.
59

Hiroshima University Doctoral Dissertation

The change in dimensionless permeance with increases in the firing temperature
became more pronounced as molecular size increased from H2 to SF6. This could haveresulted from heat-treatment at higher temperatures causing the acac- ligands todecompose within the modified pores, which would make the pores larger, which isconsistent with thermogravimetric results. The increased pore size increases theaccess to large gases. High-temperature heat-treatment of SZa2-derived membranes
fired at 250 oC, therefore, increases the pore size due to the decomposition of acacligand groups. This decomposition of organics behavior as heat-treatment temperature
increases is consistent with the templating effect summarized in Fig. 2-12. The
chelating ligands attached to the Zr+ sites occupied the network pores in the CH3-CHCH2-CH-CH3 section of the acac- ligand. As heat-treatment temperature increased,these organic sections began to decompose opening the pores for gas passage, which
increased the gas permeance, but reduced gas selectivity as larger gases were able toaccess the network pores. The occupation of network pores can be assumed to be both
intraparticle and interparticle. Intraparticle pores exist within individual particles
created by spaces between the atomic networks while interparticle pores are thespaces that exist between particles. The interparticle pore-containing ligands
decomposed before the intraparticle pore-containing versions, which was inferred by
the fact that the membranes continued to show gas selectivity higher than Knudsen
ratios even after heat-treatment at 550 oC. This could have been because theintraparticle pore-containing ligands are relatively shielded from heat by thesurrounding Si-O-Zr network, much like the encapsulation of carbon in nickel
nanoparticles [40].
SiO2-ZrO2-acac2

Zr

O

Si

O

O

O

Si

Zr
O

Si

O

Zr
O

O

O

O

Zr

Heat treatment
Decomposition of acac
ligands

Zr

O

60

O

Si

O

Zr

O

Si

Zr

Si
O

O

O

O
Si

Zr

O

O
Si

Zr

Zr

O

Si

Zr

O

O

Si

Zr

O
Zr

O

O
Si

O

O
Si

O

Zr

O

Hiroshima University Doctoral Dissertation

Fig. 2-12 Schematic diagram of the effect of heat-treatment on the pore size and gas
permeation ability of a typical SiO2-ZrO2-acac2 membrane.
Fig. 2-13 shows the single-gas permeance and permeance ratios as functions oftemperatures between 50 and 200 oC for SZa2-derived membranes fired at 250 oC (a)(He, H2, CO2, N2, CH4) and heat-treated at 300 oC (He, H2, CO2, N2, CH4, SF6) and550 oC (b) (He, H2, CO2, N2, CH4, CF4, SF6). The temperature dependence of
permeance and the activation energies of gases were fitted and calculated using themodified gas translation model (Eq. 2-1). The model showed a very good fit with theexperimental data and should be regarded as a viable representation of gas-membrane
behavior.
𝑃𝑖 =

𝑘𝑜,𝑖

𝐸

(2-1)

𝑝,𝑖
𝑒𝑥𝑝 (− 𝑅𝑇
)
𝑀 𝑅𝑇

√ 𝑖

For membranes fired at 250 oC, the permeance of all gases (He, H2, N2, CH4)
increased with increases in temperature, suggesting an activated diffusion mechanism.
The exception to this was CO2, which showed an increase in permeance as thetemperature decreased, which suggests a surface-diffusion mechanism. Gas
selectivities (H2/CH4, CO2/CH4, CO2/N2) increased as temperature decreased. In
particular, for gas permeance measured at 50 oC, H2/CH4 and CO2/CH4 selectivities
reached 100 and 60, respectively.

61

Hiroshima University Doctoral Dissertation

Selectivity [-]

10

4

10

3

10

2

10

1

200

150

100

50

CO2/N2

0

4

200

10

3

H2/SF6

10

2

10

1

10

He

10

10

-9

CO2/N2

-6

10

4

10

3

10

2

10

1

200

CO2

N2

CH4

50

CO2/N2

CO2/CH4

0

10
-5
10
10

-6

10

-7

10

-7

10

-8

10

-8

10

-9

10

-9

N2

100

H2/CH4

H2

He

CO2

H2

150

H2/SF6

N2

He

CH4

CF4

SF6

CH4

4

-10

2.0

50

CO2/CH4

CO2

10

100

10
-5
10

-7

-8

150

H2/CH4

H2

-2

-1

-1

Permeance [mol m s Pa ]

10

0

10-6
10

10

Temperature [ C]

Temperature [ C]

H2/CH4
CO2/CH4oooTemperature [ C]

-10

2.4

2.8

10
3.2 2.0-11000/T [K ]
(a)

SF6

10

3

-10

2.4

2.8
-1

1000/T [K ](b)

10
3.2 2.0

2.4

2.8

3.2

-1

1000/T [K ](c)Fig. 2-13 Temperature dependence of gas permeance at temperatures ranging from50-200 oC for SiO2-ZrO2-acac2 membranes fired at 250 oC (a) and heat-treated at 300oC (b) and 550 oC (c) under N2.
Membranes heat-treated at 300 oC showed the same trend as the original

membrane fired at 250 oC with gas permeance (He, H2, N2, CH4) increasing as thetemperature increased with the noted exception of CO2 that exhibited a surfacediffusion mechanism. However, the slopes of the temperature dependence of gas
permeance followed different trends, and membranes fired at 250 oC showed
increases in the slope as the molecular size increased, but the slope decreased withmolecular size for membranes heat-treated at 300 oC.
When heat-treated to 550 oC, the permeance of gases (N2, CF4, SF6) decreased as
temperature increased except for small gases (He, H2 and CH4) that continued to show
smaller activation energies (He: 5.7, H2: 5.1, CH4: 5.07 kJ mol-1) compared with 250oC-fired membranes (He: 12, H2: 13, CH4: 20 kJ mol-1). The transition in membrane

permeation characteristics can be clearly observed in the gas selectivity plot. As amembrane is heat-treated to 300 oC, there is a drastic decrease in gas selectivity(H2/CH4, CO2/CH4, CO2/N2), which is reduced even further for membranes heattreated at 550 oC with no considerable variation in permeation temperature. Based on
62

10
2
10
1

10
0
10

Hiroshima University Doctoral Dissertationthese results, it can be concluded that the pore morphology in a 250 oC-fired
membrane is such that gas separation with high selectivity can be expected by altering
the operating temperature, whereas membranes heat-treated to 550 oC can accomplish
no gas separation at different operating temperatures since the pores are not selective
for gases.
Fig. 2-14 shows the relationship between the activation energies of gases (He, H2,
N2, CH4, SF6) and the kinetic diameter for SZa2-derived heat-treatment ranging from250-550oC, for isoeugenol-modified TiO2-ZrO2-derived [50] and polyimide

membranes [52, 53]. For SZa2-derived membranes heat-treated at 250 oC, isoeugenolmodified, TiO2-ZrO2-derived, and polyimide membranes, the activation energy
increases as the kinetic diameter increases, as indicated in Fig. 2-14 via the data points
in the solid-line circle. These results indicate a similar gas permeation mechanism inthese membranes, although the polymer membrane showed higher activation energies
across all gases. This behavior (increased activation energy of gas permeance withkinetic diameter of gases) is a solution-diffusion mechanism that is accepted as thetransport mechanism in polymer membranes [54, 55]. As for the polyimide
membranes studied by Villaluenga [52] et al. and Hirayama et al. [56], the interaction
between the polymer and each gas (He, H2, N2, CH4) corresponding to the heat ofsorption, ΔHs, is very small so that the high activation energies of permeation
observed should be attributed to diffusivity rather than to solubility, even if thesolubility coefficient is thermodynamic in nature. The thermodynamic nature of gas
diffusivity in polymer membranes can be explained by the fact that polymer chainscorresponding to voids vibrate relative to temperature so that activation energy
increases with kinetic diameter. Thus, based on the observed trend, this phenomenon
could also be attributed to SZa2-derived membranes fired at 250 oC.

63

Hiroshima University Doctoral Dissertation-1

Activation energy, Ep [kJ.mol ]

35

CH4

30
N2

25
20

o

250 C Firedo
300 C HTo
400 C HTo
550 C HTTiO2-ZrO2-ISOH

H2

Polyimide

15 He
10
5
0
-5
0.25

0.30

0.35

0.40

0.45

0.50

0.55

Kinetic diameter [nm]Fig. 2-14 Kinetic diameter dependence of gas activation energies (He, H 2, N2, CH4,
SF6) for SiO2-ZrO2-acac2 membranes fired at 250 oC and heat-treated at 300, 400 or
550 oC under N2 compares with that of isoeugenol (ISOH)-modified TiO2-ZrO2 [50]and polyimide [52, 53] membranes.
In contrast, as indicated by data points in the dashed-line circle in Fig. 2-14, thekinetic diameter dependence of the apparent activation energy of SZa2-derived
membranes heat-treated at 300-550 oC showed a trend whereby activation energy
decreased with kinetic diameter. This was apparent because heat-treatment at higher
temperatures decomposed the organic content in the modified network structure,
which resulted in enlarged network pores that allowed easier access to larger gases forthese membranes. Fig. 2-15 also compares the heat-treatment temperature dependenceof the apparent activation energies of He and N2, which clearly displays a sharper
reduction in the activation energy of N2 than that of He and suggests the creation oflarger pores. However, small gases (He and H2) showed higher activation energies
indicating the presence of pores with sizes in the range of the kinetic diameters of He
and H2 even after firing at higher temperatures. These two observations are evidenceof a bimodal pore structure. Therefore, the heat-treatment of membranes affects thepore structure enough to affect the permeation properties of the membrane.

64

Hiroshima University Doctoral Dissertation-1

Activation energy, Ep [kJ.mol ]

20
15
10
He

5
N2

0
-5
200

300

400

500

600
o

Heat treatment temperature [ C]
Fig. 2-15 Heat treatment temperature dependence of gas activation energies (He andN2) for SZa2-derived membranes
Fig. 2-16 shows the trade-off plot for the H2/SF6 permeance ratio as a function ofH2 permeance for two different membranes fired at 250 oC and heat-treated at 300,
400 or 550 oC. It should be noted that the SF6 permeance for a membrane fired at 250o

C was assumed to be on the order of 10-11, which is below the detection limit, and,

therefore, the H2/SF6 permeance ratio was assumed to be within a reasonable limit of
104 to 105. It was apparent that all H2/SF6 permeance ratios were above the Knudsen
ratio irrespective of heat-treatment temperature. The solid line in Fig. 2-16 is thetrade-off line that can be drawn to align with the trade-off points. A membrane withhigh H2/SF6 selectivity generally shows a low H2 permeance, and vice versa. Whenthe trade-off line is considered, membranes heat-treated at 300 oC provided afavorable trade-off point by maintaining a good balance between H2/SF6 (1,800-7600)
selectivity and H2 permeance (0.4-1 x 10-6 mol m-2 s-1 Pa-1). Although membranes
fired at 250 oC exhibited a higher H2/SF6 selectivity (> 18,000), that for a membraneheat-treated at 300 oC was not far below, but with almost 10 times the order of H2
permeance (1.8 x 10-7 mol m-2 s-1 Pa-1). Thus, a membrane heat-treated at 300 oC
delivers a better trade-off than other membranes with high H2 permeance and H2/SF6
selectivity.

65

Hiroshima University Doctoral Dissertation

10

6
o

250 C

10

5

10

4

10

3

10

2

10

1

H2/SF6 [-]

o

300 C

o

400 C
o

550 C

Knudsen

0

10 -7
10

10

-6

10
-2

-1

-5

-1

H2 permeance [mol m s Pa ]Fig. 2-16 H2/SF6 permeance ratio as a function of H2 permeance at 200 oC for
membranes fired at 250 oC and heat-treated at 300, 400 or 550 oC.

2.4 Conclusions
A chemical modification of SiO2-ZrO2 via the sol-gel method was successfully
accomplished by chelating acetylacetonate (acac-) with zirconium tetrabutoxide(ZrTB) prior to co-hydrolysis and condensation with tetraethoxysilane (TEOS), which
was then used to fabricate a gas separation layer by coating the acac--modified SiO2ZrO2 sol onto a SiO2-ZrO2 intermediate layer. Acac--modified SiO2-ZrO2-derived
membranes were then fabricated by firing under N2 at 250 oC followed by heattreatments that ranged from 300-550 oC under N2. Single-gas permeation experiments
were carried out to examine the effect of chelating agents on the resultant
modification of the pores, which affected the gas permeation properties of themembrane.
The acac--modified SiO2-ZrO2-derived membrane showed H2 permeance of 9.9 x
10 -7 with a H2/SF6 permeance ratio of 7,600, which was better than the results whenusing a pure SiO2-ZrO2-derived membrane (H2 permeance: 1.4 x 10 -6 mol m-2 s-1 Pa-1,
H2/SF6 permeance ratio: 11). When an acac--modified SiO2-ZrO2-derived membrane
prepared at 250 oC (H2 permeance: 1.8 x 10-7 mol m-2 s-1 Pa-1, H2/CH4: 100, CO2/CH4:60 at 50 oC) was heat-treated at 550 oC, the result was improved H2 permeance of 5.2
66

Hiroshima University Doctoral Dissertationx 10-6 mol m-2 s-1 Pa-1 with low permeance ratios (H2/CH4: 3, H2/SF6: 9) at 50 oC. The
250oC-fired acac--modified SiO2-ZrO2-derived membranes exhibited a gas

permeation mechanism similar to polymer membranes, and the 300 oC-heat-treatedversion showed the best trade-off between H2 permeance and the H2/SF6 permeanceratio.

Cited literature
1.

N. W. Ockwig, T. M. Nenoff, Membranes for hydrogen separation, Chem.
Rev. 107 (2007) 4078-4110.

2.

S. Adhikari, S. Fernando, Hydrogen membrane separation techniques, Ind.
Eng. Chem. Res. 45 (2006) 875-881

3.

H. Li, Z. Song, X. Zhang, Y. Huang, S. Li, Y. Mao, H. J. Ploehn, Y. Bao, M.
Yu, Ultrathin, molecular-sieving graphene oxide membranes for selectivehydrogen separation, Science 342 (2013) 95-98

4.

M. Kanezashi, D. Fuchigami, T. Yoshioka, T. Tsuru, Control of Pd dispersion
in sol-gel-derived amorphous silica membranes for hydrogen separation athigh temperatures, J. Membr. Sci. 439 (2013) 78-86

5.

S-J. Ahn, A. Takagaki, T. Sugawara, R. Kikuchi, S. T. Oyama, Permeation
properties of silica-zirconia composite membranes supported on porousalumina substrates, J. Membr. Sci. 526 (2017) 409-416.

6.

S.-W. Chuang, S. L.-C. Hsu, Y.-H. Liu, Synthesis and properties of fluorinecontaining polybenzimidazole/silica nanocomposite membranes for proton
exchange membrane fuel cells, J. Membr. Sci. 305 (2007) 353-363

7.

Y. Zhang, J. Sunarso, S. Liu, R. Wang, Current status and development of
membranes for CO2/CH4 separation: a review, Int. J. Green Gas Control. 12(2013) 84-107

8.

M. A. Carreon, S. Li, J. L. Falconer, R. D. Noble, Alumina-supported SAPO34 membranes for CO2/CH4 separation, J. Am. Chem. Soc. 130 (2008) 54125413

9.

C. E. Powell, G. G. Qiao, Polymeric CO2/N2 gas separation membranes for thecapture of carbon dioxide for power plant flue gases, J. Membr. Sci. 279(2006) 1-49

67

Hiroshima University Doctoral Dissertation

10.

L. Zhao, E. Riensche, R. Menzer, L. Blum, D. Stolten, A parametric study of
CO2/N2 gas separation membrane processes for post-combustion capture, J.
Membr. Sci. 325 (2008) 284-294

11.

D. C. Nymeijer, T. Visser, R. Assen, M. Wessling, Composite hollow fiber
gas-liquid membrane contactors for olefin/paraffin separation, Sep. Purif.
Tech. 37 (2004) 209-220

12.

B. N. Nair, K. Keizar, W. J. Elferink, M. J. Gilde, H. Verweij, A. J. Burggraaf,
Synthesis, characterization and gas permeation studies on microporous silica
and alumina-silica membranes for separation of propane and propylene, J.
Membr. Sci. 116 (1996) 161-169

13.

S. Yun, J. H. Ko, S. T. Oyama, Ultrathin palladium membranes prepared by anovel electric field assisted activation, J. Membr. Sci. 369 (2011) 482-489.

14.

M. Kanezashi, Sol-Gel Derived Silica Membranes, Encyclopedia of
Membrane Science and Technology, John Wiley and Sons, 2013

15.

B. Ballinger, J. Motuzas, S. Smart, J. C. Diniz da Costa, Palladium cobalt
binary doping of molecular sieving silica membranes, J. Membr. Sci. 451(2014) 185-191

16.

R. M. de Vos, H. Verweij, High-selectivity, high-flux silica membranes forgas separation, Science. 279 (1998) 1710-1711

17.

R. M. de Vos, H. Verweij, Improved performance of silica membranes for gas
separation, J. Membr. Sci. 143 (1998) 37-51

18.

A. Peta Dral, J. E. ten Elshof, Organic groups influencing microporosity inorganosilicas, Micr. Meso. Mater. 267 (2018) 267-273.

19.

M. Kanezashi, Y. Yoneda, H. Nagasawa, T. Tsuru, K. Yamamoto, J. Ohshita,
Gas permeation properties for organosilica membranes with different Si/C
ratios and evaluation of microporous structures, AIChE J. 63 (2017) 44914498.

20.

N. K. Raman, C. J. Brinker, Organic “template” approach to molecular sieving
silica membranes, J. Membr.Sci. 105 (1995) 273-279.

21.

K. Yoshida, Y. Hirano, H. Fujii, T. Tsuru, M. Asaeda, Hydrothermal stability
and performance of silica-zirconia for hydrogen separation in hydrothermal
conditions, J. Chem. Eng. Jpn. 34 (2001) 523-530.

22.

S. Yun, S. T. Oyama, Correlations in palladium membranes for hydrogen
separation: a review, J. Membr. Sci. 375 (2011) 28-45.
68

Hiroshima University Doctoral Dissertation

23.

D. Uhlmann, S. Liu, B. P. Ladewig, J. C. Diniz da Costa, Cobalt-doped silica
membranes for gas separation, J. Membr. Sci. 326 (2009) 316-321

24.

Y. Gu, P. Hacarlioglu, S. T. Oyama, Hydrothermally stable silica-alumina
composite membranes for hydrogen separation, J. Membr. Sci. 310 (2008) 2837

25.

V. Tajer-Kajinebaf, H. Sarpoolaky, T. Mohammadi, Sol-gel synthesis ofnanostructured titania-silica mesoporous membranes with photo-degradationand physical separation capacities for water purification, Ceram. Int., 40(2014) 1747-1757

26.

Y. Gu, S T. Oyama, Permeation properties and hydrothermal stability ofsilica-titania membranes supported on porous alumina substrates, J. Membr.
Sci., 345 (2009) 267-275

27.

M. Asaeda, Y. Sakou, J. Yang, K. Shimasaki, Stability and performance ofporous silica-zirconia composite membranes for pervaporation of aqueousorganic solutions, J. Membr. Sci., 209 (2002) 163-175.

28.

W. Puthai, M. Kanezashi, H. Nagasawa, K. Wakamura, H. Ohnishi, T. Tsuru,
Effect of firing temperature on water permeability of SiO2-ZrO2 membranes
for nanofiltration, J. Membr. Sci., 497 (2016) 348-356.

29.

P. Pandey, R. S. Chauhan, Membranes for gas separation, Prog. Polym. Sci.
26 (2001) 853-893

30.

L. Li, H. Qi, Gas separation using sol-gel derived microporous zirconia
membranes with high hydrothermal stability, Ch. J. Chem. Eng., 23 (2015)
1300-1306.

31.

C. A. Milea, C. Bogatu, A. Duta, The influence of parameters in silica sol-gel
process, in: Bulletin of the Transilvania University of Brasov, Series I:
Engineering sciences Vol. 4 (53), 2011, pp. 59-66

32.

D. A. Ward, E. I. Ko, Preparing catalytic materials by the sol-gel method, Ind.
Eng. Chem. Res. 34 (1995) 421-433.

33.

N. Stock, S. Biswas, Synthesis of metal-organic frameworks (MOFs): routes
to various MOF topologies, morphologies, and composites, Chem. Rev. 112(2012) 933-969.

34.

S. Basu, A. Cano-Odena, I. F. J. Vankelecom, MOF-containing mixed-matrix
membranes for CO2/CH4 and CO2/N2 binary gas mixture separations, Sep.
Purif. Technol. 81 (2011) 31-40.
69

Hiroshima University Doctoral Dissertation

35.

O. G. Nik, X. Y. Chen, S. Kaliaguine, Functionalized metal organic
framework-polyimide mixed matrix membranes for CO2/CH4 separation, J.
Membr. Sci. 413-414 (2012) 48-61.

36.

D. Hoebbel, T. Reinert, H. Schmidt, E. Arpac, On the hydrolytic stability oforganic ligands in Al-, Ti- and Zr-alkoxide complexes, J. Sol-Gel Sci.
Technol. 10 (1997) 115-126.

37.

C. Janiak, Engineering coordination polymers towards applications, Dalton
Trans. (2003) 2781-2804

38.

D. Peter, T. S. Ertel, H. Bertagnolli, EXAFS study of zirconium alkoxides asprecursors in sol-gel process: II. The influence of the chemical modification, J.
Sol-Gel Sci. Technol. 5 (1995) 5-14.

39.

T. Tsuru, Nano/subnano-tuning of porous ceramic membranes for molecularseparations, J. Sol-Gel Sci. Technol. 46 (2008) 349-361

40.

M. Chen, L. Bin, S. Huai-he, Z. Lin-jie, Preparation of carbon encapsulated
metal magnetic nanoparticles by an instant pyrolysis method, New Carbon
Mater. 25 (3) (2010) 199-204

41.

Y. Li, S. Chen, H. Hu, C. Zou, Z. Chen, X. Ma, A meltable precursor forzirconium carbide ceramics and C/C-ZrC composites, Ceram. Int. 44 (2018)
10175-10180

42.

R. W. Silverstein, G. C. Bassler, Spectrometric identification of organic
compounds, J. Chem. Ed. 39 (1962) 546-553

43.

A. L. Smith, Infrared spectra-structure correlations for organosiliconcompounds, Spectrochimica Acta. 16 (1960) 87-105

44.

J. Coates, Interpretation of infrared spectra, a practical approach,
Encyclopedia of Analytical Chemistry, doi: 10. 1002/9780470027318. A5606

45.

Z.-G. Wu, Y.-X. Zhao, D.-S. Liu, The synthesis and characterization ofmesoporous silica-zirconia aerogels, Micr. Meso. Mater. 68 (2004) 127-132.

46.

R. Lee, M. Kanezashi, Y. Shimomura, T. Yoshioka, T. Tsuru, Evaluation andfabrication of pore-size-tuned silica membranes with tetraethoxydimethyl
disiloxane for gas separation, AIChE J. 57 (2011) 2755-2765.

47.

J. Yang, D. Gong, G. Li, G. Zeng, Q. Wang, Y. Zhang, G. Liu, P. Wu, E.
Vovk, Z. Peng, X. Zhou, Y. Yang, Z. Liu, Y. Sun, Self-assembly of thioureacrosslinked graphene oxide framework membranes toward separation of small
molecules, Adv. Mater. 30 (2018) 1705775 (1-8)
70Hiroshima University Doctoral Dissertation

48.

T. Ikeda, Y. Akiyama, Y. Oumi, A. Kawai, F. Mizukami, The topotactic
conversion of a novel layered silicate into a new framework zeolite, Angew.
Chem. Int. Ed. 43 (2004) 4892-4896

49.

T. Uragami, K. Okazaki, H. Matsugi, T. Miyata, Structure and permeation
characteristics of an aqueous ethanol solution of organic-inorganic hybrid
membranes

composedofpoly(vinyl

alcohol)andtetraethoxysilane,

Macromolecules 35 (2002) 9156-9163
50.

T. Fukumoto, T. Yoshioka, H. Nagasawa, M. Kanezashi, T. Tsuru,
Development and gas permeation properties of microporous amorphous TiO2ZrO2-organic composite membranes using chelating ligands, Journal of
Membrane Science. 461 (2014) 96-105.

51.

G. I. Spijksma, C. Huiskes, N. E. Benes, H. Kruidhof, D. H. A. Blank, V. G.
Kessler, H. J. M. Bouwmeester, Microporous zirconia-titania composite
membranes derived from diethanolamine-modified precursors, Adv. Mater. 18(2006) 2165-2168

52.

J. P. G. Villaluenga, B. Seoane, J. Hradil, P. Sysel, Gas permeation
characteristics of heterogeneous ODPA-BIS P polyimide membranes atdifferent temperatures, J. Membr. Sci. 305 (2007) 160-168.

53.

K. Haraya, K. Obata, N. Itoh, Y. Shindo, T. Hakuta, Y. Yoshitome, Gas
permeation and separation by an asymmetric polyimide hollow fiber
membrane, J. Membr. Sci. 41 (1989) 23-35.

54.

S. A. Stern, S. R. Sampat, S. S. Kulkarni, Test of a “free-volume” model ofgas permeation through polymer membranes. II. Pure Ar, SF6, CF4, andC2H2F2 in polyethylene, J. Polym. Sci., Part B: Polym. Phys. 24 (1986) 21492166

55.

J. P. G. Villaluenga, B. Soane, Experimental estimation of gas-transport
properties of linear low-density polyethylene membranes by an integral
permeation method, J. Appl. Polym. Sci. 82 (2001) 3013-3021

56.

Y. Hirayama, T. Yoshinaga, K. Kusuki, K. Ninomiya, T. Sakakibara, T.
Tamari, Relation of gas permeability with structure of aromatic polyimides I,
J. Membr. Sci. 111 (1996) 169-182

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Chapter 3
Microstructure engineering of ligand-modified silica-zirconia membranes
derived from network-forming precursors
3.1 Introduction
One critical element in separation membrane technology is the creation ofadvanced materials with tunable capabilities for molecular-level separation [1, 2] via
the use of so-called evolutionary and revolutionary strategies [3]. Inorganic
membranes are believed to possess the potential for revolutionary performances overpolymeric membranes because polymeric membranes are limited by an ‘upperboundary’ [4, 5], which several examples of inorganic membranes have definitively
surpassed. Porous inorganic membranes (silica, transition metal oxides, carbon,
zeolites, etc.), however, are plagued by shortcomings such as the hydrothermal
instability of silica and silica-based membranes [6-8], as well as that of zeolites [9],and also by low levels of packing density (membrane surface area per module
volume) [10] and issues with the control of pore sizes [11, 12]. Therefore, inorganic
materials have been modified with different organic moieties to address theseshortcomings [6, 11, 13].
These organically modified inorganic materials are diphasic hybrids that can begrouped into two large categories: class I and class II inorganic-organic hybrids. A
class I hybrid is such that the organic and inorganic moieties are linked by weakinteractive forces such as hydrogen and van der Waals bonding [14, 15]. Examples ofclass I hybrids are mixed-matrix membranes, interpenetrating polymer and inorganic
networks (IPNs) [16]. This class of hybrids is used more to improve the qualities ofthe polymer moiety. In class II hybrids, on the other hand, the interaction between the
inorganic and organic parts, which includes strong bonds like covalent and ionocovalent bonds, serve to improve the inorganic moiety qualities such as improvedhydrothermal stability [6], pore size controllability [11,13], targeted functionalization[17], etc. Examples include pendant and bridged-type organically modified silicon
alkoxides, organically modified transition metal oxides, metal-organic frameworks,etc. [14, 16]. Therefore, class II hybrids present a very wide field for membrane
technology development.

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Special types of class II hybrids that are referred to as ORMOCERs (hybrid
inorganic-organic polymers) are fabricated by the sequential synthesis of inorganicand organic networks from different precursors [14-16, 18, 19]. Figure 1 shows ageneral framework for the fabrication of hybrid inorganic-organic polymers based on
the work of K.-H. Haas [16]. This framework is comprised of four structural elementsthat include a siloxane network, a transition-metal oxide network (TMO) that provides
chelation sites, an organic modifier, and cross-linking of the organic units.
Traditionally, the organic R or (R) group can be bonded to the siloxane and/or TMO
networks with two aims: as network modifiers and network formers. When R is a nonhydrolysable group, it will have a network-modifying effect. For example, a network
canbemodifiedbyfunctionalization

with

a

hydrophilicgroup(aminopropyltrimethoxysilane), a hydrophobic group (phenyltriethoxysilane), or withacombinationofhydrophobicity

andoleophobicity(tridekafluoro-1,1,2,2-

tetrahydrooctyltriethoxysilane). In cases where R possesses a reactive group such as apolymerizable unit, it serves as a network former [15, 20]. For network formation,organic

precursors

such

as

methacryloxypropyltrimethoxysilane,

vinyltrimethoxysilane, etc., can be utilized [16, 20].
Irrespective of their modifier role, hybrid inorganic-organic polymers have beenused in many applications such as in protective and decorative coatings [21], barriersystems [22], dental applications [18], microelectronics [23] micro-optics [24],batteries [25], proton-exchange membrane fuel cells [26, 27], and gas separation
membranes [19]. Extensive reviews of these applications can be found in theliterature, with the exception of gas separation applications [15, 16, 18, 20]. It is
unclear why the application of hybrid inorganic-organic polymer membranes has notbeen the focus of many studies; however, two reasons can be surmised. First, because
the cross-linking of the organic modifiers is done mostly via photoinitiation of aradical polymerization initiator, the reaction can be very rapid, and conversion can becompleted in a matter of seconds or a few minutes [28, 29] making it difficult tocontrol the membrane microstructure. In another vein, a reported hybrid inorganicorganic polymer membrane was fabricated as a free-standing membrane with athickness of 100 µm. 19 Fabrication of free-standing membranes usually results inbrittle films that are difficult to handle. Thus, the choice of the radical polymerization73

Hiroshima University Doctoral Dissertation

initiator type, the initiation process, and the organic modifiers can dramatically
influence the positive fortunes of hybrid inorganic-organic polymer membranes.
In this study, the choice of the organic moieties as key structural elements for the
fabricationofmaterials

isillustrated

in

Fig.

3-1.

As

shown,

3-methacryloxypropyltriethoxysilane (MAPTMS) served a dual purpose, because it was
the precursor for the siloxane network, and it contained a siloxane network modifier(methacrylate group). Two chelating ligands, allyl acetoacetate (AAA) andacetylacetone (ACAC), were instrumental in modifying the zirconium oxide network(zirconium (IV) butoxide as a precursor). For organic cross-linking, 2,2’azobisisobutyronitrile (AIBN) was chosen as a thermal radical initiator, because it is
known to thermally dissociate at temperatures between 60 and 80 oC, which makes itsuitable for AAA co-polymerization reactions [30, 31].
Figure 3-2 shows the network formation schemes with and without organic crosslinking. For a network without organic cross-linking, the network structure iscomprised of only three of the four structural elements of a hybrid inorganic-organic
polymer. A network with organic cross-linking, however, possesses all the requisite
structural elements. However, the nature of this organic cross-linking component is
expected to differ based on the chelating ligand. In an ACAC-chelated network, since
there are no reactive units for cross-linking, homopolymerization of the MAPTMS
moiety results. In the case of an AAA-chelated network, polymerizable C=C units arepresent, and therefore a co-polymerization reaction can be expected with the
MAPTMS methacrylate unit.
The evolution of these organic cross-linking networks and the effects on the
microstructure were evaluated via various characterization methods. Furthermore, αAl2O3-supported membranes were fabricated, and their gas permeation characteristics
were evaluated. The performances of these membranes were compared vis-à-vis theirmicrostructures against different defect-free polymers and organosilica membranes.
Therefore, in this work we proposed a facile framework for the fabrication of defectfree, microstructure-controllable hybrid inorganic-organic composite membranes foruse as templates in future work.

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Hiroshima University Doctoral Dissertation
Structural elements
Components

Structure

This work(MAPTMS)

Siloxane network
TMO network(transition-metal oxide)(ZrTB)
M= Ti, Zr, Al, etc(MAPTMS)

Modified inorganic network(functionalization)(ACAC)(AAA)[MAPTMS]nCross-linked organic units(polymerized networks)[MAPTMS]m [AAA]nFig. 3-1 Structural elements in the formation of hybrid inorganic-organic polymer
materials
Hybrids with organic crosslinking

Hybrids without organic crosslinking

4

MAPTMS-SZ-ACAC
3*

3

Homo-polymerization

3

2
2

1

1

1. Siloxane network
2. TMO network
3. Saturated organic modifier
3*. Unsaturated organic modifier
4

4. Cross-linked organic units
3*

1

3*

2

1

MAPTMS-SZ-AAA

2

1

Co-polymerization

Fig. 3-2 Comparison of the structural elements of non-crosslinked and organic
crosslinked MAPTMS-SZ-LIGAND materials

3.2 Experimental

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Hiroshima University Doctoral Dissertation

In this study, samples (sol, film, powder, and membrane) derived from MAPTMSSZ-AAA and MAPTMS-SZ-ACAC systems are named as follows: non-crosslinked
samples are referred to as NC-AAA and NC-ACAC (sol, film, powder, or membrane),
respectively, and crosslinked samples are referred to as CL-AAA and CL-ACAC (sol,film, powder, or membrane), respectively.
3.2.1 Preparation of MAPTMS-SZ-LIGAND sols, powders, and films
5 weight% of NC-AAA and NC-ACAC sols in ethanol were prepared in twostages as follows. First, zirconium (iv) tert-butoxide solution in ethanol (ZrTB, purity
80% in 1-butanol, Sigma-Aldrich) was pre-modified differently with two chelating
ligands—acetylacetone (ACAC, purity 99%, Sigma-Aldrich) and allyl acetoacetate(AAA, purity >95 %, Tokyo Chemical Industries)—and stirred for one hour at room
temperature in a ZrTB/LIGAND molar proportion of 1/4. Subsequently, the 5weight% sols were realized by adding 3-methacryloxypropyltriethoxysilane(MAPTMS; 98% purity, Sigma-Aldrich; dissolved in ethanol), deionized water andhydrochloric acid (catalyst, purity 37%, Nacalai Tesque) dropwise into the ligandmodified ZrTB in a specific proportion (ZrTB/LIGAND/MAPTMS/H2O/HCl =1/4/1/8/0.25) to promote co-hydrolysis and co-condensation reactions. Hydrolysis andpoly-condensation proceeded for more than 12 hours with stirring of the mixture in aclosed vial at 500 rpm and room temperature.
To prepare the CL-AAA and CL-ACAC sols, the polymerization radical initiator,
2,2’-azobisisobutyronitrile (AIBN powder, purity 98 %, Sigma-Aldrich) was added tothe previously prepared NC-AAA and NC-ACAC sols under room temperature withstirring at 600 rpm. Four different samples were prepared according to AIBN/C=C(methacrylate) molar ratios (0, 0.01, 0.1 and 1). Where a sol with an AIBN ratio of 0
represented the control sample. The crosslinking reaction was subsequently carriedout at 80 oC in a thermostatic water bath for 24 hours. At the end of a 24-hour period,
the CL-AAA and CL-ACAC sols (AIBN/C=C 0.1 and 1) had formed thick gels thatwere insoluble in pure ethanol. P-xylene solvent was carefully added to form a clear
solution for further use.
Powders were prepared via a quick-drying method whereby the prepared sols(with and without crosslinking) were deposited dropwise into a platinum hot plate at
200 oC. Subsequently, the so formed powders were heated at 250 oC in a furnace
under a 0.5 L min-1 N2 stream for 20 min. This is to ensure that solvents are76

Hiroshima University Doctoral Dissertation

completely removed and to stabilize the material surface via further condensation
reactions. To prepare the films used to evaluate the surface structural properties, theas-prepared sols were spin-coated onto UV-pretreated Si wafers at 2,000 rpm and thenallowed to dry at room temperature prior to characterization.
3.2.2 Fabrication of MAPTMS-SZ-AAA and MAPTMS-SZ-ACAC membranes
Thin-layer membranes were formed onto prefabricated supports. Fabrication of
the supports has been described in section 2.2.1. The active separation layers(MAPTMS-SZ-AAA and MAPTMS-SZ-ACAC membranes) were formed by hotcoating the prefabricated support with a 0.5 weight% concentration of respective sols(with and without crosslinking). After the coating process, each supported membrane
was immediately calcined at 250 oC in a furnace with N2 gas flowed at a rate of 0.5 L
min-1 for 15-20 min. This was to ensure that the solvents were completely removed
and to stabilize the material surface via further condensation reactions. It is worth
noting that although the NC-AAA and NC-ACAC sols were soluble in ethanol, allsols (with and without crosslinking) were dispersed in p-xylene to a 0.5 weight% prior
to coating for consistent comparison.
3.2.3 Characterization of sols, powders, films, and membranes
The evolution of particle size distribution and particle dispersal in the preprepared sols, control sols, and organically crosslinked sols was evaluated at 25 oC via
dynamic light scattering (DLS) using a Malvern Zetasizer Nano ZS (Malvern
Instruments Ltd.). The surface and microstructural transformation of experimental
samples were evaluated as follows. The presence and transformation of structural
moieties in thin films supported on UV-treated Si-wafers were monitored using
Fourier Transform-Infrared spectroscopy (FT-IR, FTIR-4100, JASCO, Japan). The
amorphous structure nature and lattice spacing of as-prepared MAPTMS-SZLIGAND and MAPTMS-SZ-LIGAND+AIBN-derived powders were obtained via Xray diffraction spectroscopy (D2 PHASER X-Ray Diffractometer, Bruker, Germany)
using Cu Kα as the radiation source with a wavelength of 1.54 Å. Furthermore, N2
and CO2 sorption isotherms of powders were analyzed at -196 and 25 oC,
respectively, using BELMAX sorption equipment (BEL JAPAN INC., Japan). Prior
to this measurement, adsorbed gases and vapors were evacuated from the samples at
200 oC for at least 12 hours. The thermal decomposition properties of non-crosslinked
77

Hiroshima University Doctoral Dissertationand organic crosslinked powders were evaluated using thermogravimetry (DTG-60
Shimadzu Co., Japan). Finally, the supported membrane cross-sectional morphologies
were obtained using Field Emission-Scanning Microscopy (FE-SEM, Hitachi S-4800,
Japan).

Prior to examination, sections were carefully cut from the supported

membrane, attached to sample stages via special carbon tape, and vacuum dried at 50o

C for 24 hours.

3.2.4 Evaluation of gas permeation through membranes
A process-flow diagram of the gas permeation measurement setup has appeared in
Chapter 2 Fig. 2-2. Immediately after preparing the top-layer MAPTMS-SZ-AAA andMAPTMS-SZ-ACAC membrane, it was carefully fitted into its steel module inside
the furnace in the gas permeation measurement rig at 200 oC. Helium gas was thenflowed through the module at 100 ml min-1 for about 12 hours to remove adsorbed
water vapor. Subsequently, single-gas permeation tests of the membranes were carriedout using high-purity gases (H2, He, CO2, N2, CH4, CF4, SF6 in that order). Each gas
was fed to the upstream side of the closed-end module at 200 kPa absolute pressure
with temperatures ranging from 50 – 200 oC. Permeate side pressure was maintained
at atmospheric levels, and the permeate gas-flow rate was measured using a bubble
film flow meter (HORIBA-STEC, Japan).

3.3 Results and Discussion
3.3.1 Evolution of organic cross-linking units via thermal radical polymerization
3.3.1.1 Facile control of particle size distribution and heterogeneity in sols
As previously mentioned, it was surmised that the study of hybrid inorganicorganic polymer membranes has been limited because of the difficulty associated witharresting the rate of photopolymerization reactions during organic crosslinking in theconversion of C=C bonds [28, 29]. Many researchers have demonstrated theimportance of the particle size distribution of the coating sols in controlling amembrane’s microstructure, and, hence, the molecular sieving properties [33, 34].
Fig. 3-3 (a) shows the observed polydispersity indexes (PDI) and z-average valuesof as-prepared NC-AAA and CL-AAA sols while Fig. 3-3 (b) shows the data for asprepared NC-ACAC and CL-ACAC sols respectively. The PDI is a measure of theheterogeneity of a suspension based on particle size while the z-average represents theintensity weighted mean size of the dispersed particles to give the average size of78

Hiroshima University Doctoral Dissertation

particle distribution. The particle size distributions by intensity are shown in Fig. 3-4(a) and (b). Fig. 3-3 (a) clearly shows that after the crosslinking reaction, the PDI andz-average values of the CL-AAA (AIBN/C=C 0.1) showed a very high PDI (0.99) andz-average value (207 nm) compared with that of the control NC-AAA sol. The CLACAC sol showed a similar trend with respect to the z-average. The z-average value
was increased to almost 100 nm. This indicates that the addition of the AIBN radical
initiator induced a thermal polycondensation reaction. Furthermore, the PDI value ofabout 0.99 for the CL-AAA sol indicated a high level of polydispersity and suggesteda very heterogeneous dispersion [35, 36]. For the CL-ACAC sol shown in Fig. 3-3(b), however, the PDI showed only a slight change compared with the control NCACAC sol, which indicates that the heterogeneity that resulted from organic
crosslinking was low. Therefore, the copolymerization of AAA with the MAPTMS
created a high level of heterogeneity while the homopolymerization of MAPTMS in
the MAPTMS-SZ-ACAC system resulted in little heterogeneity.
As previously established, the CL-AAA and CL-ACAC sols (AIBN/C=C 0.1 and
1) formed thick gels with intensity sol sizes of 700 and 200 nm, respectively, andthese gels were insoluble in pure ethanol (Fig. 3-4). Particles within this range of sizes
are unsuitable for membrane fabrication. The sizes were considerably reduced towithin 2-8 nm (Fig. 3-4 (c)), however, via the careful addition of p-xylene solvent,and this range of sol sizes was suitable for coating onto the SiO2-ZrO2 intermediate

200

0.6

150

0.4

100

0.2

50

0.0

0
NC-AAA(Control)

1.0
Polydispersity index (PDI) [-]

0.8

250

(a) MAPTMS-SZ-AAA

Z-Average [nm]Polydispersity index (PDI) [-]1.0

substrate.

CL-AAA(AIBN/C=C 0.1)

250(b) MAPTMS-SZ-ACAC

0.8

200

0.6

150

0.4

100

0.2

50

0.0

Z-Average [nm]

layer

0
NC-ACAC(Control)

CL-ACAC(AIBN/C=C 0.1)

Fig. 3-3 Measured polydispersity index and Z-average values of control andcrosslinked sols after a 24 h reaction at 80 oC for (a) MAPTMS-SZ-AAA and (b)MAPTMS-SZ-ACAC hybrid materials.

79

Hiroshima University Doctoral Dissertation
16

16

Intensity [%]Intensity [%]

8

8

4

4

0 -1
10

0

10

1

10

2

10

Size [nm]

3

10

4

10

0 -1
10

30
20
10

(b)

(a)

MAPTMS-SZ-AAA + AIBN0.1 in PX
MAPTMS-SZ-ACAC + AIBN0.1 in PX

40

12

12
Intensity [%]50

MAPTMS-SZ-ACAC
MAPTMS-SZ-ACAC Control
MAPTMS-SZ-ACAC + AIBNMAPTMS-SZ-AAA
MAPTMS-SZ-AAA Control
MAPTMS-SZ-AAA + AIBN

0

10

1

10

2

10

3

10

0

(c)
2

4

Size [nm]

6

8

10

Size [nm]Fig. 3-4 Sol size based on intensity measured at room temperature for pre-prepared,
control and AIBN-containing (a) MAPTMS-SZ-AAA and (b) MAPTMS-SZ-ACAC
sols
3.3.1.2 Microstructural evolution of organically cross-linked materials
The structural transformations from non-crosslinked to organically crosslinked
materials were confirmed according to the chemical and physical changes. FTIR
spectroscopy represents a dependable method of monitoring the changes in chemical
structure and was used to ascertain whether the growth in size and the evolution ofpolydispersity in CL-AAA and CL-ACAC sols were due to organic crosslinking
reactions. Fig. 3-5 (a) and (b) show the observed FTIR profiles ranging from 1,8001,250 cm-1 of different MAPTMS-SZ-AAA and MAPTMS-SZ-ACAC films,
respectively, containing different AIBN/C=C ratios. The advancement of apolymerization reaction can be confirmed in a few ways: (1) the disappearance of
C=C peaks (1,638-1,640; and, 1,325 cm-1) belonging to MAPTMS [28, 29, 37] due tothe organic crosslinking; (2) the broadening and shift of the C=O peak (1,700-1,720
cm-1) of MAPTMS to a higher wave number due to the loss of conjugation with C=C;
[28, 29, 37] and, (3) the broadening and shift of peaks representative of ligand
chelation to zirconium (AAA: 1,287 cm-1; and, ACAC: 1,280 cm-1) in the cases where
co-polymerization occurs between the MAPTMS and a ligand.

80Hiroshima University Doctoral Dissertation

MAPTMS-SZ-AAA

1
0.1
0.01
Control

1800 1750 1700 1350 1300 1250

C=C (methacrylate)
C-O-C(ACAC)
C=O(methacrylate)(b)

Absorbance [a. u.]

Absorbance [a. u.](a)

C=C (methacrylate)
C=O (methacrylate)
C-O-C (AAA)

MAPTMS-SZ-ACAC

1
0.1
0.01
Control

1800 1750 1700

1300

1250

-1

-1

Wavenumber [cm ]Wavenumber [cm ]Fig. 3-5 Observed FTIR profiles of as-prepared (a) MAPTMS-SZ-AAA and (b)MAPTMS-SZ-ACAC films, respectively, at different AIBN/C=C ratios (0, 0.01, 0.1,1), as measured at room temperature and presented in a range of from 1,800-1,250
cm-1.
In Fig. 3-5 (a), three changes are evident as the AIBN/C=C ratio increased. The
C=O (methacrylate) peak was shifted to higher wavenumbers and became broad, the
C=C (methacrylate) peaks disappeared at an AIBN/C=C ratio of 1, and the AAA
chelation peak became broad and shifted to lower wavenumbers. These changes were
evident at the AIBN/C=C ratio of 0.1, which confirmed that organic crosslinking hadtaken place. The nature of this organic crosslinking can safely be assumed to be copolymerization since a transformation in the characteristic peak of AAA chelation to
zirconium can be confirmed. On the other hand, as Fig. 3-5 (b) shows, the MAPTMSSZ-ACAC samples displayed a similar trend with respect to the changes in C=C(methacrylate) and C=O (methacrylate) peaks but differed in the characteristic peak of
ACAC chelation with zirconium as the peak showed no shift as the AIBN/C=C ratio
increased. This underscores the fact that ACAC was not involved in organic
crosslinking. With respect to ligand type, AAA was involved in a co-polymerization
crosslinking with MAPTMS because it possessed an allyl C=C unit while ACAC was
not involved in any organic crosslinking since it possesses no organic crosslinking
reactive units. Although it has been concluded that AAA is involved in copolymerization with MAPTMS, the occurrence of side homopolymerization reactions
between AAA and MAPTMS reactive moieties cannot be ruled out. However, AAAin particular has been found to exhibit a ‘degradative monomer chain transfer’ wherethe -CH2- group adjacent to the allylic C=C bond causes a steric-induced resonance81

Hiroshima University Doctoral Dissertation

stabilization thereby resulting in poor chain propagation [38]. Therefore, it isreasonable to opine that a co-polymerization between MAPTMS and AAA reactive
moieties is the dominant reaction in this case.
After establishing the evolution of chemical structures via organic crosslinking
and how this is affected by the ligand type, it would be interesting to determine theeffect on the physical structure such as the microporosity and the correlation between
them. Fig. 3-6 (a) shows the N2 adsorption isotherms measured at -196 oC for NCAAA and CL-AAA (AIBN/C=C 0.1) powders while Fig. 3-6 (b) shows the isotherms
for NC-ACAC and CL-ACAC (AIBN/C=C 0.1) powders, all calcined at 250 oC,
respectively. Table 3-1 also shows the experimentally obtained BET (BrunauerEmmett-Teller) surface areas, pore volumes and average isosteric heats of CO2
adsorption values for all samples. To calculate the isosteric heats of adsorption ofCO2, the adsorption isotherms at three temperatures (20, 25 and 30 oC) were used andCrosslinked40(b) MAPTMS-SZ-ACAC

30

3

150

-1

(a) MAPTMS-SZ-AAA

N2 adsorbed [cm (STP) g ]

200

3

-1

N2 adsorbed [cm (STP) g ]the respective isotherms and heats of adsorption plots are shown in Fig. 3-7.

100
50
0
0.0

Non-crosslinked

0.2

0.4

0.6

0.8

Crosslinked

20
10

0
1.0 0.0

Non-crosslinked

0.2

0.4

0.6

0.8

1.0

p/po [-]

p/po [-]

Fig. 3-6 (a) N2 adsorption-desorption isotherms for non-crosslinked and organic
cross-linked MAPTMS-SZ-AAA powders measured at -196 oC; (b) CO2 adsorptiondesorption isotherms for non-crosslinked and organic cross-linked MAPTMS-SZAAA samples measured at 25 oC.

82

Hiroshima University Doctoral Dissertation

Table 3-1 Experimentally obtained BET (Brunauer-Emmett-Teller) surface areas(SBET), pore volumes and average isosteric heats of CO2 adsorption (Qst, Av) values for
powders calcined at 250 oC under a nitrogen gas stream.
Non-crosslinked
Material
MAPTMSSZ-AAA
MAPTMSSZ-ACAC

Crosslinked

SBET

Pore vol.

Qst,Av (CO2) SBET

Pore vol.

Qst, Av (CO2)[m2 g-1][cm3 g-1][kJ mol-1][m2 g-1][cm3 g-1][kJ mol-1]

<1

0.1

20.8

320

73.4

27.7

<1

0.2

23.3

40

8.6

28.0

Fig. 3-6 (a) and (b) show that NC-AAA and NC-ACAC powders adsorbed almost
no N2. Correspondingly, the measured surface area and pore volume according to the
BET method were less than 1 m2 g-1 and 0.2 cm3 g-1 respectively, as shown in Table 31, indicating a non-porous structure. On the other hand, after crosslinking, both
samples showed a considerable increase in the amount of N2 adsorbed compared with
non-crosslinked samples with a corresponding increase in the BET surface area andpore volumes of 320 m2 g-1 and 73.4 cm3 g-1 for CL-AAA and 40 m2 g-1 and 8.6 cm3
g-1 for CL-ACAC powders, respectively (Table 3-1).
Moreover, at very low relative pressure ranges (P/P0: 0 – 0.001) a steep
adsorption of N2 was seen for organic crosslinked samples indicating the creation of
microporosity. This was expected since mere organically modified -Si-O-Zr- chainsor planar networks can be crosslinked into either extended planar networks or 3dimensional structures creating a greater number of pores and preventing thedensification of the -Si-O-Zr- networks, as illustrated in Fig. 3-8. This is a common
phenomenon in hypercrosslinked polymers where the surface areas of non-crosslinked
polymers increase immediately following chain crosslinking [39, 40]. The surface
area measured for the CL-AAA powder sample was much higher than that measured
for the CL-ACAC sample, however, which may be because following thehomopolymerization of MAPTMS in a CL-ACAC sample, the ACAC ligands still
occupied the intra-network spaces thereby reducing the effective surface areas (Fig. 38).
83

6

o

20 C
o
25 C
o
30 C
Organic crosslinked

5
4
3
2
1
0

Non-crosslinked

0

20

40

60

7

MAPTMS-SZ-ACAC-1

MAPTMS-SZ-AAA

3

3

CO2 adsorbed [cm (STP) g ]7-1

CO2 adsorbed [cm (STP) g ]Hiroshima University Doctoral Dissertation

(a)
80

100

6

o

20 C
o
25 C
5
o
30 C
Organic crosslinked

4
3
2
1
0

Non-crosslinked

(b)

0

20

-1

Qst, CO [kJ mol ]40
Organic crosslinked

80

100

MAPTMS-SZ-ACAC

40
Organic crosslinked

30

2

2

-1

Qst, CO [kJ mol ]50

MAPTMS-SZ-AAA

30

60

Pressure [kPa]

Pressure [kPa]
50

40

20

Non-crosslinked

10
0

1

2

3

4

5
3

Non-crosslinked

10

(c)
0

20

0

6

(d)
0

1

2

3

4
3

-1

5
-1

Amount of CO2 adsorbed [cm (STP) g ] Amount of CO2 adsorbed [cm (STP) g ]Fig. 3-7 CO2 adsorption-desorption isotherms for non-crosslinked and organic crosslinked samples in (a) MAPTMS-SZ-AAA and (b) MAPTMS-SZ-ACAC measured at
20, 25 and 30 oC. Corresponding isosteric values for heat of CO2 adsorption plots
calculated based on the Clausius-Clapeyron equation are presented in panels (c) and(d), respectively.
Furthermore, the average isosteric heats of CO2 adsorption increased from 20.8
for NC-AAA sample to 27.7 kJ mol-1 for CL-AAA sample and from 23.3 for NCACAC sample to 28.0 kJ mol-1 for CL-ACAC sample. This may be due to theincreased attraction of CO2 to Zr coordinatively unsaturated sites, which are mademore accessible to CO2 because the increased surface areas in crosslinked samples
presents less hinderance for CO2 access to Zr centers [41] Therefore, a correlation
between the chemical and physical changes following crosslinking is established suchthat new connections accessible to N2 are formed that prevent the tight ordering of Si-O-Zr- chains or networks.

84

Hiroshima University Doctoral Dissertation

MAPTMS-SZ-AAA
Zr O ZrO SiOSiOO O Zr O
Si
Si
O
O Zr
O Si

Zr O Zr

Crosslinking
Si
Si

O

O Zr

O Si

Zr O SiO Si

MAPTMS-SZ-ACAC

Zr
Zr O
O Si
Zr O ZrOO
Si Crosslinking
O
Si
Si
O
O
O O Zr O
Si
Zr O Si
Si
O
O Zr
Si
O Si
O Zr
O Si
ACAC still occupying
the intra-network space

Fig. 3-8 Illustration of the network formation for non-crosslinked and organic
crosslinked MAPTMS-SZ-LIGAND materials.
3.3.2 Thermal stabilities and decomposition mechanisms: effect of organic
crosslinking
A key characteristic of membranes for robust applications is thermal stability atelevated temperatures, which has been widely substantiated for inorganic ceramic
membranes. When inorganic materials are modified with organic moieties however,
the thermal stability of the composite material becomes greatly dependent on thethermal stability of the organic modifier [42]. Fig. 3-9 (a) and (b) show the maximum
decomposition temperatures and calculated activation energies of thermal
decompositionfor

MAPTMS-SZ-AAA

andMAPTMS-SZ-ACAC

samples(crosslinked with different AIBN/C=C ratios of 0, 0.1 and 1), obtained fromdifferential thermogravimetry (DTG) data.

85

400

30

300
20

200

10

100

0

0
0

0.1

1

500

50

MAPTMS-SZ-ACAC

-1

(b)

Ea (decomposition) [kJ mol ]

-1

40

600

o

50

MAPTMS-SZ-AAA

Peak decomposition temperature [ C]

500

(a)

Ea (decomposition) [kJ mol ]600

oPeak decomposition temperature [ C]

Hiroshima University Doctoral Dissertation

40

400

30

300
20

200

10

100

0

0
0

0.1

1

AIBN/C=C ratio [-]AIBN/C=C ratio [-]Fig. 3-9 DTG peak decomposition temperatures and calculated activation energies ofthermal decomposition for MAPTMS-SZ-AAA and MAPTMS-SZ-ACAC powders
crosslinked using different AIBN/C=C ratios.
The activation energies of thermal decomposition were calculated using the
Horowitz-Metzger equation [43, 44] as shown below.
𝑊

𝜃𝐸

(3-1)

𝑓
𝑎
𝑙𝑜𝑔. [𝑙𝑜𝑔 𝑊 −𝑊
] = 2.303𝑅𝑇
2 − 𝑙𝑜𝑔2.303
𝑓𝑠

The equation describes a linear relationship such that the activation energy of𝐸𝑎thermal decomposition, 𝐸𝑎 , can be calculated from the slope 2.303𝑅𝑇
2 when𝑠

𝑊𝑓𝑙𝑜𝑔. [𝑙𝑜𝑔 𝑊 −𝑊
] is plotted against 𝜃, where 𝑊𝑓 denotes the final weight loss after𝑓complete thermal decomposition, 𝑊is the weight loss at any temperature 𝑇, 𝑇𝑠 is the
maximum decomposition temperature as obtained from the DTG curve, 𝜃 denotes thedifference (𝑇 − 𝑇𝑠 ) and 𝑅 is the universal gas constant. As shown in Fig. 3-9 (a) and(b) the peak decomposition temperatures and activation energies of thermal
decomposition follow a similar trend. As the AIBN/C=C ratio in MAPTMS-SZ-AAA
increased from 0 to 0.1, both 𝑇𝑠 and 𝐸𝑎 increased from 270 to 400 oC and 24.3 to 34.7
kJ mol-1, respectively. However, as the AIBN/C=C ratio increased to 1, 𝐸𝑎 and 𝑇𝑠
decreased substantially to 17.4 kJ mol-1 and 216 oC, respectively. This drastic
reduction may be explained by a combined decomposition of excess AIBN and free
non-crosslinked AAA. The same trend occurred in the case of MAPTMS-SZ-ACAC
where both 𝑇𝑠 and 𝐸𝑎 increased from 320 to 393 oC and 29.6 to 40.0 kJ mol-1,
respectively, as AIBN/C=C increased from 0 to 0.1. Then, 𝑇𝑠 and 𝐸𝑎 were reduced to
385 oC and 32.3 kJ mol-1, respectively, at an AIBN/C=C ratio of 1. This shows thatthe thermal stability of organic crosslinked materials increased and then decreased86Hiroshima University Doctoral Dissertation

with respect to AIBN content. This shows that the thermal stability of organic
crosslinked materials increased and then decreased with respect to AIBN content. As
indicated by FTIR, the presence of characteristic C≡N peaks when the AIBN/C=C
ratio was increased to 1 were identified. Considering the presence of free noncrosslinked AAA (having a lower thermal stability compared to ACAC), thedecomposition of unreacted IBN radicals contributed to affect the MAPTMS-SZAAA at an AIBN/C=C ratio of 1 the most. Consequently, AIBN/C=C ratio close to 1
could be detrimental to the thermal stability of organic crosslinked materials.
Therefore, for the fabrication of thermally stable and defect-free thin-layer
membranes, an AIBN/C=C ratio of 0.1 guaranteeing no free unreacted IBN radicals
was adopted.
Furthermore, the chemical structures of the organic moieties in NC-ACAC and
CL-ACAC films for example decomposed differently when compared before andafter calcination at 300 oC under N2, as shown in Fig. 3-10. MAPTMS-SZ-ACAC waschosen because it provides distinct peaks of the organic moieties involved incrosslinking (methacrylate) and those that are not (acetylacetonate). Since the C-Ccrosslinks cannot be identified via any peaks in FTIR, the stability of methacrylate
C=O (1,716-1,725 cm-1) was used as a representation since it is directly conjugatedwith the reactive C=C (methacrylate) involved in the crosslinking reaction. On theother hand, the characteristic peak of ACAC at C=C (1,590 cm-1) represents theorganic moiety not involved in crosslinking. For a consistent comparison the peakheights (absorbance values) of the organic moieties before and after calcination werenormalized
(

with

the

thermally

𝐴𝐶=𝑂 (𝑚𝑒𝑡ℎ𝑎𝑐𝑟𝑦𝑙𝑎𝑡𝑒) 𝐴𝐶=𝐶 (𝐴𝐶𝐴𝐶)
𝐴𝑆𝑖−𝑂−𝑆𝑖, 𝐴

𝑆𝑖−𝑂−𝑆𝑖stableSi-O-Si

peaks(1,080-1,090

cm-1)) and the obtained values are presented in Table 3-2.

87

Hiroshima University Doctoral Dissertation

Absorbance [a. u.]

C=O (methacrylate) C=C (ACAC) C=O (ACAC)
MAPTMS-SZ-ACAC
Non-crosslinkedo300 C
As-prepared

Crosslinkedo300 C
As-prepared

2000

1800

1600

1400

1200

-1

Wavenumber [cm ]Fig. 3-10 Observed FTIR profiles of non-crosslinked and organic cross-linked
MAPTMS-SZ-ACAC samples before and after calcination at 300 oC under a N2
atmosphere.
Table 3-2 Normalized absorbance of organic moieties in non-crosslinked andcrosslinked MAPTMS-SZ-ACAC films obtained from the FTIR spectra of Fig. 3-10.
Non-crosslinked
Condition

𝐴𝐶=𝑂 (𝑚𝑒𝑡ℎ.)
𝐴𝑆𝑖−𝑂−𝑆𝑖

Crosslinked

𝐴𝐶=𝐶 (𝐴𝐶𝐴𝐶)
𝐴𝑆𝑖−𝑂−𝑆𝑖

𝐴𝐶=𝑂 (𝑚𝑒𝑡ℎ.)
𝐴𝑆𝑖−𝑂−𝑆𝑖

𝐴𝐶=𝐶 (𝐴𝐶𝐴𝐶)
𝐴𝑆𝑖−𝑂−𝑆𝑖

[-]

[-]

[-]

[-]

As prepared

1.1

1.9

0.8

1.7

Calcined (300 oC)

0.4

0.9

0.7

1.1

In the case of NC-ACAC before and after calcination at 300

oC, C=O(methacrylate) and C=C (ACAC) peaks both had reductions in intensity by 63 and52%, respectively. However, for CL-ACAC after calcination at 300 oC, only the C=C(ACAC) peak was reduced by 39%. The C=O (methacrylate) peak, however, wasonly slightly decreased by 10% indicating that the C=O (methacrylate) peak was morestable at temperatures as high as 300 oC and by extension the organic crosslinked
network to which it was conjugated. This indicates that the ACAC moieties in CLACAC samples that were uninvolved in organic crosslinking still decomposed within
the same temperature range as that in the NC-ACAC sample. The MAPTMS moieties
88

Hiroshima University Doctoral Dissertationin CL-ACAC and NC-ACAC, however, did not decompose in the same manner.
Decomposition was delayed to a higher temperature in CL-ACAC, at which point thedegradation was very rapid. Therefore, in organically crosslinked materials, thecrosslinked moieties resulted in better stability.
Interestingly, the primary thermogravimetry (TG) profiles reveal different
observations. The final residual weights after complete thermal decomposition
become lower as the AIBN/C=C ratio increases, as shown in Fig. 3-11. It should be
noted that the TG profile of a control sample is overlapped by that of a crosslinked
sample with AIBN/C=C 0.01. This is a trend that is associated with a decrease inthermal stability. This disagrees with the previously established fact that organicallycrosslinked samples are more thermally stable. This can be explained by the additionof a crosslinking agent. AIBN radical initiators generate IBN radicals for the
crosslinking process thus creating more crosslinking initiation sites as more AIBN is
added. Therefore, as more AIBN radical initiators are added, more organically richmaterial results, which leads to more weight loss at the final decomposition stage.

Residual weight [-]1.0

MAPTMS-SZ-ACAC

0.8
0.6
0.4

AIBN/C=C
0
0.01
0.1
1

0.2
0.0
100

300

500

700

900

o

Temperature [ C]
Fig. 3-11 Thermogravimetry (TGA) curves of MAPTMS-SZ-ACAC samples withdifferent AIBN/C=C ratios fired under a N2 atmosphere.
3.3.3 Membranes fabricated from cross-linked hybrids and gas permeation
analysis
3.3.3.1 Morphologies of supported membranes

89Hiroshima University Doctoral Dissertation

Fig. 3-12 (a) and (b) show the cross-sectional morphology of NC-AAA and CLAAA membranes taken by FE-SEM on a scale of 100 nm (x300, 000 magnification).
These images present a clear distinction between the active separation and SiO2-ZrO2
substrate layers. The SiO2-ZrO2 intermediate layer and active top separation layer
were adequately controlled to within 400 nm. This is instructive since controlling the
fabrication of membranes that are as thin as possible reduces permeation resistance
while simultaneously maintaining a high level of molecular sieving properties. By
contrast, a crack-free, free-standing, inorganic-organic hybrid membrane with athickness of 100 µm had a very low H2 permeance of 1.14 x 10-13 mol m-2 s-1 Pa-1
[19].
As Fig. 3-12 (a) clearly demonstrates, the CL-AAA membrane formed a distinct
layer about 20 nm at its thickest point over the SiO2-ZrO2 substrate. On the other
hand, the NC-AAA membrane formed a much thinner, almost indistinct, layer lessthan 15 nm at its thickest point on the SiO2-ZrO2 substrate. This can be explained by
the fact that the CL-AAA has a much higher polydispersity (Fig. 3-3 (a)) with anextensively crosslinked structure that is deposited much thicker and distinctly over a
porous surface. When the NC-AAA is coated over a porous surface, its much lower
polydispersity (Fig. 3-3 (a)) and non-crosslinked structure allows the particles to settle
into the underlying surface thereby forming a thinner layer that is almost indistinct.
To support this explanation, Fig. 3-12 (c) and (d) compare the deposition behaviors of
CL-AAA and NC-AAA on a copper grid taken via transmission electron micrography(TEM). As shown in Fig. 3-12 (c) and (d), after deposition the CL-AAA with its
extended structure was deposited across the grid supports and spaces while the NCAAA was deposited only on the grid supports and penetrated the spaces.

90Hiroshima University Doctoral Dissertation

Fig. 3-12 Cross-sectional morphologies of (a) CL-AAA- and (b) NC-AAA-derived
membranes obtained via field emission-scanning electron microscopy (FE-SEM300,000x magnification) and the transmission electron microscopy (TEM- 30,000x
magnification) images of deposited (c) CL-AAA- and (b) NC-AAA-derived films.
3.3.3.2 Gas sieving properties at 200 oC and membrane microstructures
Four different types of membranes were fabricated: non-crosslinked and organic
crosslinked MAPTMS-SZ-AAA-derived membranes (NC-AAA and CL-AAA); noncrosslinked and organic crosslinked MAPTMS-SZ-ACAC-derived membranes (NCACAC and CL-ACAC). Fig. 3-13 (a) shows single-gas permeance as a function of
kinetic diameter (He: 0.26 nm, H2; 0.289 nm, CO2: 0.33 nm, N2: 0.364 nm, CH4: 0.38
nm, CF4: 0.48 nm, SF6: 0.55 nm) at 200 oC. From Fig. 3-13 (a) it is evident that allmembranes displayed molecular sieving properties such that gases with smaller
kinetic diameters generally showed a higher permeance than gases with much larger
kinetic diameters. Another observation is that non-crosslinked membranes generally
showed higher permeance (H2 permeance; NC-AAA: 5.4 x 10-7 mol m-2 s-1 Pa-1, NCACAC: 8.7 x 10-7 mol m-2 s-1 Pa-1) than the organic crosslinked versions (H2permeance; CL-AAA: 2.4 x 10-7 mol m-2 s-1 Pa-1, CL-ACAC: 1.9 x 10-7 mol m-2 s-1
Pa-1). This is consistent with the morphology of each membrane type as seen in Fig.
91

Hiroshima University Doctoral Dissertation

3-12 (a) and (b). The NC-AAA formed a thinner layer on the SiO2-ZrO2 substrate,and, thus, was expected to present a lower resistance to gas permeation. On the other
hand, the CL-AAA formed a thicker layer and was expected to present higher
resistance to gas permeation.

10

-7

10

-8

He H2CO2CH4
N210

CF4 SF6

MAPTMS-SZ-AAA
MAPTMS-SZ-ACAC

-9

10

-10

10

-11

Open: Crosslinked
Closed: Non-crosslinked

0.2 0.3 0.4 0.5 0.6 0.7 0.8(b)Dimensionless permeance [-]

10

-6

-2

-1

-1

Permeance [mol m s Pa ](a)

He H2CO2CH4

10

0

10

-1

10

-2

10

-3

10

-4

CF4 SF6

N2

Knudsen
MAPTMS-SZ-AAA
MAPTMS-SZ-ACAC

Open: Crosslinked
Closed: Non-crosslinked

0.2 0.3 0.4 0.5 0.6 0.7 0.8

Kinetic diameter [nm]Kinetic diameter [nm]Fig. 3-13 Kinetic diameter dependence of (a) observed single-gas permeance and (b)dimensionless permeance (normalized on helium basis) of non-crosslinked andorganic

crosslinked

MAPTMS-SZ-AAA-

andMAPTMS-SZ-ACAC-derived

membranes at 200 oC.
Table 3-3 H2 permeance and H2/X permeance ratio values for each membrane type at
200 oC

Membrane typeH2 permeance

H2/X permeance ratio[10-7 mol m-2 s-1 Pa-1][-]
He

CO2 N2

CH4

CF4

SF6

NC-AAA

5.4

1.3

2.1

10.9

7.2

188.3

683.7

CL-AAA

2.4

1.0

4.9

32.9

29.0

2117.0-NC-ACAC

8.7

1.5

1.9

7.4

4.3

114.0

407.4

CL-ACAC

1.9

1.0

3.7

19.8

14.9

230.6

319.4

To get an idea of the pore size distribution and membrane microstructures afterorganic crosslinking, the plot of dimensionless permeance (normalized based on He
92

Hiroshima University Doctoral Dissertation

permeance) as a function of kinetic diameter of the different gases for different
membranes at 200 oC is instructive, as shown in Fig. 3-13 (b). The plot reveals that
NC-AAA- and NC-ACAC-derived membranes possessed similar dimensionless gas
permeance values indicating similar pore size distribution. In contrast, CL-AAA- and
CL-ACAC-derived membranes showed lower dimensionless gas permeance values
indicating narrower pore size distribution. This suggests that the crosslinking in CLAAA and CL-ACAC membranes modifies the microstructure enough to improve themolecular sieving properties (H2/N2 selectivity; CL-AAA: 33, CL-ACAC: 20)compared with non-crosslink-derived membranes (H2/N2 selectivity; NC-AAA: 11,
NC-ACAC: 7). Table 3-3 presents the H2/X permeance ratios for all gases. It is also
worth noting that the molecular sieving properties were better for CL-AAA comparedwith that of CL-ACAC in that while the H2 permeance was comparable H2/N2
selectivity was higher. This observation is instructive because according to the workof Schmidt [45] network formation by polymerization of modifying ligands can becomplex, and some polymerization leads to a mere organic polymer chain that is
almost independent of the main inorganic network while others will form truly
interpenetrating chains with an inorganic superstructure, and both possibilities occur
in the same material.
The modified gas translation model proposed by Lee et al. [46] was used topropose a way to estimate the mean pore size of <1 nm for membranes by employing
a NKP (normalized Knudsen permeance) method based on the modified gas
translation model expressed in Equation 3-2. Fig. 3-14 (a)-(d) show the calculated
NKP values plotted as a function of kinetic diameter fitted to Equation 3-3 for NCAAA, CL-AAA, NC-ACAC, and CL-ACAC-derived membranes, respectively. The
figures reveal a rather interesting observation. Fig. 3-14 (a) and (c) show that two
different mean pore diameters were estimated for the NC-AAA- (0.56 and 0.62 nm)
and NC-ACAC-derived (0.65 and 0.70 nm) membranes, respectively, at 50 and 200
oC. In contrast, the estimated mean pore diameters for both CL-AAA- and CL-ACAC-

derived membranes at 50 and 200 oC were very close (Fig. 3-14 (b) and (d)). This
indicates that non-crosslink-derived membranes express different pore sizes
depending on the permeation temperature. This may be explained by the fact that the
non-crosslinked organic moieties in NC-AAA- and NC-ACAC-derived membranes
are flexible and thus vibrate with an increase in temperature similar to polymer
93

Hiroshima University Doctoral Dissertation

matrices. Therefore, at low temperatures the flexible organic moieties are stiff andoccupy much of the intraparticle free volume. High selectivity (H2/SF6 ⁓1000) can beobtained, but at high temperatures their continuous vibrations create more free volume
for gas permeation. Therefore, it is difficult to estimate a single mean pore diameter incases where organic moieties are not crosslinked and are free to vibrate.
1 𝜀

𝑃𝑖 = 3 𝜏𝐿 (𝑑𝑝 − 𝑑𝑖 )

𝑀

𝑠

𝑆

1.0

𝐸𝑝,𝑖

8
𝑖

𝑑 3
𝑑𝑝
3
𝑑
(1− 𝑠 )
𝑑𝑝

He

(3-3)

1.0
CL-AAA

H2

0.56 nm

0.4

CL-AAA

H2

o

50 C
o
200 C

0.6
0.4

0.62 nm

(a)

He

0.8

o

50 C
o
200 C

0.6

0.2

(3-2)

√𝜋𝑀 𝑅𝑇 𝑒𝑥𝑝 (− 𝑅𝑇 )(1− 𝑖 )

0.8

NKP [-]

2𝑑𝑝

2

NKP [-]𝑃

𝑓𝑁𝐾𝑃 = 𝑃𝑖 √𝑀 𝑖 ≈(𝑑𝑝 −𝑑𝑖 )

0.2

SF6

0.43 nm0.45 nm

(b)

CF4

0.0
0.0
0.2 0.3 0.4 0.5 0.6 0.7 0.2 0.3 0.4 0.5 0.6 0.7
1.0 Kinetic diameter [nm] 1.0 Kinetic diameter [nm]NC-ACAC

0.6

0.65 nm0.4
0.2

0.70 nm(c)

CL-ACAC

0.8

o

50 C
o
200 C

NKP [-]NKP [-]

0.8

o

50 C
o
200 C

0.6

0.57 nm

0.4
0.2

(d)

0.56 nm0.0
0.0
0.2 0.3 0.4 0.5 0.6 0.7 0.2 0.3 0.4 0.5 0.6 0.7
Kinetic diameter [nm]Kinetic diameter [nm]Fig. 3-14 NKP as a function of kinetic diameter fitted to Equation 4 for noncrosslinked and organic crosslinked MAPTMS-SZ-AAA and MAPTMS-SZ-ACAC.
The solid and dashed lines are theoretical curves.

94

Hiroshima University Doctoral Dissertation

An alternative approach for estimating the mean pore diameter of membranes is
the k0 plot [46]. The k0 plot presents a temperature-independent method for
calculating the pore size by employing the apparent activation energies of gas
permeation, Ep, and pre-exponential factors, k0, of a gaseous species, as expressed in
Equation 3-4.
⁄

1 3
𝑘0,𝑖
= 𝑎1⁄3 𝑑𝑝 - 𝑎1⁄3 𝑑𝑖(3-4)

Therefore, the estimated pore diameters are the pore sizes for an infinite range oftemperatures. Fig. 3-15 (a) and (b) show the k0 plots as a relationship between the1⁄

𝑘0,𝑖3 and the kinetic diameter for NC-AAA, CL-AAA, NC-ACAC and CL-ACACderived membranes, respectively. It is interesting to note that the estimated mean pore
diameters of CL-AAA and CL-ACAC-derived membranes according to k0 (CL-AAA:
0.48 nm; CL-ACAC: 0.55 nm) were very similar to the NKP-estimated pore
diameters at both 50 and 200 oC. On the other hand, the k0-estimated mean pore
diameters for NC-AAA and NC-ACAC-derived membranes (NC-AAA: 0.76 nm; NCACAC: 0.84 nm) were considerably higher than those estimated via NKP. Theseobservations establish that the NKP method of estimating the mean pore diameters of
NC-AAA and NC-ACAC-derived membranes more accurately defines the highselectivity of small gases over that of large gases, which cannot be expected frompore diameters that are estimated using the k0 plots.

(a)

(b)

0.12

0.12

MAPTMS-SZ-AAA

He

CH4

0.04

Non cross-linked

CF4

Organic cross-linked
N2
CH4

1/3

[-]

0.08

CO2

k0, i

k0, i

Organic cross-linked

CO2

MAPTMS-SZ-ACAC

H2

N2

0.08

1/3

[-]

H2

He

CF4

0.04

Non cross-linked

SF6

0.00
0.2

0.48 nm

0.4

0.76 nm

0.6

0.8

1.0

Kinetic diameter [nm]

0.00
0.2

0.55 nm

0.4

SF6

0.6

0.84 nm0.8

Kinetic diameter [nm]

95

1.0

Hiroshima University Doctoral Dissertation
1⁄

Fig. 3-15 Relationship between 𝑘0,𝑖3 and kinetic diameter for (a) non-crosslinked andorganic crosslinked MAPTMS-SZ-AAA and (b) non-crosslinked and organic
crosslinked MAPTMS-SZ-ACAC.
3.3.3.3 Analysis of microporous structures via the gas permeance ratio and activationenergyThe microstructures of membranes can be evaluated by examining the relationshipbetween the permeation properties of gas molecules. Specifically, the relationshipbetween He and H2 permeation behaviors give insight into the sub-ultra-microporous
nature of membranes [47]. Fig. 3-17 (a) and (b) plot the activation energy of He
permeance and the He/H2 permeance ratio as functions of the activation energy of H2
permeance, and the activation energy of N2 permeance along with the H2/N2
permeance ratio as functions of the activation energy of H2 permeance, respectively.
The activation energies were calculated by fitting the plots of the temperaturedependence of gas permeance (shown in Fig. 3-16 (a)-(d) for NC-AAA, CL-AAA,
NC-ACAC and CL-ACAC-derived membranes, respectively) to Equation 2-1, using
activation energy, Ep, and k0 as fitting parameters. Fig. 3-17 (a) and (b) show that theresults of this work, which were reproduced, compare with various other state-of-theart membranes such as glassy polymer membranes [48-51] rubbery polymer
membranes [51] and silica-based membranes [52-54]. All related data are presented inTable 3-4. These various state-of-the-art membranes were carefully chosen to
represent various membrane types and microstructures.

96

Hiroshima University Doctoral Dissertation
MAPTMS-SZ-AAA+AIBN

MAPTMS-SZ-AAA+AIBN0.1ooTemperature [ C]
200 150

100

50

(a)

10

-1

H2

CO2

-1

-7

10

He
CH4

-8

10

N2

CF4
-9

10

SF6

-10

10

100

-7

10

-8

10

(b)

He
H2
CO2
CH4
N2

-9

10

-10

10

50

CF4

He
-1

-7

10
10

CF4

-9

10

SF6

(c)

2.0

10

-7

10

-2

-2

N2

-8

200 150

2.8

50

H2
He
CO2

-8

N2

-9

SF6

10
10

-10

2.4

100

-6

-1

Permeance [mol m s Pa ]

-1

CO2

CH4

-10

50

MAPTMS-SZ-ACAC 10MAPTMS-SZ-ACAC+AIBN0.1
2.4
2.8
3.2
2.0
2.4
2.8
3.2
o
o
-1[ C]-1
Temperature
Temperature
1000/T [K ]1000/T [K [] C]

200 150
H2
10

10

100

-11

2.0

-6

-1

200 150-6-2-2Permeance [mol m s Pa ]

10

-1

-1

Permeance [mol m s Pa ]-6

Permeance [mol m s Pa ]Temperature [ C]

CH4

(d)

10
3.2 2.0

2.4

2.8

3.2

-1

-1

1000/T [K ]1000/T [K ]Fig. 3-16 Plots of the temperature dependence of single-gas permeance formembranes derived from (a) non-crosslinked MAPTMS-SZ-AAA, (b) organic
crosslinked MAPTMS-SZ-AAA, (c) non-crosslinked MAPTMS-SZ-ACAC, and (d)organic crosslinked MAPTMS-SZ-ACAC.
Fig. 3-17 (a) shows that all membranes displayed very similar He-H2 correlations
without substantial deviations, which suggests that the activation energies ofpermeance, and by extension the permeation paths, for He and H2 are very similarirrespective of membrane type. Fig. 3-17 (a) also shows that the He/H2 permeanceratio generally increased as the activation energy of H2 permeance increased for all
membranes with polymer membranes showing the highest values. Therefore, in thisrespect, non-crosslinked and organic crosslink-derived membranes follow a similar
trend.
97

Hiroshima University Doctoral Dissertation

A deviation in the characteristics of the membranes began to appear, however, as
shown in Fig. 3-17 (b) as H2 and N2 began to permeate via different paths. As shownby the dashed lines, the non-crosslink-derived membranes showed a different trend
from the organic-crosslink-derived membranes. The trends for NC-AAA and NCACAC tended towards that of polymer membranes that show a very high Ep for N2
with respect to the Ep for H2. This provided direct evidence of the similarities in the
microstructural behavior of non-crosslinked organic moieties and polymer matrices.
On the other hand, however, the trends for CL-AAA and CL-ACAC followed that ofsilica-based membranes, as evident from Fig. 3-17 (b). In this case, both values for Ep(N2 and H2) tended to be similar. Fig. 3-17 (b) also shows that, irrespective of themembrane type, the H2/N2 permeance ratio increased as the Ep (H2) of a membrane
became higher. It is worth noting that CL-AAA-derived membranes displayed better
performance than CL-ACAC-derived membranes because of a greater amount ofextended crosslinking, which combined with the parent -Si-O-Zr- network to form

1.5

 (H2/N2) [-] (He/H2) [-]better molecular sieving networks.
(a)

1.0
0.5

10

2

10

1

(b)
This work
CL-AAA
CL-ACAC
NC-AAA
NC-ACAC0cross-linked
10AAA-Organic
ACAC-Organic cross-linked
35AAA-Non cross-linked
ACAC-Non cross-linked-1

Ep (N2) [kJ mol ]

20
15
10
5
0

3

30

25

-1

Ep (He) [kJ mol ]30

10Polymers
25 Polyimide

PDMS

15 BTESE5

10 15 20 25-1
Ep (H2) [kJ mol ]30This work

BTESO

10 POSS-PI-Ph
50

Polyimide
Cellulose acetate
Polysulfone
PDMS
BTESE
BTESO
POSS-PI-Ph
CVD-Silica

Cellulose acetate

20 Polysulfone

0

CVD-silica

05Silica-based

10

15

20

25

30

-1

Ep (H2) [kJ mol ]Fig. 3-17 (a) Activation energy of He permeance and the He/H2 permeance ratio at
200 oC as functions of the activation energy of H2 permeance; and, (b) Activationenergy of N2 permeance and H2/N2 permeance ratios at 200 oC as functions of theactivation energy of H2 permeance. Dashed lines are for a visual aid.
Table 3-4 Apparent activation energies of H2 and N2 permeation for different
reference membranes

98Hiroshima University Doctoral Dissertation

Activation energy [kJ mol-1]

Permeance ratios [-]

He

H2

N2

He/H2

H2/N2

Polyimide

17.0

17.6

25.9

1.2

162.7

[48]

Polyimide

-

14.1

23.1

-

-

[49]

Polyimide

-

13.7

21.7

-

-

[49]

Cellulose acetate

14.4

14.8

21.5

1.3

46.7

[50]

Polysulfone

-

13.2

23.2

-

50.0

[51]

PDMS

-

6.1

9.5

-

3.3

[51]

BTESE

1.8

1.0

0.6

-

16.2

[52]

BTESO

6.3

5.7

6.6

-

5.3

[52]

POSS-PI-Ph

19.5

21.5

31.5

1.0

33.3

[53]

CVD-silica

10.4

10

9.6

1.0

53.0

[54]

Membrane

Reference

Fig. 3-18 (a) and (b) show the plots of H2 and N2 permeance and H2/N2 selectivity
as functions of temperature for CL-AAA and CL-ACAC-derived membranes and NCAAA and NCL-ACAC-derived membranes, respectively. The H2-N2 pair was chosen
since they possess very different kinetic diameters and thus travel through different
permeation paths. As expected, the H2/N2 selectivity for all membranes was slightly
reduced with temperature while the H2/N2 selectivity showed greater temperaturedependence NC-AAA and NC-ACAC-derived membranes. Close examination of theplots also reveals that for both CL-AAA- and CL-ACAC-derived membranes H2 andN2 permeance showed a higher dependence on temperature with corresponding higher
apparent activation energies compared with the non-crosslinked alternatives, as shownin Fig. 3-18. It is generally agreed that higher activation energy is required topermeate smaller pores [47] and the higher level of activation energy values for CLAAA- and CL-ACAC-derived membranes support the estimations for mean pore
diameters. Associating a smaller pore size accompanied by higher apparent activationenergies of gas permeation with the membrane microstructural characteristics may be
misleading, however, when comparing membranes—particularly those of different
classes.
99

Hiroshima University Doctoral Dissertation

1

10

Ep = 10.5 kJ/ mol

H2 Ep = 15.2 kJ/ mol

10

-2

-2

-1

-7

MAPTMS-SZ-AAA

-8

10

-9

10

Ep = 14.0 kJ/ mol

N2

(a)
10
50
-10

Ep = 17.5 kJ/ mol

Red: Crosslinked
Black: Non-crosslinked

100

150
o

Temperature [ C]200

MAPTMS-SZ-ACAC

1

10

0

10-6
10
H2

Ep = 8.1 kJ/ mol

-1

-1
-1

0

10-6
10

Permeance [mol m s Pa ]2

10Permeance [mol m s Pa ] H2/N2 [-]H2/N2 [-]

2

10

-7

10

Ep = 15.1 kJ/ mol
Ep = 11.0 kJ/ mol-8N2

Ep = 16.0 kJ/ mol

10

Red: Crosslinked
Black: Non-crosslinked

(b)
10
50
-9

100

150

200

o

Temperature [ C]

Fig. 3-18 H2 and N2 permeance and H2/N2 selectivity as functions of temperature for(a) non-crosslinked and organic crosslinked MAPTMS-SZ-AAA; and (b) noncrosslinked and organic crosslinked MAPTMS-SZ-ACAC.
Resolving this problem with comparisons required defining the intrinsic
microstructural characteristics based on the normalized values of the apparent
activation energies of the chosen gas pairs. An intrinsic characteristic is the ‘intrinsic
rigidity’ (or flexibility). For membranes with very flexible microstructures such as
rubbery polymer membranes, the activation energy of gas permeance increases withkinetic diameter. In rigid microstructures such as glassy polymer membranes, the gas
kinetic diameter has less influence on the activation energies [3]. Based on this, theintrinsic rigidity of a membrane can be defined as the normalized difference in theactivation energies of the permeance of a large gas and that of a smaller gas through amembrane—in this case, 1 – [Ep (H2) / Ep (N2)]. The relative rigidity is determined by
the value of this expression.
Fig. 3-19 plots the defined ‘intrinsic rigidity’ as a function of Ep (N2) / Ep (H2) for
various membranes. Details of the activation energy values are presented in Table 3-4.
A correlation was established such that increasingly ‘flexible’ membranes show
increasingly polymer-like characteristics. On the other hand, increasingly ‘rigid’
membranes show increasingly silica-like microstructures. Organosilica membranes
are regarded as rigid, porous membranes with some degree of flexibility depending on
the bridging organic group, and the greater the length of the organic C-C bridge100Hiroshima University Doctoral Dissertation

between Si atoms the more flexibility it will demonstrate [52, 55]. In Fig. 3-19, a
BTESO-derived membrane with an octyl hydrocarbon bridge tended towards a more
flexible and dense structure compared with a BTESE-derived membrane with an ethylhydrocarbon bridge. The NC-AAA- and NC-ACAC-derived membranes showed atrend closer to polymer membranes confirming the theory of flexible non-crosslinked
moieties. The CL-AAA- and CL-ACAC-derived membranes on the other hand tended
towards more ‘rigid’ structures.

Material spectrum
Silica-rich Hybrid

Polymer
Flexible

0.4
0.3
0.2
0.1
0.0
-0.1
-0.2
0.8

Rigid

1 - [Ep(H2)/Ep(N2)] [-]

0.5

1.0

1.2

1.4

1.6

CL-AAA
CL-ACAC
NC-AAA
NC-ACAC
Polyimide
Cellulose acetate
Polysulfone
PDMS
BTESE
BTESO
POSS-PI-Ph
CVD-silica

1.8

Ep(N2)/Ep(H2) [-]Fig. 3-19 The correlation between ‘intrinsic rigidity’ 1 – [Ep (H2) / Ep (N2)] andmaterial class calculated based on Ep (N2) / Ep (H2) values for different membranes.
Absolute values are shown in Fig. 3-18 and Table 3-4.

3.4

Conclusions
In this work, SiO2-ZrO2-based organically crosslinked hybrid membranes were

fabricated from class II hybrid inorganic-organic polymer materials prepared via thesol-gel crosslinking of ligand (AAA or ACAC)-chelated zirconium oxide networks
and MAPTMS derived siloxane networks. The membrane microstructuralcharacteristics were evaluated by analyzing the results of single-gas permeation
experiments for temperatures ranging from 50 to 200 oC. The following conclusions
were drawn.

101

Hiroshima University Doctoral Dissertation

1. The AAA-chelated system generated a copolymerization crosslinking with the
Si precursor- 3-methacryloxypropyltriethoxysilane (MAPTMS)- while the ACACchelated system generated only a homopolymerization within the MAPTMS.
2. New pore structures were formed as organic crosslinked materials showed asignificant increase in measured BET surface areas (OCL-AAA: ≈320 m2 g-1, OCLACAC: ≈40 m2 g-1) compared with non-crosslinked samples (< 1 m2 g-1).
3. The organic crosslink-derived membranes showed better molecular sieving
properties with higher H2/N2 selectivity (OCL-AAA: 33, OCL-ACAC: 20) comparedwith non-crosslink-derived membranes (H2/N2 selectivity; NCL-AAA: 11, NCLACAC: 7).
4. Organic-crosslinked and non-crosslink-derived membranes revealed different
membrane microstructures. Two different mean pore diameters were estimated for the
NCL-AAA- (0.56 and 0.62 nm) and NCL-ACAC-derived (0.65 and 0.70 nm)
membranes, respectively, at 50 and 200 oC. On the other hand, the estimated meanpore diameters for both OCL-AAA- and OCL-ACAC-derived membranes at 50 and
200 oC were very similar.
5. Finally, characterization of the membrane microstructure based on thepermeation properties of H2 and N2 revealed that organic crosslink-derived
membranes showed a change in microstructural properties from a flexible to a rigid
structure that ranged between that of silica-rich materials and polymer materials.
Considering that a facile method was established to fabricate and tune themicrostructure of supported class II hybrid inorganic-organic highly permeable
membranes, this study should serve as a template for future work where tailored
network tuning and functionalization can be achieved for specific applications.

Cited literature
1.

W. J. Koros, Evolving beyond the thermal age of separation processes:
membranes can lead the way, AIChE J., 50 (2004) 2326-2334

2.

W. J. Koros, A more aggressive approach is needed for gas separation
membrane development, Membrane, 31 (2006) 157-160

3.

M. R. Abdul Hamid, H. -K. Jeong, Recent advances on mixed-matrix
membranes for gas separation: opportunities and engineering challenges,
Korean J. Chem. Eng. 35 (2018) 1577-1600

102

Hiroshima University Doctoral Dissertation

4.

L. M. Robeson, Polymer membranes for gas separation, Curr. Opin. Solid St.
M. 4 (1999) 549-552

5.

L. M. Robeson, The upper bound revisited, J. Membr. Sci., 320 (2008) 390400

6.

H. L. Castricum, A. Sah, R. Kreiter, D. H. A. Blank, J. F. Vente, J. E. ten
Elshof, Hydrothermally stable molecular separation membranes fromorganically linked silica, J. Mater. Chem., 18 (2008) 2150-2158

7.

K. Yoshida, Y. Hirano, H. Fujii, T. Tsuru, M. Asaeda, Hydrothermal stability
and performance of silica-zirconia for hydrogen separation in hydrothermal
conditions, J. Chem. Eng. Jpn. 34 (2001) 523-530

8.

G. P. Fotou, Y. S. Lin, S. E. Pratsinis, Hydrothermal stability of pure andmodified microporous silica membranes, J. Mater. Sci., 30 (1995) 2803-2808

9.

Y. Li, H. Zhou, G. Zhu, J. Liu, W. Yang, Hydrothermal stability of LTA
zeolite membranes in pervaporation, J. Membr. Sci., 297 (2007) 10-15

10.

T. Tsuru, Nano/subnano-tuning of porous ceramic membranes for molecularseparations, J. Sol-Gel Sci. Technol., 46 (2008) 349-361

11.

N. K. Raman, C. J. Brinker, Organic “template” approach to molecular sieving
silica membranes, J. Membr. Sci., 105 (1995) 273-279

12.

Y. Ohta, K. Akamatsu, T. Sugawara, A. Nakao, A. Miyoshi, S. -I. Nakao,
Development of pore size-controlled silica membranes for gas separation bychemical vapor deposition, J. Membr. Sci., 315 (2008) 93-99

13.

M. Kanezashi, M. Kawano, T. Yoshioka, T. Tsuru, Organic-inorganic hybrid
silica membranes with controlled silica network size for propylene/propane
separation, Ind. Eng. Chem. Res., 51 (2012) 944-953

14.

P. Judeinstein, C. Sanchez, Hybrid organic-inorganic materials: a land of
multidisciplinarity, J. Mater. Chem., 6 (1996) 511-525

15.

C. Sanchez, J. de A. A. Soler-Illia, F. Ribot, T. Lalot, C. R. Mayer, V. Cabuil,
Designed

Hybridorganic-inorganic

nanocomposites

fromfunctional

nanobuilding blocks, Chem. Mater. 13 (2001) 3061-3083
16.

K. -H. Haas, Hybrid inorganic-organic polymers based on organicallymodified Si-alkoxides, Adv. Eng. Mater., 2 (2000) 571-582

17.

L. Yu, M. Kanezashi, H. Nagasawa, J. Oshita, A. Naka, T. Tsuru, Pyrimidinebridged organoalkoxysilane membrane for high-efficiency CO2 transport via
mild affinity, Sep. Purif. Technol., 178 (2017) 232-241
103

Hiroshima University Doctoral Dissertation

18.

K. -H. Haas, H. Wolter, Synthesis, properties and applications of inorganicorganic copolymers (ORMOCERs), Curr. Opin. Solid St. M., 4 (1999) 571580

19.

S. M. Kumbar, T. Selvam, C. Gellermann, W. Storch, T. Ballweg, J. Breu, G.
Sextl, ORMOCERs (organic-inorganic hybrid copolymers)-zeolite beta (BEA)
nanocomposite membranes for gas separation applications, J. Membr. Sci.,
347 (2010) 132-140

20.

C. Sanchez, B. Julian, P. Belleville, M. Popall, Applications of hybrid organicinorganic nanocomposites, J. Mater. Chem. 15 (2005) 3559-3592

21.

G. Schottner, K. Rose, U. Posset, Scratch and abrasion resistant coatings onplastic lenses-state of the art, current developments and perspectives, J. SolGel Sci. Technol., 27 (2003) 71-79

22.

S. Amberg-Schwab, H. Katschorek, U. Weber, A. Burger, Inorganic-organic
polymers as migration barriers against liquid and volatile compounds, J. SolGel Sci. Technol., 26 (2003) 699-703

23.

M. Popall, J. Kappel, M. Pilz, J. Schulz, G. Feyder, A new inorganic-organic
polymer for the passivation of thin film capacitors, J. Sol-Gel Sci. Technol., 2(1994) 157-160

24.

R. Houbertz, G. Domann, J. Schulz, B. Olsowski, L. Frohlich, W. -S. Kim,
Impact of photoinitiators on the photopolymerization and the opticalproperties of inorganic-organic hybrid polymers, Appl. Phys. Lett., 84 (2004)
1105

25.

J. F. Velez, J. M. Aparicio, J. Mosa, Covalent silica-PEO-LiTFSI hybrid solid
electrolytes via sol-gel for Li-ion battery applications, Electrochim. Acta., 213(2016) 831-841

26.

K. Prater, The renaissance of the solid polymer fuel cell, J. Power Sources., 29(1990) 239-250

27.

B. Bonnet, D. J. Jones, J. Roziere, L. Tchicaya, G Alberti, M. Casciola, L.
Massinelli, B. Bauer, A. Peraio, E. Rammuni, Hybrid organic-inorganic
membranes for a medium temperature fuel cell, J. New Mat. Electrochem.
System., 3 (2000) 87-92

28.

O. Soppera, C. Croutxe-Barghorn, D. J. Lougnot, New insights into
photoinduced processes in hybrid sol-gel glasses containing modified titanium
alkoxides, New J. Chem., 25 (2001) 1006-1014
104

Hiroshima University Doctoral Dissertation

29.

O. Soppera, C. Croutxe-Barghorn, Real-time Fourier transform infrared studyof free-radical UV-induced polymerization of hybrid sol-gel. I. Effect ofsilicate backbone on photopolymerization kinetics, J. Polym. Sci. A1., 41(2003)716-724

30.

J. M. Guissi, I. Blaszyczyk-Lezak, B. Sanz, P. E. Allegretti, C. Mijangos, M.
S. Cortizo, Tautomeric acetoacetate monomers as building units of functional
copolymers, Eur. Polym. J., 59 (2014) 84-93

31.

K. Grzybowska, Z. Wojnarowska, A. Grzybowski, M. Paluch, J. M. Guissi,
M. S. Cortizo, I. Blaszyczyk-Lezak, C. Mijangos, Effect of polymer structureon the molecular dynamics and thermal behavior of poly (allyl acetoacetate)
and copolymers, Polymer, 55 (2014) 1040-1047

32.

S. Lawal, M. Kanezashi, H. Nagasawa, T. Tsuru, Development of anacetylacetonate-modified silica-zirconia composite membrane applicable togas separation, J. Membr. Sci., 599 (2020) 117844

33.

M. ten Hove, A. Nijmeijer, L. Winnubst, Facile synthesis of zirconia doped
hybrid organic inorganic silica membranes, Sep. Purif. Technol., 147 (2015)372-378

34.

N.

Moriyama,

H.

Nagasawa,

M.

Kanezashi,

T.

Tsuru,

Bis(triethoxysilyl)ethane (BTESE)-derived silica membranes: pore size
formation mechanism and gas permeation properties, J. Sol-Gel Sci. Technol.,
86 (2018) 63-72
35.

S. T. Balke, T. H. Mourey, Local polydispersity detection in size exclusion
chromatography: method assessment, J. Appl. Polym. Sci., 81 (2001) 370-383

36.

M. Danaei, M. Dehghankhold, S. Ataei, F. H. Davarani, R. Javanmard, A.
Dokhani, S. Khorasani, M. R. Mozafari, Impact of particle size andpolydispersity index on the clinical applications of lipidic nanocarrier systems,
Pharmaceutics, 10 (2018) 57

37.

X. Le Guevel, C. Palazessi, P. Prosposito, G. D. Giustina, G. Brusatin,
Influence of chelating agents on the photopolymerization of hybrid Ti-based
waveguides J. Mater. Chem., 18 (2008) 3556-3562

38.

A. Matsumoto, T. Kumagai, H. Aota, H. Kawasaki, R. Arakawa,
Reassessment of free-radical polymerization mechanism of allyl acetate basedon end-group determination of resulting oligomers by MALDI-TOF-MS
spectrometry, Polym. J., 41 (2009) 26-33
105

Hiroshima University Doctoral Dissertation

39.

R. Joseph, W. T. Ford, S. Zhang, M. P. Tsyurupa, A. V. Patsukhov, V. A.
Davankov, Solid-state 13C-NMR analysis of hypercrosslinked polystyrene, J.
Polym. Sci. A: Polym. Chem., 35 (1997) 695-701

40.

B. Li, R. Gong, Y. Luo, B. Tan, Tailoring the pore size of hypercrosslinked
polymers Soft Matter, 7 (2011) 10910-10916

41.

C. Morterra, G. Cerrato, C. Emanuel, End-on surface coordinated (adsorbed)
CO2: a specific ligand for surface Lewis acidic centres, Mater. Chem. Phys.,
29 (1991) 447-456

42.

S. Lawal, L. Yu, H. Nagasawa, T. Tsuru, M. Kanezashi, A carbon-silicazirconia membrane with CO2 flow-switching behavior promising hightemperature H2/CO2 separation, J. Mater. Chem. A. 8 (2020) 23563

43.

H. H. Horowitz, G. Metzger, A new analysis of thermogravimetric traces,
Anal. Chem., 35 (1963) 1464-1468

44.

A. Gronowski, Z. Wojtczak, The thermal decompositions of some transition
metal acrylates and polyacrylates, J. Thermal Anal., 26 (1983) 233-244

45.

H. Schmidt, Sol-gel: science and technology; proceedings of the winter schoolon glass and ceramics from gels, São Carlos (SP), Brazil, 14 - 19 August
1989. Ed. by M. A. Aegerter. Singapore: World Scientific. 1989, 432-469

46.

H. R. Lee, M. Kanezashi, Y. Shimomura, T. Yoshioka, T. Tsuru, Evaluationand fabrication of pore-size-tuned silica membranes with tetraethoxydimethyl
disiloxane for gas separation, AIChE J., 57 (2011) 2755-2765

47.

P. Hacarlioglu, D. Lee, G. V. Gibbs, S. T. Oyama, Activation energies forpermeation of He and H2 through silica membranes: An ab initio calculation
study, J. Membr. Sci., 313 (2008) 277-283

48.

K. Haraya, K. Obata, T. Hakuta, H. Yoshitome, The permeance of gases
through a new type of polyimide membrane, Membrane, 11 (1986) 48-52

49.

Y. Yampolskii, S. Shishaskii, A. Alentiev, K. Loza, Correlations with andprediction of activation energies of gas permeation and diffusion in glassy
polymers, J. Membr. Sci., 148 (1998) 59-69

50.

K. Haraya, K. Obata, T. Hakuta, T. Yoshitome, Permeation of gases through asymmetric cellulose acetate membrane, J. Chem. Eng. Jpn., 19 (1986) 464-466

51.

X. Feng, P. Shao, R. Y. M. Huang, G. Jiang, R. -X. Xu, A study of silicone
rubber/polysulfone composite membranes: H2/N2 and O2/N2 permselectivities,
Sep. Purif. Technol., 27 (2002) 211-223
106

Hiroshima University Doctoral Dissertation

52.

H. L. Castricum, G. G. Paradis, M. C. Mittelmeijer-Hazeleger, R. Kreiter, J. F.
Vente, J. E. ten Elshof, Tailoring the separation behavior of hybridorganosilica membranes by adjusting the structure of the organic bridging
group, Adv. Funct. Mater., 21 (2011) 2319-2329

53.

M. Kanezashi, Y. Tomarino, H. Nagasawa, T. Tsuru, Tailoring the molecularsieving properties and thermal stability of carbonized membranes containing
polyhedral oligomeric silsesquioxane (POSS)-polymer via the introduction of
norbornene, J. Membr. Sci., 582 (2019) 59-69

54.

H. Nagasawa, Y. Yamamoto, M. Kanezashi, T. Tsuru, Atmospheric-pressureplasma-enhanced chemical vapor deposition of hybrid silica membranes, J.
Chem. Eng. Jpn., 51 (2018) 732-739

55.

M. Kanezashi, Y. Yoneda, H. Nagasawa, T. Tsuru, Gas permeation properties
for organosilica membranes with different Si/C ratios and evaluation of
microporous structures, AIChE J., 63 (2017) 4491-4498

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Chapter 4
Enhanced H2-permselectivity in a carbon-silica-zirconia membrane
derived from an acetylacetonate-ceramic composite pyrolysis
4.1 Introduction
A 2015 report estimates that the world’s energy demand satisfied by fossil fuelssuch as gas, coal and oil stand at about 85 % [1]. Global annual energy demand of 12
billion tons equivalent of oil (BTOE) results in the emission of about 40 giga tons of
CO2 [1] while the energy forecasted to be demanded by just 8 countries (China, theUSA, India, Brazil, Japan, South Korea, Canada, and Mexico) will be as high as 12.5
BTOE by 2035, the equivalent of 2015 total CO2 emission-resulted consumptions [2].
With the continued dependence on fossil fuel and the ever-increasing demand ofenergy, the danger of increased CO2 emission looms. Moreover, other factors driving
the move towards alternative fuel sources are energy security and climate change [3].
Hydrogen is widely considered as a sustainable and clean alternative source of energy[3, 4].
Hydrogen not being a primary energy source needs to be produced from othersources. The chemical industry serves as a good source of hydrogen production [5].
However, it is required to purify hydrogen from other gases encountered in typicalindustrial gas mixtures. Several techniques exist for gas separation in the chemicalindustries including but not limited to traditional processes such as pressure swing
adsorption (PSA), cryogenic distillation [6-8], and in recent decades the energy
conservative membrane separation [7, 8].
Diverse membrane types such as polymeric membranes, metallic membranes,
porous inorganic membranes (silica, zeolites, metal oxides) and carbon molecularsieve membranes have been studied for hydrogen separation and extensive reviews
have been made [7, 9]. Porous inorganic membranes are considered as having
superior molecular permselectivity compared to polymeric membranes, durable overlonger periods and can be operated at much higher temperatures [10]. Particularly,
silica membranes have shown very high H2-permselective ability [11, 12]. However, asignificant problem with pure silica-derived membranes has to do with network
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densification at elevated temperatures [10, 13]. According to the work of de Vos andVerweij, H2 permeance in a silica membrane prepared at 600 oC was about 5 x 10-7
mol m-2 s-1 Pa-1 at 200 oC which is one order of value lower compared to thatmeasured for a silica membrane prepared at 400 oC (1.7 x 10-6 mol m-2 s-1 Pa-1) [11].
Lin also demonstrated that by subjecting an unsupported silica to 600 oC for 30 hours,
the pore volume of silica reduced significantly [13]. The densification of silica at very
high temperatures is attributed to the continued thermal condensation of -Si-OH
groups [13, 14]. Vahlas and co-workers processed silica films at 400 oC and observed
a continuous decrease in the -Si-OH peaks detected by FTIR as temperature increased
until completely disappearing at 550 oC [14].
Several techniques have thus been proposed to solve this issue. Such includes
spacer methods (incorporating an alkyl group to bridge Si atoms in the precursor) [15,16], templating methods (incorporating unhydrolyzable organic group attached as apendant to the Si atom) [17, 18], anion/cation doping [19, 20] and the use of stable
metal oxides such as Al2O3, ZrO2 and TiO2 as co-composites [21-23]. At temperatures
reaching 550-600 oC, the most potent means of inhibiting the thermal densification ofsilica is by compositing with metal oxides because of their extreme thermal stability.
The modification of pure silica with a metal oxide to form a ceramic composite has
been well studied and the most popular oxide being ZrO2 [22, 24, 25]. The stability ofsilica is greatly improved, and stability improvement increased as Zr content wasincreased. Yoshida and co-workers [22] as well as Choi and co-workers [24] showedthat for the best results, a mixture ratio of 1:1 was desired. A SiO2-ZrO2 (1:1)
membrane prepared by Choi and co-workers by the CVD method at 600 and 650 oC
only reduced in H2 permeance from 7.3 to 2.7 x 10-8 mol m-2 s-1 Pa-1. Thereafter, themembrane kept a stable value at 700 oC. However, increasing the ZrO2 content
resulted in loss of selectivity and a low initial H2 permeance [22]. This is borne out ofthe difficulty in achieving an ideal mixed composite because the hydrolysis rate of thezirconia precursor is much faster than a silica precursor [26]. This can result in aquick formation and aggregation of ZrO2 nanoparticles thereby forming an unevencomposite.
Therefore, to achieve even hydrolysis reactivity, a modification of the zirconium
precursor with an organic coordination ligand has been suggested and proven [27, 28].
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Utilizing organic coordination/complexing ligands offers another advantage asreported by Spijksma and co-workers [29]. They showed that by modifying Ti and Zr
atoms in a TiO2-ZrO2 composite with diethanolamine (DEA), a membrane’s
permselectivity could be improved. Since then, the successful improvement of a SiO2ZrO2 composite membrane with an acetylacetone ligand has been reported [30]. But
the ligands do not remain intact at the elevated temperatures in which a SiO2-ZrO2
composite membrane is needed.
Recently, our group studied an acetylacetonate-modified SiO2-ZrO2 composite(1:1) after carbonization by the pyrolysis of the acetylacetone ligand at 550 oC in anitrogen atmosphere [31]. It was confirmed that the acetylacetone ligands pyrolyzed
to form free carbon nanoparticles within the SiO2-ZrO2 composite network.
Characterizations revealed an interesting electrostatic interaction between the freecarbon nanoparticles and CO2 whereby a pressure-induced flow switching of CO2
could be observed. Combined with an improved H2 permeance over an unmodified
SiO2-ZrO2 composite membrane, enhanced H2/CO2 separation ability was achieved.
However, the pyrolysis of the highly concentrated acetylacetonate ligands resulted information of faults in the membrane. Hoebbel and co-workers however demonstrated
that by utilizing the same stoichiometric water/alkoxide molar ratio during thehydrolysis process, different β-diketone and β-ketoester-type coordination ligands
chelated to Zr displayed different hydrolytic stability with the most stable being
acetylacetone [32]. It is thus reasonable to theorize that utilizing different
water/alkoxide molar ratios beyond the stoichiometric amount for hydrolysis could
tune the eventual concentration of an acetylacetone ligand within the resulting SiO2ZrO2 composite. This may be able to control the extent and severity of the defectforming pyrolysis.
In this study, the combination of compositing metal oxide and carbonization ofacetylacetone ligand for the modification of silica microstructure was investigated.
Specifically, the effect of hydrolysis water/alkoxide molar ratio (WR) on the resulting
carbonized SiO2-ZrO2 composite network was investigated and optimized. Based on
the optimized water ratio, the synergistic combination of the relative amounts of SiO2
over ZrO2 and the residual free carbon for H2-permselectivity was investigated andcompared to the state-of-the-art sol-gel-derived SiO2-ZrO2 membrane. It is resolved
110Hiroshima University Doctoral Dissertationthat free residual carbon deposited within the SiO2-ZrO2 composite provided
microstructural benefits.

4.2 Experimental
4.2.1 Materials preparation by sol-gel
5 wt% in ethanol medium of SiO2-ZrO2-acetylacetone (represented as SZA
hereafter) sols were prepared according to a previous report [30]. In this work
however, two different sol-gel experiments were made.
In the first experiment, SZA sols of different Si/Zr molar ratios were prepared.
First, 80% pure zirconium (iv) tert-butoxide (ZrTB, Aldrich) dissolved in ethanol(Aldrich) was modified with a 99% pure acetylacetone (ACAC, Aldrich)(ACAC/ZrTB molar ratio = 4/1) for one hour at room temperature. Subsequently, a
98% pure tetraethoxysilane (TEOS, Aldrich) dissolved in ethanol was co-hydrolyzed
with the previously acetylacetone-modified ZrTB in four places so that the atomic
ratios of Si/Zr were 9/1, 7/3, 5/5 and 3/7.

For hydrolysis, deionized water(water/alkoxide molar ratio = 60) and hydrochloric acid (37%, Nacalai Tesque) as acatalyst (acid/alkoxide molar ratio 0.25) were used. Hydrolysis and poly-condensation
were carried out in a closed vial by stirring the mixture at 550 rpm for more than 12hours at room temperature.
In the second sol-gel experiment in a similar manner but fixing the atomic ratio of
Si/Zr at 5/5, three different SZA sols were prepared depending on the water/alkoxide
molar ratio (WR). Hydrolysis proceeded using three different WR of 4, 60 and 240.
The acid/alkoxide ratio was still fixed at 0.25.
Films, powders and membranes were subsequently prepared from these sols. To
prepare films, 50 µl of the respective sols were drop-coated on UV-pretreated siliconwafers and allowed to dry at room temperature. The dried films were then calcined at
the needed temperatures before characterization. Powders were prepared by a slow
drying process whereby the sol was allowed to dry freely under a controlledtemperature range of 50-80 oC giving rise to unsupported xerogels. These were
subsequently crushed and calcined appropriately. All calcinations were done in a N2

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atmosphere. The samples prepared at 300 oC are referred to as SiO2-ZrO2-ACAC(SZA) while those prepared at 550 oC are referred to as carbon- SiO2-ZrO2 (CSZ).
4.2.2 Preparation of membrane separation layers
Membranes were fabricated on support layers fabricated according to thedescription in section 2.2.1. Two types of membrane separation layers were fabricatedbased on the calcination temperature. The first type was the membrane derived from
SiO2-ZrO2-ACAC. A 5 wt% SiO2-ZrO2-ACAC sol prepared as described earlier was
diluted to a 0.5 wt% with ethanol and wipe-coated onto the SiO2-ZrO2 intermediate
layer. This was followed by firing in 600 ml min-1 nitrogen gas stream at 300 oC for
20 min. This process was repeated several times to obtain the top separation layer.
The second type of membrane layer was the carbon-SiO2-ZrO2 separation layer
based on a carbonization/pyrolysis process. The same 0.5 wt% sol was wipe-coatedon the SiO2-ZrO2 intermediate layer in the same manner. After this, the membrane
was allowed to dry at 50 oC in a stream of nitrogen gas for a few minutes before flash
firing at the carbonization temperature of 550 oC for 20 minutes. This process was
also repeated several times.
4.2.3 Material characterizations
The particle sizes of the as-prepared sols were measured at room temperature via
dynamic light scattering (DLS) using a Malvern Zetasizer Nano ZS (Malvern
Instruments Ltd.). The presence and intensities of different functional groups weredetected on film-coated silicon wafers via Fourier Transform Infrared Spectroscopy(FTIR, FTIR-4100, JASCO, Japan). The thermogravimetric profiles of powders wereanalyzed using Differential Thermogravimetry/Thermogravimetric Analysis (DTGTGA, TG-50 Shimadzu Co., Japan) under nitrogen or air gas flow of approximately
60 ml min-1 and a heating rate of 10 oC min-1.
Furthermore, the lattice structures were analyzed using the x-ray diffractometry(XRD, D2 PHASER X-Ray Diffractometer, Bruker, Germany) with Cu Kα as theradiation source at a wavelength of 1.54 Å. The microporosities of the powders were
evaluated by N2 adsorption-desorption experiments at -196 oC (BELSORP MAX,112

Hiroshima University Doctoral Dissertation

BEL Co., Japan). Prior to this measurement, samples were evacuated of adsorbed
gases and vapors at 200 oC for at least 12 hours.
4.2.4 Evaluation of membrane gas permeation characteristics
A schematic representation of the gas permeation measurement setup is shown in
Fig. 2-2. Water vapor adsorbed into the membrane pores after fabrication wasdesorbed by flowing He gas through the membrane module setup at 200 oC for 12hours.
The single gas performance of membranes was evaluated with high-purity gases(H2, He, CO2, N2, CH4, CF4, SF6 in that order) similar to the procedure described in
section 2.2.3.

4.3 Results and discussion
4.3.1 Control and optimization of ligand chelation degree via hydrolysis
Since Hoebbel et al [32] demonstrated that more than 97% of ACAC chelating
ligand remained bonded to Zr after 7 days of hydrolysis with a water/alkoxide molar
ratio of 6, ACAC ligand is a very hydrolytically stable chelating ligand. Therefore, by
replacing the butoxyl groups in the Zr precursor by ACAC chelating ligands, thehydrolysis rate can be considerabely slowed down [26, 27]. Nonetheless, it isexpected that this resistance can be finely controlled by varying the water/alkoxide
molar ratio other than the stoichiometric amount and the concentration of the ligand
chelated to Zr in a SiO2-ZrO2-ACAC composite may be tuned. Three different
water/alkoxide molar ratios (WR) have been adopted: 4, 60 and 240 to determine theoptimal WR required to form a carbon-SiO2-ZrO2 membrane. On one hand, it isanticipated that a lower WR will result in retaining a higher amount of ligand andtherefore possibly result in a higher amount of residual carbon, on the other hand, ahigher WR result in higher microporosity for a better molecular sieving effect. For
this purpose, the ratio of Si to Zr was kept at 1:1.
4.3.1.1 Characterization of sols, films, and powders before carbonization
Fig. 4-1 shows the sol size measured by intensity for SZA sols prepared at WR 4,
60 and 240. The mean values and standard deviations are presented in Table 1. It canbe observed that the sol size increased as WR increased from 4 (3.4±1.0 nm) to 60
113

Hiroshima University Doctoral Dissertation(8.1±1.5 nm) and then to 12±0.3 nm for a WR of 240. This therefore points to morecondensations between hydrolyzed groups (-SiOH and -ZrOH) as WR increased
resulting from increased hydrolysis.

Intensity [%]

12
8

WR240
WR60
WR4

4
0 -1
10

10

0

10

1

10

2

Size [nm]Fig. 4-1 Sol size by intensity measured by DLS for different water/alkoxide molar
ratios.
Table 4-1 Average sol size by intensity values of SiO2-ZrO2-ACAC sols prepared at
different water/alkoxide molar ratios.
Water ratio

Average sizeStandard deviation (+/-)

[-]

[nm]

4

3.4

1.0

60

8.1

1.5

240

12

0.3

FT-IR provides definitive evidence for the observation made with the sol size
measurement. Fig. 4-2 (a) shows the observed FT-IR spectra in the range of 3200 to
400 cm-1 measured for as-prepared SZA films before and after firing at 300 oC in N2
atmosphere. The two peaks that appear around 1,600 cm-1 and 1,525 cm-1 are assigned
to the enol C=O and C=C bonds that are characteristic of the acetylacetone,
respectively [34]. The peak observed around 1000 cm-1 can be assigned to thesiloxane bond -Si-O-Si- [35]. This indicates that hydrolysis-produced -Si-OH species
were able to undergo condensation with other -Si-OH species or -Si-OEt species. The
spectra of the unfired films show that acetylacetone remained chelated irrespective of
the WR used only differing in intensity. Fig. 4-2 (b) shows the plot of evaluated114Hiroshima University Doctoral Dissertation

absorbance ratios of the -C=C- peak of chelating ACAC ligand to siloxane as afunction of water/alkoxide molar ratios. The -C=C- peak intensity reduced as WR
increases. Therefore, the increased hydrolysis and condensation between SiO2 andZrO2 inferred from Fig. 4-1 indicates that the amount of ACAC chelated to Zr was
reduced by increasing the WR.
-C=O -C=C--CH2-(b) 1.50

WR240

1.25
A-C=C-/A-SiOSi- [-]Absorbance [a. u.](a)-SiOSi-WR60

1.00
Uncalcined

0.75
0.50

o

Calcined (300 C)

WR4

0.25
3200 2800 2400 2000 1600 1200 800

400

0

50

100

150

200

250

Water molar ratio [-]-1

Wavenumber [cm ]Fig. 4-2 (a) FT-IR spectra observed within a range of 3200- 400 cm-1 for uncalcined(green line) and calcined (300 oC in N2; blue line) SiO2-ZrO2-ACAC films preparedwith different water/alkoxide molar ratios; (b) plot of the evaluated absorbance ratiosof the representative peaks of the chelating ACAC ligand to siloxane as a function ofwater/alkoxide molar ratios.
The same comparisons were made for films fired at 300 oC in N2 atmosphere(represented by the blue solid lines). After firing at 300 oC, the intensities of therepresentative acetylacetone peaks reduced irrespective of WR. This is consistent withthe thermogravimetric profile reported for a SZA powder prepared with a WR of 4
which showed weight loss at 300 oC [30]. However, the characteristic peaks
belonging to ACAC can still be observed and the relative ACAC content followed thesame trend as the unfired films.
Fig. 4-3 shows the N2 adsorption-desorption isotherms for SZA powders preparedwith different water/alkoxide molar ratios and fired at 300 oC. This figure shows thatthe amount of N2 adsorbed increased as WR for hydrolysis increased. At WR of 4 theamount of N2 adsorbed was small because the high amount of ACAC ligands present
115

Hiroshima University Doctoral Dissertationin the SiO2-ZrO2 matrix based on Fig. 4-2 (b). These occupied the intra-particle
network spaces that allowed only a little amount of N2 adsorbed. At WR of 240
however, the high amount of hydrolysis that caused a reduced amount of ACAC
ligands improved the amount of N2 adsorbed. Nonetheless, WR of 240 still resulted in

3-1

Amount of adsorbed [cm (STP) g ]the highest microporous volume.

30
25

Solid: adsorption
Open: desorption

WR240

20
15
10

WR60

5

WR4

0
0.0

0.2

0.4
0.6
p/po [-]

0.8

1.0

Fig. 4-3 N2 adsorption-desorption isotherms at -196 oC for SiO2-ZrO2-ACAC powders
prepared with different water/alkoxide molar ratios and fired at 300 oC.
The gas permeation properties of membranes fabricated from SZA at the different
WR were also examined at 200 oC. Fig. 4-4 (a) shows the kinetic diameter
dependence of single gas permeance of different gases (He (0.26 nm), H2 (0.289 nm),
CO2 (0.33 nm), N2 (0.364 nm), CH4 (0.38 nm), CF4 (0.48 nm), and SF6 (0.55 nm)) for
SZA-derived membranes fabricated at 300 oC. Fig. 4-4 (b) shows the dimensionlesspermeance normalized on the basis of He permeance in Fig. 4-4 (a) as a function of
kinetic diameter. It is obvious that the permeance values of all gases increased as WR
increased from 4 to 240. However, the ideal selectivity of H2 over larger gases
reduced with WR but still well above Knudsen-based values as presented in Table 42.

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Hiroshima University Doctoral Dissertation
He H2 CO2 CH4

CF4

SF6

Dimensionless permeance [-]

N2

-1

-6

2

10
10

-7

10

-8

10

-9

10

WR 4
WR 60
WR 240

-10

2

3

4

5

He H2 CO2 CH4

(b)

-1

Permeance [mol m s Pa ](a) 10-5

6

10
10

-1

10

-2

10

-3

10

-4

CF4

SF6

N2

0

WR 4
WR 60
WR 240

2

3

4

5

6

Kinetic diameter [Å]Kinetic diameter [Å]Fig. 4-4 (a) Single gas permeance at 200 oC as a function of kinetic diameter of SiO2ZrO2-ACAC-derived membranes fabricated at 300o

C under N2 for different

water/alkoxide molar ratios. (b) corresponding dimensionless permeance as a functionof kinetic diameter normalized by He permeance.
Table 4-2 Gas permeance of H2, N2 and SF6 gases at 200 oC; ideal selectivity of H2
over N2 and SF6 for SiO2-ZrO2-ACAC-derived membranes prepared with different
water/alkoxide molar ratios.
Permeance [10-7 mol m-2 s-1 Pa-1]

Ideal selectivity [-]

H2

N2

SF6

H2/N2

H2/SF6

4

5.2

0.2

0.003

30

1600

60

8.8

0.6

0.032

14

280

240

22

2.8

0.21

8

110

Knudsen

-

-

-

4.3

8.5

WR

Furthermore, Fig 4-4 (b) which can reveal the pore size distribution comparison of
the membranes showed that a membrane prepared with a WR of 4 showed thenarrowest distribution compared with water/alkoxide molar ratios of 60 and 240. This
trend is opposite to that observed for organosilica membranes whereby an increase in
WR increased the selectivity of membranes. For example, a triethoxysilane (TRIES)derived membrane showed that membranes prepared with water ratio of 240 exhibited
the narrowest pore distribution compared to a WR of 60 and 6 (having the broadest
117

Hiroshima University Doctoral Dissertationpore distribution) as reported by Kanezashi et al. [36]. Kanezashi and co-workers
opined a “template effect” to explain this phenomenon. At low water ratio, theincomplete hydrolysis of ethoxy groups (-SiOEt) resulted in the persistence of these
ethoxy groups in the networks. At a high water/alkoxide molar ratio however, these
ethoxy groups were completely hydrolyzed and therefore replaced by an abundance ofhydroxyl groups (silanol -Si-OH). After calcination at elevated temperatures, thesilanol (-Si-OH) rich, high WR-derived membrane underwent a thermal condensationof silanol into siloxane crosslinks thereby resulting in a -Si-O-Si- dense network. On
the other hand, after calcining an ethoxy-rich network derived from low WR, thethermally unstable ethoxy groups simply decompose without a condensation
crosslinking, resulting in a loose network. This trend was also reported for abis(triethoxysilyl propoxypyridine) (BTPP)-derived membrane by Yu and co-workers
utilizing water/BTPP ratios of 12, 40 and 240 [37].
In a chelating ligand-ceramic composite however, the mechanism of thermal
network formation is different. Fig. 4-5 summarizes the effect of water/alkoxide
molar ratio on the ligand concentration and network formation in SiO2-ZrO2-ACAC
before and after calcination. FT-IR results in Fig 4-2 (a) and (b) revealed that evenafter calcination at 300 oC, a substantial amount of ligands was still present. A WR of
4 retains the highest amount of ligands, and these network-occupying ligands result ina quasi-dense network that behaves like a polymer-derived structure [38] that gives anarrow pore distribution. On the other hand, when the WR is increased, the amount ofhydroxyl groups increases but with a persistent presence of ligands in the network as
Fig. 4-5 illustrates. At the calcination temperature, some -OH groups are removed but
condensation to form -Si-O-Si-, -Si-O-Zr- or -Zr-O-Zr- linkages is inhibited because
the ligands still present prevented the proximity of condensable species. Therefore, it
might be reasonably concluded that for a coordination ligand-ceramic composite
network, an increase in WR results in membranes with looser network and vice versa.

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Hiroshima University Doctoral Dissertation
Water/alkoxide
60

Water/alkoxide
4
OEt
Si
O Zr
O
O Si
OEt
O
Zr
Si
O
O
O
O
Zr
O
Si
Zr
Si
O
O
O Zr
O Si
Calcination
OEt
Si
O Zr
O
O Si
O
Zr
Si
O
O
O
O
Zr
O
Si
Zr
Si
O
O
O Zr
O Si

Water/alkoxide
240

OH

O Si
OEt

O

Zr

Zr

OH

O Zr

O

Si

O

OH
O

Si HO
O
O Zr

O

O Si

Si

O
Zr O Si

O Si
O
OH
OH Si
O
O
O
O
Zr
O
Si
Zr
Si HO OH
O
O
O Zr
O Si
Zr

OH

O
Si

O

Calcination

O

Zr

Zr

Si

O
O

OH

O Zr

O

OH

Calcination

OH

O Si

O Zr

O

Si

O Zr

O

O Si

Si

O
Zr O Si

OH

O

Zr

O
O

Zr

Si

O

O Zr
Si

O

Si

O
Zr O Si
Si
O
O
O
O Zr
O Si

O

O

Fig. 4-5 Schematic illustration of the effect that the water/alkoxide molar ratio exerts
on the relative ligand population before and after calcination.
4.3.1.2 Characterization after carbonization
The results of characterizations before carbonization have revealed that theconcentration of ACAC ligand in the membrane can be successfully determined by
the water/alkoxide molar ratio applied for hydrolysis and affects SZA membranes’
pore size distribution. However, the optimal WR required for a carbonized SiO2-ZrO2
can only be determined by analyzing the results of characterization after carbonization
at 550 oC.
The amount of residual carbon after 550 oC pyrolysis can be evaluated bymeasuring the thermogravimetry in N2 and air atmospheres. Fig. 4-6 (a) – (c) show
the thermogravimetric profiles in N2 and air. The amount of residual carbon(evaluated as mg of carbon per mg of residual material) is assumed to be thedifference between the normalized residual weight in N2 at 550 oC and the normalized
residual weight in air at 1000 oC after all carbon is removed. From Fig. 4-6 (a) – (c) it
can be observed that the amount of residual carbon estimated was the highest for WRof 4 and smallest for the WR of 240. This is consistent with the ACAC ligand content

119Hiroshima University Doctoral Dissertationfrom the precursor material. Therefore, WR of 240 is not suitable to retain a high

WR4

0.9

N2
Air

0.8
0.7
0.068 mg/mg

0.6
0.5
0

(a)

1.0
WR60

0.9

N2
Air

0.8
0.7
0.061 mg/mg

0.6

0.5
200 400 600 800 1000 0(b)

Normalised residual weight [-]1.0

Normalised residual weight [-]Normalised residual weight [-]amount of carbon material in the SiO2-ZrO2 matrix.
1.0
WR240

0.9

Air

0.8

0.6

o

Temperature [ C]0.053 mg/mg

0.7

0.5
200 400 600 800 1000 0

o

N2

(c)
200 400 600 800 1000oTemperature [ C]

Temperature [ C]

Fig. 4-6 Thermogravimetric profiles of SiO2-ZrO2-ACAC powders under N2 (solidline) and air (dash line) up to 1000 oC.
Furthermore, the XRD diffractograms of carbon-SiO2-ZrO2 powders prepared at
550 oC as shown in Fig. 4-7 (a) and (b) shows that the WR of 240 resulted inaggregation of tetragonal ZrO2 nanoparticles. WR of 4 and 60 retained amorphous
structures when compared to the powders before carbonization shown in Fig. 4-7 (a).
Therefore, WR 4 and 60 are more likely to form defect-free thin separation layers.
However, despite a WR of 4 retaining the highest residual carbon but resulting in avery low microporosity as shown in Fig. 3, the WR of 60 serves as the best balance
between microporosity, amorphousness and amount of residual carbon.

(b)
WR240

WR60

WR60

20 30

40 50 60 70o
2 theta [ ] Tetragonal ZrO2





30 40 50o
2 theta [ ]

60

WR4

WR4

10



WR240

Intensity [a. u.]

Intensity [a. u.]

(a)

80 10

120

20

70

Hiroshima University Doctoral Dissertation

Fig. 4-7 XRD patterns of (a) SiO2-ZrO2-ACAC powders prepared at different
water/alkoxide molar ratios after firing under a N2 atmosphere at 300 oC (b) carbonSiO2-ZrO2 powders prepared at different water/alkoxide molar ratios after firing undera N2 atmosphere at 550 oC.
4.3.2 Microstructural transformation between SiO2-ZrO2-ACAC and CarbonSiO2-ZrO2: effect of Si/Zr ratio
An optimal water/alkoxide molar ratio for preparing SiO2-ZrO2-ACAC as the
precursor to carbon-SiO2-ZrO2 has been chosen to be 60 at a fixed ratio of Si to Zr
atoms of 1:1. However, the effects of the relative ratios of Si to Zr at a fixed WR of
60 on the microstructural properties of CSZ-derived bulk materials and membranes isvery important. It is important to note that the theoretical amount of ACAC present ina given SZA matrix can also be dictated by the ratio of Si to Zr atoms present when at
any given time the theoretical ACAC/Zr ratio is kept at 4:1. Therefore, in a silica-richmaterial the theoretical ACAC amount reduces and, for a zirconia-rich material, theACAC amount increases, all while keeping the theoretical ACAC/Zr ratio at 4:1.
Fig. 4-8 shows the thermogravimetric profiles in an N2 atmosphere for as-prepared
SZA powders with different Si/Zr atomic ratios of 3/7, 5/5, 7/3 and 9/1. The final
residual weight represents the undecomposable residual mass in the given thermal
atmosphere. In consistence with the previous discussion, the amount of material
decomposed corresponded with the amount of ACAC contained in the starting
material. Therefore, the residual weight increased as Si/Zr ratio increased- the mostsilica-rich SZA powder (Si/Zr of 9/1) shows the highest residual weight and the mostzirconia-rich SZA powder (Si/Zr of 3/7) shows the lowest residual weight. The CSZ
materials considered in this work were prepared at 550 oC and the above observations
hold at 550 oC on the thermogravimetric profile. However, the absolute amount of theresidual carbon at 550 oC cannot be determined even if the residual weights
correspond to the initial ACAC amount at each Si/Zr ratio.

121

1.0
0.9

In N2

0.8
0.7
0.6

3/7
5/5
7/3
9/1

0.5
0

200

400

600

(b)

1.00

Residual weight [-](a)Residual weight [-]Hiroshima University Doctoral Dissertation

0.95

In air

0.90
0.85
0.80
0

800 1000
o

3/7
5/5
7/3
9/1

200

400

600

800

o

Temperature [ C]

Temperature [ C]

Fig. 4-8 Thermogravimetric profiles of (a) SiO2-ZrO2-ACAC powders (uncalcined)
prepared at different Si/Zr atomic ratios under a N2 atmosphere up to 1000 oC and (b)carbon-SiO2-ZrO2 powders (prepared at 550 oC under N2) under an air atmosphere up
to 800 oC. (Gas flow: 50-60 ml min-1; Ramping rate: 10 oC min-1.)
CSZ powders prepared by pyrolyzing SZA powders in N2 atmosphere at 550 oC
were thereafter subjected to a thermogravimetric process in an air atmosphere. The
hot air atmosphere resulted in the complete oxidation and removal of the carbon
nanoparticles in the CSZ powders. Therefore, the amount of mass reduction
corresponds to the residual carbon in the starting CSZ powders prepared at 550 oC.
Table 4-3 presents the results of this thermogravimetric experiment. The general trend
is that the residual carbon amount increases with the zirconia mole fraction. For themost zirconia-rich precursor where the theoretical C/Si atomic ratio is 28/3, the result
is the highest residual carbon of 0.103 mg of carbon per mg of CSZ. The most silicarich precursor with theoretical C/Si atomic ratio of 4/9 resulted in the lowest residualcarbon of 0.065 mg of carbon per mg of CSZ.

122

Hiroshima University Doctoral Dissertation

Table 4-3 Evaluation of carbon content in carbon-SiO2-ZrO2 powders prepared at 550oC

Si/Zr atomic ratio

Theoretical starting C/Si atomic ratio

Residual carbon at 550 oC

[-]

[-]

[mg/mg]

3/7

28/3

0.103

5/5

20/5

0.081

7/3

12/7

0.083

9/1

4/9

0.065

The microstructural transformation between SZA and CSZ was also examined by
x-ray diffraction (XRD). Fig. 4-9 shows the x-ray diffractograms of SZA (green lines)
powders and CSZ powders (red lines) prepared at different Si/Zr atomic ratios. Before
carbonization, all the spectra of SZA show amorphous structures irrespective of Si/Zr
atomic ratio. A broad peak is common to all the samples before carbonization whichshifts very slightly from higher 2θ to a lower 2θ as Si/Zr ratio increases. The
calculated d-spacing values are presented in Table 4-4 and reveal that the d-spacing
value of 0.37 nm for SZA9/1 is only slightly higher than that of SZA3/7 at 0.34 nm.
Nonetheless, the slightly lower d-spacing of SZA3/7 may be attributed to the lower
inter-plane distance caused by the higher concentration of ACAC ligands between thelattices.

(d) 9/1

Intensity [a. u.](c) 7/3(b) 5/5(a) 3/7

10

20

30
o

2 theta [ ]

123

40

Hiroshima University Doctoral Dissertation

Fig. 4-9 X-ray diffractograms of SiO2-ZrO2-ACAC powders (before carbonization;
green) and Carbon-SiO2-ZrO2 powders (after carbonization in N2 at 550 oC; red)prepared at different Si/Zr atomic ratios. Symbol ◊ represents the appearance of
aggregated tetragonal ZrO2 particles.
The results recorded after for the CSZ powders reveal drastically different
observations. It is important to note that all but the Si/Zr ratio of 3/7 have completely
amorphous phases. CSZ3/7 showed a little aggregation of tetragonal ZrO2 particles
with a pointed peak at 2θ of 30 o.
According to the d-spacing calculated for CSZ powders a larger difference in thevalues can be observed compared to the uncarbonized SZA powders and the peakshifts are obvious in Fig 4-9. The d-spacing value of 0.37 nm for CSZ9/1 is muchhigher than that of CSZ3/7 at 0.29 nm. This indicates that a considerable
microstructural differentiation can be achieved once carbonization takes place. A
second interesting observation is revealed when comparing SZA and CSZ profiles at
each Si/Zr ratio (from green to red). For example, the d-spacing of SZA3/7 at 0.34 nm
is much higher than that of a corresponding CSZ3/7 at 0.29 nm. The CSZ7/3 powder
showed only a slight shift compared to the SZA7/3 precursor resulting in d-spacing
values of 0.32 nm and 0.35 nm, respectively. As for CSZ9/1, there was no observablepeak shift compared to SZA9/1 corresponding to a similar d-spacing of 0.37 nm in
both cases.
This result is also consistent with a N2 adsorption experiment. Fig. 4-10 presents
the N2 adsorption-desorption isotherms measured at -196 oC for all SZA and CSZ
powders with different Si/Zr ratios. The obtained BET surface areas are also shown in
Table 4-4. Generally, it can be observed that a reduction of surface areas occurs from
SZA powders to CSZ powders attributed to the densification of the networks byincreased -Zr-O-Zr- thermal condensations [39] and the possible deposition of the freecarbon products of pyrolysis. Nonetheless, this densification is more pronounced inzirconium rich cases than in the silica rich cases. SZA5/5 with BET surface area of 14
m2 g-1 became CSZ5/5 with Sa,BET of 0.7 m2 g-1, representing about 95 %densification. On the other hand, SZA9/1 with a Sa,BET of 456 m2 g-1 was transformed
to CSZ9/1 with Sa,BET of 413 m2 g-1, representing only about 9 % densification.
Therefore, it is noteworthy that in a silica rich case, even if densification occurs, no124

Hiroshima University Doctoral Dissertation

substantial change was effected on the ceramic composite superstructure. In contrast
to this, a pure silica bulk prepared by Lin et al showed a surface area reduction of
about 34 % and a reduction of 15 % for 6% aluminum-doped silica [21].

160

-1

Solid: adsorption
Open: desorption

9/1

120

3

200

(b)

80
7/3
40
0
0.0

3/7

5/5
0.2

0.4

Amount adsorbed [cm (STP) g ]3-1

Amount adsorbed [cm (STP) g ]

(a)

0.6

0.8

200
160

Solid: adsorption
Open: desorption

9/1

120
80
40
0
0.0

1.0

5/5

7/3
0.2

p/p0 [-]

0.4

0.6

3/7
0.8

1.0

p/p0 [-]

Fig. 4-10 N2 adsorption-desorption isotherms at -196 oC for (a) SiO2-ZrO2-ACAC and(b) carbon-SiO2-ZrO2 powders prepared with different Si/Zr atomic ratios
Table 4-4 Effect of Si/Zr atomic ratio on microstructural properties of SiO2-ZrO2ACAC and carbon-SiO2-ZrO2 powders
SiO2-ZrO2-ACAC

Carbon-SiO2-ZrO2

atomicd spacing

Sa,BETd spacing

Sa,BET

ratio

[nm]

[m2 g-1][nm][m2 g-1]

3/7

0.34

2.0

0.29

0.9

5/5

0.34

14

0.29

0.7

7/3

0.35

178

0.32

43

9/1

0.37

456

0.37

413

Si/Zr

[-]

Sa,BET- BET surface area
The effect of microstructural changes from SZA structure to CSZ structure on
membrane performance was also determined by preparing SZA and CSZ membranes
with the different Si/Zr atomic ratios. SZA membranes were fabricated at 300 oC
125

Hiroshima University Doctoral Dissertation

while CSZ membranes were fabricated at 550 oC based on the procedures previously
described in this work.
Fig 4-11 (a) shows the kinetic diameter dependence of single gas permeance for
SZA membranes measured at 200 oC while Fig. 4-11 (b) shows that for CSZ
membranes measured at the same temperature. H2 permeance and selectivities overCH4 for the two classes of membranes are presented in Table 4-5. It is immediately
evident from Fig. 4-11 (a) that all SZA membranes are capable of molecular sieving
with H2/CH4 ideal selectivity ranging between 6.6 and 11 (Table 4-5) well above theKnudsen ratio of 2.8. At the same time, the H2 permeances showed similar range ofvalues between 4.4 to 12 x 10-7 mol m-2 s-1 Pa-1. Therefore, it is difficult to observe
any substantial difference in the performance of SZA membranes with different Si/Zr
ratios. This can be explained by the fact that a SZA9/1 membrane prepared with a WRof 60 and calcined at 300 oC presents silica-rich phase that possesses a strong
molecular sieving ability. On the other hand, a SZA3/7 membrane prepared at thesame conditions presents ACAC-rich networks that also serve as molecular sieving
agents within the quasi-dense networks.
He H2 CO2 CH4
N2

-7

10

-8

10

-9

Si/Zr
9/1
5/5
3/7

2

2

3

4

5

-1

10

Tcal= 300 CHe H2 CO2 CH4

-5

N2

10

-6

10

-7

10

-8

10

-9

CF4 SF6o

Tcal= 550 C
Si/Zr
9/1
7/3
5/5
3/7

-1

-6

(b)

o

10

2

10

CF4 SF6

Permeance [mol m s Pa ]

-5

-1

-1

Permeance [mol m s Pa ]

(a)

10

6

Kinetic diameter [Å]

2

3

4

5

6

Kinetic diameter [Å]Fig. 4-11 (a) Single gas permeance at 200 oC as a function of kinetic diameter of
SiO2-ZrO2-ACAC-derived membranes fabricated at 300 oC in N2 for different Si/Zr
atomic ratios. (b) Single gas permeance at 200 oC as a function of kinetic diameter of126

Hiroshima University Doctoral Dissertation

carbon-SiO2-ZrO2-derived membranes fabricated at 550 oC in N2 for different Si/Zr
atomic ratios.
After carbonization however, CSZ membranes show an entirely different trend.
The membrane performance depends on the Si/Zr ratio. A CSZ9/1 shows a very highH2 permeance of 16 x 10-7 mol m-2 s-1 Pa-1 together with a similarly high H2/CH4 ideal
selectivity of 148. The trend then follows that as Si/Zr ratio reduces, H2 permeanceand H2/CH4 ideal selectivity reduce as shown by the values in Table 4-5.
Table 4-5 Comparison of H2 permeance and H2/CH4 permeance ratio of SiO2-ZrO2ACAC and carbon-SiO2-ZrO2 membranes at 200 oC
SiO2-ZrO2-ACAC
Si/Zr atomic

Carbon-SiO2-ZrO2

H2 permeanceα (H2/CH4)a, b H2 permeance[10-7 mol m-2 s-1 Pa-1][-][10-7 mol m-2 s-1 Pa-1]

[-]

3/7

4.4

6.6

0.16

2.7

5/5

8.8

11

1.1

48

7/3

-

-

3.2

55

9/1

12

7.4

16

148

ratio
[-]

a

α (H2/CH4) a, b

α = ideal selectivity, b H2/CH4 Knudsen ratio = 2.8
Fig. 4-12 summarizes the microstructural transformation from a SiO 2-ZrO2-

ACAC to carbon-SiO2-ZrO2 based on the relative Si/Zr atomic ratios. This illustration
also visualizes the effect on membrane performance. At a low Si/Zr ratio, the medium
is zirconium and therefore was ACAC rich. In this case, the high population of ACAC
physically modifying the network size can result in a quasi-dense network providing
molecular sieving. After carbonization, the high population of zirconium and ACAC
give rise to a permanently dense network and at the same time resulting in defectformation during the pyrolysis process. These combine to result in low H2permselectivity.
At the other end of the spectrum, when Si/Zr ratio is high, the medium is silica
rich with some ACAC still present as shown in Fig 4-12. The molecular sieving
ability of silica rich medium combined with some ACAC ligands cause H2permselectivity. After carbonization, the resulting CSZ structure becomes more127

Hiroshima University Doctoral Dissertation

siloxane (-Si-O-Si-) rich as a result of more thermal condensation. Also, since thecontent of zirconium and ACAC are limited in this case, defect formation is also
minimal. These will greatly improve the H2-permselectivity of the resulting
membrane. An organosilica membrane has also been shown to improve H2perselectivity upon high temperature calcination up to 700 oC which was attributed torearrangements of the silicon-carbon bonds upon pyrolysis of the organic linking units[40].
Intermediate Si/Zr ratio

Low Si/Zr ratio

OH

O
O Si
O

Zr

O Zr

O

Zr

O

High Si/Zr ratio

Zr

O

Zr
O Zr

O

Zr

O
Zr O SiOO

Carbonization

Si

OH

O Si O Zr
O O Zr
O
Zr
Zr
O O
O O
O
O
Zr O Si
Zr
Zr
O
O
O
O Zr
Si
O

OH

O Si
O

Zr

Zr

O

O Zr
O
O OH
Si
O
O
Zr

Si

O Zr

Si
O
O Si OH

O
O

Si

Carbonization

HO O Si
O

Si

Zr

O Si

O

Si

O
O

O
O

Si
O Zr

O

Si

O

O
Si O Si

OH
O

Si

OH

O

Carbonization

O Si O Zr
O O Si
O
Zr
Si
O O
O O
O
O
Zr
O Si
Zr
Si
O
O
O
O Zr
O Si

O Si O SiO O Si
O
Si
Si
O O
O O
O
O
Si
O Si
Zr
Si
O
O
O
O Zr
Si
O

Fig. 4-12 Schematic illustration of the effect of Si/Zr atomic ratio on ligand-to-carbon
and microstructural transformations.
4.3.3 Carbonization-enhanced H2-permselectivity in a SiO2-ZrO2 composite
membrane
Silica-derived membranes have impressively high H2-permselectivity, that is, they
can achieve simultaneous high H2 permeance and selectivity of H2 over larger gases.
Higher selectivities can be achieved when silica-derived membranes are fabricated ateven higher temperatures between 500 and 700 oC. However, the membranes begin to
trade off H2 permeance as a result of the densification and thus disappearance of H2permeating pores [14, 21]. For example, the H2 permeance for a silica membrane
prepared at 400 oC was 20 x 10-7 mol m-2 s-1Pa-1 with H2/N2 selectivity of 64. When
prepared at 600 oC however, the permeance reduced by 75 % to 5 x 10-7 mol m-2 s-1Pa128

Hiroshima University Doctoral Dissertation1

while the H2/N2 selectivity increased to more than 135 [12]. Kanezashi et al also

demonstrated the densification effect of temperature on silica by evaluating theactivation energies of gas permeation through silica 550 oC- and 750 oC-derived
membranes. They observed almost double increase of activation energies of gases in
750 oC-derived membranes compared to 550 oC-derived [41]. These are consistent
with the observation made by Lin et al. [21] regarding the microstructural propertiesof a bulk silica gel. A silica powder prepared at 40 oC had a pore volume of 0.458 cm3
g-1 which reduced by 75 % to 0.114 cm3 g-1 at 400 oC. In comparison to this, a SiO2ZrO2 composite in the Si/Zr ratio 9/1 showed only a 21 % reduction in pore volume
from 0.21 to 0.165 cm3 g-1 when calcined through a similar range of temperatures(from 200 to 550 oC) [42].
Yoshida et al. [22] prepared a SiO2-ZrO2 membrane with a Si/Zr ratio of 9/1 at
570 oC. The single gas permeance as a function of kinetic diameter is displayed inFig. 4-13 (a).

At 3 x 10-7 mol m-2 s-1Pa-1, the H2 permeance is lower than acorresponding pure SiO2 membrane prepared at 600 oC as reported by de Vos andVerweij. This is because ZrO2 being crystalline is a naturally dense material and while
the addition to SiO2 to form a composite matrix reduces the densification rate atelevated temperatures, it reduces the H2 permeance at the same time.
Fig. 4-13 (a) and (b) compare the CSZ9/1 and unmodified SiO2-ZrO2 (9/1)
membranes. Fig. 4-13 (a) shows the comparison of single gas permeance at 200 oC as
a function of kinetic diameter. By comparison, a SZA9/1 membrane prepared at 300oC has an H2 permeance of 12 x 10-7 mol m-2 s-1Pa-1, while a CSZ9/1 membrane

prepared at 550 oC has an H2 permeance of 16 x 10-7 mol m-2 s-1Pa-1 corresponding toalmost no change in H2 permeance. This is an interesting observation. The H2
permeance for a pure SiO2-ZrO2 (9/1) membrane is 5 times lower at 3 x 10-7 mol m-2
s-1Pa-1. Fig. 4-13 (b) also shows that both CSZ9/1 and SiO2-ZrO2 (9/1) membranes
have similar distribution of small pores.

129

Hiroshima University Doctoral Dissertation

He

-5

10

-6

10

-7

10

-8

(b)

CF4 SF6N2

2-1

CO2 CH4

H2-1

Permeance [mol m s Pa ]10

Normalised permeance [-]

(a)

10

10

-9

C-SiO2-ZrO2
SiO2-ZrO2

-10

2

3

4

5

He
10

Kinetic diameter [Å]CO2

N2 CH4

0

10

-1

10

-2

C-SiO2-ZrO2
SiO2-ZrO2

10

6

H2

-3

------- Knudsen

2.5

3.0

3.5

4.0

Kinetic diameter [Å]Fig. 4-13 (a) Single gas permeance at 200 oC as a function of kinetic diameter ofcarbon-SiO2-ZrO2-derived membranes (this work) compared with pure unmodified
SiO2-ZrO2-derived membranes [22] prepared with a Si/Zr atomic ratio of 9/1. (b)Corresponding dimensionless permeance as a function of the kinetic diameternormalized by He permeance.
Fig. 4-14 shows the temperature dependence of single gas permeance for He, H2,
CO2, N2 and CH4. All gases show increased permeance as temperature increases
indicating activated diffusion mechanism. Table 4-6 presents the calculated apparent
activation energy when the data points in Fig. 4-14 are fitted to the theoretical lines as
expressed by Equation 2-1 [43]. The evaluated activation energies of He and H2 for a
CSZ9/1 membrane (8.3 and 7.1 kJ mol-1, respectively) reveal very similar values tothose evaluated for a pure SiO2-ZrO2 (9/1) membrane (8.1 and 7.0 kJ mol-1,
respectively). Therefore, it can be confirmed that both a CSZ9/1 and SiO2-ZrO2 (9/1)
membrane possess similar distribution of small pores. In spite of this similarity inpore distribution, a CSZ9/1 displayed enhanced H2-permselectivity compared to a
SiO2-ZrO2 (9/1) membrane- H2 permeance and H2/CH4 selectivity were higher for theCSZ9/1 membrane.

130

Hiroshima University Doctoral Dissertation
o

10

-6

10

-7

10

-8

10

-9

200

150

H2/CH450

100

CO2/CH4H2
He-2Permeance [mol m s Pa ]

 [-]
-1

150
100
50
0
-5
10

-1

Temperature [ C]

CO2
N2
CH4

2.0

2.4
2.8
-1
1000/T [K ]3.2

Fig. 4-14 Temperature dependence of gas permeance for carbon-SiO2-ZrO2
membranes prepared with a Si/Zr atomic ratio of 9/1.
Table 4-6 Comparison of He and H2 permeance, activation energies and H2/CH4
permeance ratio for pure SiO2-ZrO2 [22] and Carbon-SiO2-ZrO2 membranes (thiswork); pore volume comparison between a SiO2-ZrO2 [42] and Carbon-SiO2-ZrO2
gels prepared at a Si/Zr atomic ratio of 9/1 at 550 oC.
Permeance
-7

-2 -1

Activation energy-1-1[10 mol m s Pa ][kJ mol ]α at 200 oC

Pore

[-]

volumeMembrane type

He

H2

He

H2

H2/CH4[cc g-1]

SiO2-ZrO2a

4.5

3

8.1

7.0

83

0.165

Carbon-SiO2-ZrO2

14

16

8.3

7.1

148

0.236

a

Membrane data from ref. [21] and pore volume data from ref. [42]; α- ideal

selectivity
A comparison of the pore volumes of bulk CSZ9/1 and SiO2-ZrO2 (9/1) materials
thus explains the improved H2-permselectivity of CSZ9/1 membrane. The respective
pore volumes are indicated in Table 4-6. The CSZ9/1 with a pore volume of 0.236
cm3 g-1 shows a much higher pore volume than a pure SiO2-ZrO2 by 43 %. This canbe explained by the presence of carbon nanoparticles within the SiO2-ZrO2 structure.
131

Hiroshima University Doctoral Dissertation

As is widely reported, pyrolyzed carbon nanoparticles possess ultra-micropores andmicropores which connect to provide combined molecular sieving and improved
permeability and are attractive for application as carbon molecular sieve membranes[44, 45]. Fig. 4-15 illustrates the microstructural difference between a pure SiO2-ZrO2
composite and a CSZ composite. In the pure composite matrix, the likelihood of -ZrO-Zr- linkages increase at high temperatures to form a crystalline tetragonal phase by
sintering mechanism [46]. This tetragonal phase may be undetected in an XRD
diffractogram as a result of the overwhelming silica amorphous phase. The carbon
nanoparticles in the CSZ network prevents the densification of ZrO2 into its
crystalline phase and occupies the aggregation volume that would have been occupied
by aggregated ZrO2 nanoparticles. Thus, while providing increased permeation pathfor H2, larger gases are sieved out by the ultra-micropore-micropore connections.
Carbon-SiO2-ZrO2

Unmodified SiO2-ZrO2
CH4

H2

O Si
O O SiOSiO OOOSi O Si
Zr
O
O
O Zr
Zr
O

Si

H2
Si

Aggregationof ZrO2CH4O SiO O SiSiO OOOSi O Si
Zr
O
O
Zr
O
Zr
O
O
O

Carbonization
reduction of ZrO2
aggregation

Fig. 4-15 Schematic illustration of enhanced H2-permselectivity enabled by the
carbonization of a SiO2-ZrO2 microstructure

4.4 Conclusions
In this work an attempt was made to derive the best sol synthesis and chemostructural parameters necessary for the fabrication of an H2-permselective carbonSiO2-ZrO2 (CSZ) composite membranes by pyrolysis of a SiO2-ZrO2-ACAC (SZA)
composite at 550 oC in N2 atmosphere. The hydrolysis water/alkoxide molar ratio(WR) and relative Si/Zr atomic ratios were chosen as study parameters. The effects ofwater/alkoxide molar ratios of 4, 60 and 240 as well as the Si/Zr atomic ratios werestudied and optimized.
132

Hiroshima University Doctoral Dissertation

Characterization

of

CSZ

powders

preparedwithdifferentWRbythermogravimetry, XRD and N2 adsorption reveal that the most suitable WR tomaintain the balance between carbon nanoparticles retention, amorphousness andmicroporosity is water/alkoxide molar ratio of 60.
Furthermore, fixing the WR at 60 and characterizing films, powders andmembranes prepared using Si/Zr ratios of 9/1, 7/3, 5/5 and 3/7 reveal that a properdifferentiation in microstructure and membrane performance will be achieved when
SZA is pyrolyzed to CSZ. A CSZ9/1 membrane showed the best H2-permselectivity
with H2 permeance of 16 x 10-7 mol m-2 s-1 Pa-1 and H2/N2 and H2/CH4 ideal
selectivities of 75 and 148, respectively. Therefore, the best combination ofparameters is a WR of 60 applied to prepare a carbon-SiO2-ZrO2 composite
containing a relative Si/Zr of 9/1.
Lastly, the carbonization in a CSZ9/1 membrane improved the H2-permselectivity
over a pure unmodified SiO2-ZrO2 containing a similar Si/Zr ratio and similar
distribution of the small pores. While the CSZ9/1 membrane showed an H2
permeance of 16 x 10-7 mol m-2 s-1 Pa-1 the unmodified SiO2-ZrO2 membrane was 3 x
10-7 mol m-2 s-1 Pa-1. This was attributed to the increased pore volume afforded by the
carbon nanoparticles through ultra-micropores inherent in them. It is therefore
resolved that the densification of silica at elevated temperatures can be ameliorated by
a combination of ZrO2 composition and a carbonization of the resulting composite.

Cited literature
1.

N. Abas, A. Kalair, N. Khan, Review of fossil fuels and future energy
technologies, Futures, 69 (2015) 31–49

2.

Y. Chang, J. Lee, H. Yoon, Alternative projection of the world energy
consumption-in comparison with the 2010 international energy outlook,
Energ. Policy, 50 (2012) 154-160

3.

B. Johnston, M. C. Mayo, A. Khare, Hydrogen: the energy source for the 21st
century, Technovation, 25 (2005) 569–585

4.

G. Nicoletti, N. Arcuri, G. Nicoletti, R. Bruno, A technical and environmental
comparison between hydrogen and some fossil fuels, Energ. Convers.
Manage., 89 (2015) 205–213
133

Hiroshima University Doctoral Dissertation

5.

M. Ball, M. Wietschel, The future of hydrogen – opportunities and challenges,
Int. J. Hydrogen Energ., 34 (2009) 615 – 627

6.

R. W. Rousseau, Handbook of separation process technology, John Wiley andSons, New York, 1987

7.

N. W. Ockwig, T. M. Nenoff, Membranes for hydrogen separation, Chem.
Rev. 107 (2007) 4078-4110

8.

A. P. Simpson and A. E. Lutz, Exergy analysis of hydrogen production via
steam methane reforming, Int. J. Hydrogen Energ., 32 (2007) 4811-4820

9.

S. C. A. Kluiters, presented under the 5th Research Framework Programme ofEuropean Union, Intermediate report EU project MIGREYD NNE5-2001-670,
Petten, December 2004

10.

G. Q. Lu, J. C. Diniz da Costa, M. Duke, S. Giessler, R. Socolow, R. H.
Williams, T. Kreutz, Inorganic membranes for hydrogen production andpurification: a critical review and perspective, J. Colloid Interf. Sci. 314(2007) 589–603

11.

R. M. de Vos, H. Verweij, Improved performance of silica membranes for gas
separation, J. Membr. Sci., 143 (1998) 37-51

12.

R. M. de Vos, H. Verweij, High-selectivity, high-flux silica membranes forgas separation, Science. 279 (1998) 1710-1711

13.

Y. S. Lin, Microporous and dense inorganic membranes: current status andprospective, Sep. Purif. Technol. 25 (2001) 39–55

14.

S. Ponton, F. Dhainaut, H. Vergnes, D. Samelor, D. Sadowski, V. Rouessac,
H. Lecoq, T. Sauvage, B. Caussat, C. Vahlas, Investigation of thedensification mechanisms and corrosion resistance of amorphous silica films,
J. Non-Cryst. Solids., 515 (2019) 34–41

15.

A. P. Dral, J. E. ten Elshof, Organic groups influencing microporosity inorganosilicas, Micropor. Mesopor. Mater. 267 (2018) 267–273

16.

M. Kanezashi, Y. Yoneda, H. Nagasawa, T. Tsuru, K. Yamamoto, J. Ohshita,
Gas permeation properties for organosilica membranes with different Si/C
ratios and evaluation of microporous structures, AIChE J. 63 (2017) 44914498.

17.

N. K. Raman, C. J. Brinker, Organic “template” approach to molecular sieving
silica membranes, J. Membr. Sci., 105 (1995) 273-279
134

Hiroshima University Doctoral Dissertation

18.

Y. E. Lee, B. S. Kang, S. H. Hyun, C. H. Lee, Organic‐templating approach tosynthesis of nanoporous silica composite membranes (ii): MTES‐templating
and CO2 separation, Sep. Sci. Technol., 39 (2004) 3541–3557

19.

M. Kanezashi, M. Asaeda, Hydrogen permeation characteristics and stabilityof Ni-doped silica membranes in steam at high temperature, J. Membr. Sci.,
271 (2006) 86–93

20.

M. Takenaka, H. Nagasawa, T. Tsuru, M. Kanezashi, Hydrocarbon
permeation properties through microporous fluorine-doped organosilica
membranes with controlled pore sizes, J. Membr. Sci., 619 (2021) 118787

21.

G. P. Fotou, Y. S. Lin, S. E. Pratsinis, Hydrothermal stability of pure andmodified microporous silica membranes, J. Mater. Sci. 30 (1995) 2803-2808

22.

K. Yoshida, Y. Hirano, H. Fujii, T. Tsuru, M. Asaeda, Hydrothermal stability
and performance of silica-zirconia membranes for hydrogen separation inhydrothermal conditions, J. Chem. Eng. Jpn., 34 (2001) 523-530

23.

C. Guizard, A. Ayral, A. Julbe, Potentiality of organic solvents filtration withceramic membranes. A comparison with polymer membranes, Desalination
147 (2002) 275-280

24.

H.-S. Choi, C.-H. Ryu, G.-B. Hwang, Obtention of ZrO2-SiO2 hydrogenpermselective membrane by chemical vapor deposition method, Chem. Eng.
J., 232 (2013) 302-309

25.

S.-J. Ahn, A. Takagaki, T. Sugawara, R. Kikuchi, S. T. Oyama, Permeation
properties of silica-zirconia composite membranes supported on porousalumina substrates, J. Membr. Sci., 526 (2017) 409-416

26.

J. Livage, C. Sanchez, Sol-gel chemistry, J. Non-Cryst. Solids., 145 (1992) 1119

27.

J. Livage, M. Henry, C. Sanchez, Sol-gel chemistry of transition metal oxides,
Prog. Solid St. Chem., 18 (1988) 25-341

28.

D. A. Ward, E. I. Ko, Preparing catalytic materials by the sol-gel method, Ind.
Eng. Chem. Res. 34 (1995) 421-433

29.

G. I. Spijksma, C. Huiskes, N. E. Benes, H. Kruidhof, D. H. A. Blank, V. G.
Kessler, H. J. M. Bouwmeester, Microporous zirconia–titania composite
membranes derived from diethanolamine-modified precursors, Adv. Mater.,18 (2006) 2165–2168
135

Hiroshima University Doctoral Dissertation

30.

S. Lawal, M. Kanezashi, H. Nagasawa, T. Tsuru, Development of anacetylacetonate-modified

silica-zirconia

membraneapplicable

to

gas

seoaration, J. Membr. Sci. 599 (2020) 117844
31.

S. Lawal, L. Yu, H. Nagasawa, T. Tsuru, M. Kanezashi, A carbon-silicazirconia membrane with CO2 flow-switching behavior promising hightemperature H2/CO2 separation, J. Mater. Chem. A. 8 (2020) 23563

32.

D. Hoebbel, T. Reinert, H. Schmidt, E. Arpac, On the hydrolytic stability oforganic ligands in Al-, Ti- and Zr-alkoxide complexes, J. Sol-Gel Sci.
Technol., 10 (1997) 115-126

33.

W. Puthai, M. Kanezashi, H. Nagasawa, K. Wakamura, H. Ohnishi, T. Tsuru,
Effect of firing temperature on water permeability of SiO2-ZrO2 membranes
for nanofiltration, J. Membr. Sci., 497 (2016) 348-356

34.

D. Peter, T. S. Ertel, H. Bertagnolli, EXAFS study of zirconium alkoxides asprecursors in sol-gel process: II. The influence of the chemical modification, J.
Sol-Gel Sci. Technol., 5 (1995) 5-14

35.

S. K. Parida, S. Dash, S. Patel, B. K. Mishra, Adsorption of organic moleculeson silica surface, Adv. Colloid Interface Sci., 121 (2006) 77–110

36.

M. Kanezashi, R. Matsugasako, H. Tawarayama, H. Nagasawa, T. Tsuru, Poresize tuning of sol-gel-derived triethoxysilane (TRIES) membranes for gas
separation, J. Membr. Sci., 524 (2017) 64-72

37.

L. Yu, M. Kanezashi, H. Nagasawa, J. Ohshita, A. Naka, T. Tsuru, Fabricationand microstructure tuning of a pyrimidine-bridged organoalkoxysilane
membrane for CO2 separation, Ind. Eng. Chem. Res., 56 (2017) 1316-1326

38.

S. O. Lawal, H. Nagasawa, T. Tsuru, M. Kanezashi, Facile development of
microstructure-engineered, ligand-chelated SiO2-ZrO2 composite membranes
for molecular separations, Mol. Syst. Des. Eng., 6 (2021) 429-444

39.

Z.-G. Wu, Y.-X. Zhao, D.-S. Liu, The synthesis and characterization ofmesoporous silica–zirconia aerogels, Micropor. Mesopor. Mater. 68 (2004)
127–132

40.

X. Yu, H. Nagasawa, M. Kanezashi, T. Tsuru, Improved thermal andoxidative stability of bis(triethoxysilyl)ethane (BTESE)-derived membranes,
and their gas-permeation properties, J. Mater. Chem. A, 6 (2018) 23378-23387

41.

M. Kanezashi, T. Sasaki, H. Tawarayama, H. Nagasawa, T. Yoshioka, K. Ito,
T. Tsuru, Experimental and theoretical study on small gas permeation
136

Hiroshima University Doctoral Dissertation

properties through amorphous silica membranes fabricated at different
temperatures, J. Phys. Chem. C, 118 (2014) 20323-20331
42.

W. Puthai, M. Kanezashi, H. Nagasawa, T. Tsuru, SiO2-ZrO2 nanofiltration
membranes of different Si/Zr molar ratios: stability in hot water andacid/alkaline solutions, J. Membr. Sci. 524 (2017) 700-711.

43.

H. Nagasawa, T. Niimi, M. Kanezashi, T. Yoshioka, T. Tsuru, Modified gastranslation model for the prediction of gas permeation through microporous
organosilica membranes, AIChE J., 60 (2014) 4199-4210

44.

M. Kiyono, P. J. Williams, W. J. Koros, Effect of pyrolysis atmosphere on
separation performance of carbon molecular sieve membranes, J. Membr. Sci.,
359 (2010) 2-10

45.

X. Ning, W. J. Koros, Carbon molecular sieve membranes derived from
Matrimid® polyimide for nitrogen/methane separation, Carbon, 66 (2014) 511522

46.

Y.-S. Lin, C.-H. Chang, R. Gopalan, Improvement of thermal stability ofporous nanostrucutured ceramic membranes, Ind. Eng. Chem. Res. 33 (1994)
860-870137

Hiroshima University Doctoral Dissertation

Chapter 5CO2 flow-switching behaviour in a pyrolysis-derived carbon-silicazirconia membrane promises a versatile high-temperature H2/CO2
separation
5.1 Introduction
Separation processes have already made a significant impact on the world [1]. The
operating expenses of these processes, however, represent exorbitant investments ofcapital in industries such as petroleum, chemical, petrochemical, pharmaceutical,
pulp, mineral and others [2]. Currently, about 80% of the world’s energy demand issatisfied using fossil fuels [3]. These facts suggest that the world’s industrial energy
demand (and by extension fossil fuel consumption) derives from separation processes.
By 2035, it is estimated that 8 countries (China, the USA, India, Brazil, Japan, South
Korea, Canada, and Mexico) will demand as much as 5.24 x 1017 kJ (4.97 x 1017
BTU) of energy [4], and 37% of that is estimated to come from the industrial sector
alone of which about 79% is foreseen to be fossil fuel-derived by 2030. 5 Therefore,
two interwoven problems must be solved: 1) reliance on fossil fuels; and, 2) energy
requirement for industrial separation processes. Fortunately, new separation
techniques such as membrane separation have been developed over the years to tackle
these problems. The development of hydrogen separation membranes could be a twopronged solution to the problems highlighted above [6]. On the one hand, membrane
separation could be an energy-efficient way to separate H2 from different mixtures [2,7-10], while on the other hand the resultant H2 could provide a sustainable and‘cheap’ source of energy [11, 12].
Hydrogen production from the steam reforming of natural gas currently produceslevels of production efficiency that average between 65 and 75%, which is the highestof all non-renewable production sources [11]. Currently, steam reforming of natural
gas also has the lowest cost, uses existing infrastructure [13], requires no oxygen forprocessing, and operates at the lowest temperature of all methods [14]. Obviously,
steam reforming is the most widely used method for hydrogen production as it
accounts for 80-85% of total hydrogen production [15]. As shown below, steam
reforming of natural gas consists mainly of methane reforming and is basicallyaccomplished in two steps [6, 16]:138

Hiroshima University Doctoral Dissertation

CH4 + H2O(g) → CO + 3H2 (Steam methane reforming step)
CO + H2O(g) → CO2 + H2 (Water-gas shift step)
The overall reaction becomes
CH4 + 2H2O(g) → CO2 + 4H2
The final step in steam methane reforming (SMR) is separation of the effluent
stream, which is comprised mostly of H2, H2O and CO2, the composition of which
depends on factors such as the reformer operating temperature and pressure and thesteam-to-carbon ratio [6, 16]. Techniques utilized for separation of H2 from CO2 andsteam (for >99 % pure H2 stream) have traditionally consisted of venerable energy
intensive approaches such as the pressure swing adsorption (PSA) of H2 or CO2,
cryogenic distillation [2, 6, 16, 17], and the more recent process of energy
conservative membrane separation [6, 16].
The composition of the final effluent gas stream is dependent on variableconditions such as temperature, pressure, and the steam-to-carbon ratio, which meansthat the separation technique of choice must be versatile enough to deal with thesecomposition variations. This poses a challenge for the development of membrane
separation techniques, and researchers must develop robust and versatile materials
that simultaneously deliver the core targets of viable H2 separation. Integrated
gasification combined cycle applications (IGCC) also utilize the water-gas shift(WGS), and candidate materials must achieve levels of H2 recovery and H2/CO2
selectivity of 70-90% and 21-62, respectively, under high temperatures and pressures[18], although SMR/WGS targets in industrial sources of H2 production have more
stringent requirements for H2 stream purity at > 99.99% [6, 16, 19]. In addition tothese targets, we believe taking H2 separation membranes a step further would entail
versatility in order to deal with H2/CO2 mixtures of widely varying compositions.
Several H2 separation membranes that are applicable in a wide variety oftemperature ranges have been developed over the years. The developers of thesemembranes have aspired to meet the set targets, and beyond, with varying degrees ofsuccess, which is associated with the material make-up of the membranes. Various H2separating membrane materials and processes have been reviewed [2, 6, 20]. Overall,H2 separating membranes have yet to meet the desired potential, and as of recently,
only two industrial applications of H2 separation membranes are known to exist(recovery of H2 from off-gases in the ammonia industry and production of pure H2 inthe electronic industry) [20]. Thus, the principal driver of success in this quest is139

Hiroshima University Doctoral Dissertation

confined to creativity in materials research. An interesting group of materials thathave received much attention in recent times are flexible metal-organic frameworks(FL-MOFs) [21-28], specifically because of their ability to structurally respond toexternal stimuli such as pressure, temperature, light and electric fields [26]. Such‘dynamic’ responses are exploitable in separation processes. For example, afluorinated MOF derived from a self-assembly of 2,2’-bis(4-carboxyphenyl)
hexafluoro propane and zinc nitrate hexahydrate has shown high selectivity of CO2
over N2 because perfluorinated compounds exhibit electrostatic interactions with CO2[23]. According to Kitagawa and Uemura [21], a construction of ‘dynamic’ porous
coordination polymers is rational via the placement of weak ‘interaction devices’ suchas van der Waals force, π- π stacks and hydrogen bonds between stiff 1-D, 2-D and 3D building blocks. This strategy offers a viable design for new porous materials by
utilizing existing ceramic materials that have shown promise for H2 separation. Thus,the ‘interaction devices’ employed must selectively engage with the mixture to be
separated.
Composite SiO2-ZrO2 ceramic membranes have exhibited desirable levels ofhydrothermal stability for H2 separation, but expensive molecular sieving is required[29]. Hence, our group recently reported the successful chemical modification of acomposite SiO2-ZrO2 ceramic membrane capable of molecular sieving [30]. This
membrane in its ligand-modified form offered no promise for H2/CO2 separation. A
successful modification with a chelating ligand of acetylacetone, however, has opened
the possibility for in-situ carbonization (pyrolysis of the organic ligand). If thedeposited free carbon during pyrolysis approximates the anisotropic characteristics ofgraphite [31], then the resultant van der Waals forces between the π- π stacks ofcarbon layers could serve as ‘interaction devices’ and endow new ‘dynamic’characteristics into the SiO2-ZrO2 matrix. Van der Waals interactions could also
develop between the carbon particles and the metal centers due to shifts in theelectron densities. Moreover, a recent report of the carbonization of a molecular sieve
silica membrane has introduced interesting properties to ordinary molecular sieve
silica [32, 33].
Herein, we describe how SiO2-ZrO2-acetylacetonate was successfully pyrolyzed
to form a carbonized form of SiO2-ZrO2 (hereafter called C-SiO2-ZrO2). Fig. 5-1
presents the expected pyrolysis and network formation pathway. The in-situ
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carbonization was evaluated by characterization methods. C-SiO2-ZrO2 interaction
with CO2 at various temperatures was investigated via sorption experiments. Single
and binary gas mixture experiments at 50 and 300 oC were performed to determine thepermeation characteristics of H2 and CO2 as well as the separation performance of analumina-supported C-SiO2-ZrO2 membrane. Finally, we discuss the place of a CSiO2-ZrO2 membrane in the scheme of H2 separation in industrial H2 production
systems and postulate as to the possibility that a versatile H2 separation can beachieved. In the present study, we believe that we achieved a rational design for
selective dynamism in a porous ceramic membrane.

Si

O

O

O

O

O
O

Zr

O O
Zr
O
O
Si

O

Si
O
O

O

O

Zr

O

O
O

O

Si

O

O

Si

Si

O

O

Zr
O

Zr

O

Zr

O

O

O
O

O
Zr

Zr

O

O
Si

Zr

O

O

Zr
O

Si

Zr

O

Si

O

O

Zr

O

Zr

O

Zr

O

Si

O
O

O

O
O

Zr

O Zr

O

O
Zr

Fig.5-1 Illustration of expected pyrolysis and network formation pathway of SiO2ZrO2-acac to C-SiO2-ZrO2

5.2

Experimental

5.2.1 Preparation and pyrolysis of SiO2-ZrO2-acac powder and C-SiO2-ZrO2
membrane fabrication
The preparation of SiO2-ZrO2-acac (SZa4) sol was described in Chapter 2 section
2.2.1. SZa4 (SiO2-ZrO2-acectylacetonate) powders were subsequently prepared from a
SZa4 sol by slowly drying the sol under an atmosphere controlled at 40-50 oC.
Carbon-SiO2-ZrO2 (CSZ550) powders were then prepared from SZa4 powders via
calcination in a 3-zone furnace at 550 or 750 oC under a N2 stream at 600 ml min-1 for
20-25 min. For convenience the resultant pyrolyzed products will hereafter be referred
to as CSZ550 or CSZ750 depending on the calcination temperature.
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To prepare CSZ550 membranes, it was necessary to fabricate membrane supports.
Details of this fabrication process has been described in Chapter 2 section 2.2.1. The
active separation layer was prepared thus: 1 wt% SZa4 sol in ethanol was wipe-coatedonto the pre-prepared SiO2-ZrO2 intermediate layer followed by calcination under a
600 ml min-1 nitrogen gas flow at 550 oC for 20-25 min. in a 3-zone continuous flow
furnace. This process was repeated 6–8 times to obtain the final defect-free CSZ550
membrane ready for single-gas and binary-mixture permeation measurements. It
should be noted that after fabrication the CSZ550 membrane was kept in themembrane module under a He flow for 12 hours at 200 oC.
5.2.2 Characterization of materials
The pyrolysis route of the acetylacetonate ligand in SZa4 powders was analyzed
and monitored using thermogravimetry and mass spectroscopy (TGA-DTA-PIMS
410/S, Rigaku, Japan; DTG-60 Shimadzu Co., Japan) under nitrogen or He gas flows
of 80 ml min-1 with a heating rate of 10 oC min-1. In a similar manner, the combustion
test for free carbon under air was performed using thermogravimetry (DTG-60
Shimadzu Co., Japan). The presence and the chemical state of free carbon was
confirmed using Fourier Transform-Infrared spectroscopy (FT-IR, FTIR-4100,
JASCO, Japan) and solid-state 13C cross polarization/dipolar decoupling magic-angle
spinning nuclear magnetic resonance (CP/DD MAS-NMR), which was recorded using
a NMR spectrometer (Varian 600PS, Agilent Inc., U. S. A.). 13C-CP/DD MAS-NMR
measurement was carried out at a frequency of 150.87 Hz using both 13C-1H dipolar
coupling and decoupling modes. The spectra were acquired in 360 scans each with a
90o pulse of 3.2 µs and recycle delays of 70 s. Hexamethylbenzene was used as areference against which the observed peaks were positioned, and a 6 mm zirconia
rotor spun at a magic angle of 57.4o at 7 kHz contained the sample. The physicalevidence of the presence of carbon nanoparticles was confirmed by obtaining the
micrographs of the SZa4 and CSZ550 particles using transmission electron
microscopy. The samples to be resolved were prepared on ultra-high-resolution
carbon supports (STEM 100Cu Grids) by dropping approximately 10 µl of a 2 wt%
dispersion of the fine particles in butanol onto the grids. Prior to examination, theprepared grids were vacuum-dried at 50 oC for 48 hours. The crystal/amorphous
structure and lattice spacing values of SZa4 and CSZ550 powders were obtained using
X-ray diffraction spectroscopy (D2 PHASER X-Ray Diffractometer, Bruker,
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Germany) with Cu Kα as the radiation source at a wavelength of 1.54 Å. Furthermore,
the sorption behaviors of CO2 on both SZa4 and CSZ550 powders were analyzed with
CO2 adsorption/desorption experiments carried out at 25, 30 and 35 oC. Prior to thismeasurement, samples of adsorbed gases and vapors were evacuated: SZa4 at 100 oC
and CSZ550 at 300 oC for at least 6 hours. Temperature-programmed desorption of N2
and CO2 was carried out using a custom-built equipment setup fitted with a mass
spectrometer (Dymaxion Dycor, Ametek Process Instruments, U. S. A.). Lastly, thecross-section morphology of the composite membrane was examined by Field
Emission-Scanning Electron Microscopy (FE-SEM, Hitachi S-4800, Japan). Prior to
examination, chipped samples of the membrane were attached to sample holders via
carbon tape and vacuum-dried at 50 oC for 24 hours.
5.2.3 Gas permeation measurement
A schematic representation of the gas permeation measurement setup is shown in
Fig. 5-2. Prior to the gas permeation measurement, a CSZ550 membrane was fitted
into its module immediately after fabrication and placed inside the furnace in the gas
permeation measurement rig at 200 oC under a moderate helium flow of 100 ml min-1
for about 12 hours to remove the adsorbed vapor. Single-gas permeation tests of theCSZ550 membranes were made using high purity gases (H2, He, CO2, N2, CH4, CF4,SF6 in that order). Each gas was fed to the upstream of the closed-end membrane
module at 200-500 kPa of absolute pressure under temperatures ranging from 50 –300 oC. Permeate side pressure was kept at atmospheric pressure and the permeate gas
flow was measured using a bubble film flow meter (HORIBA-STEC, Japan).
Binary mixture separation performance was evaluated using the setup shown in
Figure 5-2 at 50 and 300 oC. In the measurement of binary gas mixture separation,
H2/CO2 (80/20, 50/50, 20/80) mixtures were fed to the upstream of the module while
maintaining the pressure by setting the retentate side needle valve at 200 and 500 kPa
with the downstream kept at atmospheric pressure. The compositions of the feed,
retentate and permeate streams were analyzed using a pre-calibrated gas
chromatograph (GC-14B, Shimadzu, Japan) equipped with a TCD detector (column:
Porapak N (GL Science, Japan)). To evaluate the transmembrane pressure-drop ofeach component, the log mean pressure difference (ΔPi,lm) was applied (Equation 5-1).
𝚫𝑷

−𝚫𝑷

(5-1)

𝚫𝑷𝒊,𝒍𝒎 = 𝐥𝐧(𝚫𝑷𝒊,𝒊𝒏 ⁄𝚫𝑷𝒊,𝒐𝒖𝒕 )
𝒊,𝒊𝒏

𝒊,𝒐𝒖𝒕

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In Equation 5-1, ΔPi,in and ΔPi,out represent the difference in partial pressures ofcomponent i between the retentate side and the permeate side at the inlet and outlet ofthe module, respectively.
4
2

5

6

1. Gas cylinder
2. Drier
3. Pressure controller
4. Pressure gauge
5. Mass flow controller
6. Stop valve
7. Membrane module

3

1

4
2

5

6

3

1

4

8. Membrane
9. Furnace
10. Temperature controller
11. Back pressure valve
12. Gas sampler
13. Gas chromatograph
14. Soap film meter

11

14

14

8

7

12
9
13

10

Fig. 5-2 Schematic flow diagram of the gas permeation testing setup (Test pressures:
Pu = 200-500 kPa, Pd = 100 kPa)
5.3 Results and discussion
5.3.1 Carbonization of SiO2-ZrO2-acac into C-SiO2-ZrO2 and identification of
the carbon state
In Fig. 5-3 (a), the decomposition products of Acac- (acetylacetonate) wereconfirmed via TG-MS (thermogravimetry-mass spectroscopy under an inert He
atmosphere. In this figure the final decomposition temperature was around 500-550oC, which corresponds to the profile for SiO2-ZrO2-acac thermal decomposition under

an inert N2 atmosphere, as presented in Fig. 5-3 (b), which shows the DTG-DTA
profiles for the thermal decompositions of SiO2-ZrO2-acac and C-SiO2-ZrO2 under N2
and air, respectively. The final decomposition temperature for SiO2-ZrO2-acac underN2 was approximately 500-550 oC with the DTA curve showing no exothermal peaks,
which indicated no combustion reactions and possible deposition of non-volatile
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decomposition products. Under air, C-SiO2-ZrO2 (prepared by pyrolyzing SiO2-ZrO2acac under N2 at 550 oC for 20 minutes, as discussed under the experimental section)
showed thermal decomposition until the temperature reached 600 oC with exothermal
peaks in the DTA curve indicating the combustion of deposited non-volatile carbon to
CO and/or CO2. The weight loss was calculated at 20%, which corresponded to 203
mg carbon/g of C-SiO2-ZrO2.

-9m/z= 2 (H2)m/z 27 (CO)10m/z=13 (CH)m/z= 44 (CO2)

-10

10

b)

-12

10
1.0
0.8
0.6
0.4

0.2
200

SiO2-ZrO2-acac in He

400

600

800

1000

100

1.0

0

0.8

-100
0.6
-200
0.4
C-SiO2-ZrO2 in air

DTA [V]-11

10

Normalised residual weight [-]Normalised residual weight [-]Intensity [A]

a)

-300

SiO2-ZrO2-acac in N2

0.2

-400
200 400 600 800 1000ooTemperature [ C]

Temperature [ C]

Fig. 5-3 Observed TG-MS profile of SiO2-ZrO2-acac under He, and (a) TGA-DTA
profiles of SiO2-ZrO2-acac under N2 and C-SiO2-ZrO2 under air (b) (ramping rate: 10oC min-1)
It is important to determine the state of carbon in the SiO2-ZrO2 matrix as this willaid in understanding the behavior of the composite material. Fig. 5-4 shows thechange in the CP-MAS-13C-NMR (cross polarization) spectra of the SiO2-ZrO2-acactransformation to C-SiO2-ZrO2 (pyrolyzed at 550 and 750 oC). For the spectrum of
SiO2-ZrO2-acac, the different peaks indicate the intensity of the different functionalgroups of the enol and keto forms of the acac-- chelate resulting from 13C-1H dipolar
coupling. Acetylacetone in its keto tautomer form exists with two carbonyl groups(C=O), and as such is not able to properly chelate with transition metals. However, in
its enol form, acetylacetone presents with one carbonyl group (C=O) and one enolgroup (C=C-OH) that more easily forms acetylacetonate, thus facilitating chelation.
145

Hiroshima University Doctoral Dissertation

The magnified chemical shifts of peaks 5 and 6 on the SiO2-ZrO2-acac spectra show
very small intensities indicating that the enol forms of acac- were adequately chelated
to SiO2-ZrO2 by comparison with the spectrum of pure acetylacetone [35]Furthermore, the spectra of 550 and 750 oC-pyrolyzed C-SiO2-ZrO2 showed broad
peaks in the aromatic chemical shift region (~130 ppm) with two spinning side bands.
The spinning side bands are peaks due to the chemical shift in anisotropy that is
associated with sp2 hybridized C species such as aromatic and carbonyl C [36, 37].
This indicates the transformation of acac- into aggregated sp2 hybridized graphitic
carbon. The DD-MAS-13C-NMR (dipolar decoupling mode) that resolved themagnetic environment of 13C by decoupling 13C-1H revealed very sharp peaks (Fig. 55). This is reasonable since free carbon has no bonded 1H. The 550 and 750 oCpyrolyzed C-SiO2-ZrO2 samples showed similar spectra, further indicating thatpyrolysis up to 750 oC results in little or no change in the chemical and physical statesof the pyrolysis product.



C=C(aromatic) spinning side bandso

C-SiO2-ZrO2 (750 C pyrolysed)

1: -CH3(keto)o

C-SiO2-ZrO2 (550 C pyrolysed)

2

3: -CH3(enol)

1

4: -CH2(keto)56

5: =CH-(enol)
6: -C=O(enol)

43

Unfired SiO2-ZrO2-acac

250

200

150

100

50

0

Chemical shift [ppm]Fig. 5-4 CP-MAS-13C-NMR spectra of fresh SiO2-ZrO2-acac and that pyrolyzed at
550 and 750 oC.

146

Hiroshima University Doctoral Dissertation

CSZ550CSZ750



250

200

150

 spinning side bands** C=C(aromatic)



100

50

0

Chemical shift [ppm]Fig. 5-5 DD-MAS-13C-NMR spectra of C-SiO2-ZrO2 prepared at 550 and 750 oC
under N2.
The physical evidence of the presence of free carbon was obtained bytransmission electron microscopy. Fig. 5-6 (a)-(d) compare the electron diffraction
patterns in the TEM micrographs of 5 nm-scale SiO2-ZrO2-acac (Fig. 5-6 (a) and (b))
and C-SiO2-ZrO2 (Fig. 5-6 (c) and (d)) that were used to detect the presence of carbon
nanoparticles. Fig. 5-6 (b) and (d) are high-contrast conversions of Fig. 5-6 (a) and(c). The images show the edges of the particles. The micrographs reveal dark patchesof ~1 nm or less dispersed in the C-SiO2-ZrO2 matrix. Similar TEM images have beenreported for fluorescent carbon nanoparticles (CNPs) [38]. The inset images of both
SiO2-ZrO2-acac and C-SiO2-ZrO2 show halo electron diffraction patterns indicatingthat the carbonization process had little effect on the amorphous structure of the SiO2ZrO2 matrix.

147

Hiroshima University Doctoral Dissertation

a

b

50 nm

c

50 nm

d

50 nm

50 nm

Fig. 5-6 TEM images of SiO2-ZrO2-acac: (a) original; (b) high-contrast and 550 oCderived C-SiO2-ZrO2: (c) original; (d) high-contrast. Insets: Electron Diffraction
Images
5.3.2 Unique permeation properties of CO2 in C-SiO2-ZrO2 membranes
5.3.2.1 Understanding the interaction between CO2 and C-SiO2-ZrO2.
Fig. 5-7 (a) and (b) illustrate the observed evolution of the CO2/C-SiO2-ZrO2
interaction. In Fig. 5-7 (a), the adsorption isotherms of SiO2-ZrO2-acac and C-SiO2ZrO2 powders at 25 and 35 oC are shown. Both samples show gentle Langmuir-like(almost Henry) type I adsorption isotherms in the pressure range observed. The
adsorption-desorption isotherms of SiO2-ZrO2-acac at both 25 and 35 oC follow
reversible paths indicating almost no special effect of CO2 adsorption on neither the
SiO2-ZrO2-acac structure nor the surface. In contrast, after carbonization of SiO2ZrO2-acac resulted in C-SiO2-ZrO2, the adsorption-desorption isotherms of CO2
showed significant hysteresis despite the smaller amount of adsorbed CO2 indicating a
structural conformation of the C-SiO2-ZrO2 to allow CO2 adsorption. The existence of
this hysteresis at a temperature above the critical temperature of CO2 (31 oC) rules out
the possibility of capillary condensation, and instead suggests this is the result of a148

Hiroshima University Doctoral Dissertation

structural transformation [22]. Culp et al. [22] proposed that when host-guest
interaction can energetically compensate for the shape transformation of the hostlattice, the host structure could conform to the shape of the guest molecules.

b) 3002.5 Open: desorptionoBlack: 25 C
SZa4
o
Blue:
35
C
2.0

-1

1.5
CSZ5501.0

SZa4
CSZ 550250
Qst [kJ mol ]2-1

Va, CO [ml (STP) g ]a) 3.0 Closed: adsorption

200
150
100

0.5

50

0.0
0

20

40

60

80

0

100

0

20

40

60

80

100

CO2 uptake [%]Absolute pressure [kPa]

Fig. 5-7 (a) Adsorption-desorption isotherms of CO2 onto SiO2-ZrO2-acac and CSiO2-ZrO2 powders measured at 25 and 35 oC; (b) Calculated isosteric heats ofadsorption of CO2 onto SiO2-ZrO2-acac and C-SiO2-ZrO2 powders using temperaturesof 25, 30 and 35 oC
Fig. 5-7 (b) presents the isosteric heat of the adsorption of CO2 onto SiO2-ZrO2acac and C-SiO2-ZrO2 calculated using the adsorption isotherms at 25, 30 and 35 oC
and the Clausius-Clapeyron equation. The isosteric heat of CO2 adsorption for CSiO2-ZrO2 shows an average value of ~170 kJ mol-1 compared with ~34 kJ mol-1 for
SiO2-ZrO2-acac. Lin et al [39] calculated the theoretical potential energy surface(PES) barrier required for CO2 chemisorption on a pristine graphite to be ~350 kJ
mol-1 (~84 kcal mol-1). With CO2 adsorption on C-SiO2-ZrO2 releasing a maximum of~290 kJ mol-1, this suggests that CO2 adsorption onto C-SiO2-ZrO2 can produce
enough energy to deform the lattice structure to allow adsorption. The drastic
decrease in the Qst of C-SiO2-ZrO2 indicates an energetic heterogeneity of the C-SiO2ZrO2 surface.
Fig. 5-8 shows the observed XRD patterns and lattice spacing values calculated
from the centers of the amorphous peaks of SiO2-ZrO2-acac, C-SiO2-ZrO2, and pure
SiO2-ZrO2 powders (with and without firing) measured at room temperature. The d149

Hiroshima University Doctoral Dissertation

spacing value was reduced when SiO2-ZrO2-acac (0.356 nm) was pyrolyzed to CSiO2-ZrO2 (0.295 nm) at 550 oC under N2. On the other hand, pure SiO2-ZrO2
displayed a d-spacing shift from 0.314 nm to 0.291 nm after calcination at 550 oC.
The resultant d-spacing of C-SiO2-ZrO2 was similar but somewhat higher in value
compared with that of pure SiO2-ZrO2 (0.291 nm). The similar d-spacing values of CSiO2-ZrO2 and pure SiO2-ZrO2 may suggest that free carbon layers are well integrated
into the SiO2-ZrO2 lattice and do not form separate structures. The spacing between
the successive layers of carbon can thus be safely assumed to be less than 0.295 nm.
In Fig. 5-9, a scheme illustrating the interaction between CO2 and C-SiO2-ZrO2 isshown. As established from characterization methods, free carbon is integrated ingraphitic form in the C-SiO2-ZrO2 lattice. Graphitic carbon is anisotropic so that thelayers of carbon are closely bound together in the vertical axis by weak van der Waals
forces produced by delocalized π-orbital electrons (π-π stacking) and sp2-hybridized
C-C σ-bonding in the horizontal direction [40]. These delocalized electrons can
induce dipole moments on approaching CO2 molecules specifically because of the
known vibrational modes of CO2 which make it susceptible to dipole induction. The
induced CO2 molecules become attracted and electronically adsorbed onto thegraphitic carbon layers thereby undergoing structural changes because the weak van
der Waals forces allow the lattice to conform to the functionality of CO2 guest
molecules [22].
d = 0.291 nmoSiO2-ZrO2 (550 C, N2)

Intensity [a. u.]

d = 0.314 nm
SiO2-ZrO2 (unfired)d = 0.295 nmod = 0.356 nm

C-SiO2-ZrO2 (550 C)
SiO2-ZrO2-acac(Unfired)

10

20

30

40
o

2 [ ]

150

50

60

Hiroshima University Doctoral Dissertation

Fig. 5-8 Observed XRD patterns and calculated d-spacing values of SiO2-ZrO2-acac,
C-SiO2-ZrO2 and pure SiO2-ZrO2 powders measured at room temperature.

Induced dipole-induction

Graphitic layer-dipole induction> 2.95 Å≤ 2.95 Å

-

+

-

+

+

CO2

-

+

-

+

C-C σ-bonding
π-delocalization

Weak Van der waal’s forces
C-C σ-bonding

Fig. 5-9 Scheme illustrating the mechanism of the interaction between CO2 and CSiO2-ZrO2
5.3.2.2 Pressure-induced transition of CO2 flow in C-SiO2-ZrO2 membranes
Fig. 5-10 (a) and (b) show the time courses for the single-gas permeance of He, H2
and CO2 at 50 and 300 oC, respectively, through an alumina-supported C-SiO2-ZrO2
membrane, as measured at an upstream pressure of 200 kPa. In both cases, a dynamic
permeation trend is observed for CO2 whereby the permeance reduces drastically withtime before reaching a steady state. This points to a strong permeation hindrance to
CO2 flow due to blocking by adsorbed immobilized CO2. At 50 oC, He and H2
recovered about 65 and 79% of their initial values, respectively, after the feeding ofCO2, which effectively blocked certain pores due to adsorption by the C-SiO2-ZrO2
pores. This shows that adsorbed CO2 has a blocking effect on the permeation paths of
He and H2 and that He, H2 and CO2 share same permeation paths. At 300 oC thevalues for He and H2 permeance are recovered at rates of 80 and 83%, respectively,
and it would be safe to assume that the recovery ratio increases as temperature
increases so that adsorption would be stronger at 50 oC than at 300 oC given thecontinued occurrence of adsorption.
151

Hiroshima University Doctoral Dissertation

-6

10

-7

o

10
10

-9

-10

-1

50 C

-8

10

b)
-1

10

H2

He H

-2

He

2

CO2
Pu= 200 kPa
0Permeance [mol m s Pa ]

-2

-1

-1

Permeance [mol m s Pa ]

a)

1

2

3

4

5

Time [h]

10

-6

10

-7

10

-8

10

-9

10

o

300 C
He

-10

He

H2

H2

CO2
Pu= 200 kPa
0

1

2

3

4

5

Time [h]Fig. 5-10 Time courses for single-gas (He, H2 and CO2) permeation at 50 (a) and 300oC (b)Fig. 5-11 (a) shows the single-gas permeance for a C-SiO2-ZrO2 membrane as afunction of the kinetic diameter of different gases (He (0.26 nm), H2 (0.289 nm), CO2(0.33 nm), N2 (0.364 nm), CH4 (0.38 nm), CF4 (0.48 nm), and SF6 (0.55 nm))measured at 300 oC under 200 and 500 kPa of upstream pressure. Also shown in Fig.
5-11 (a) are the calculated values for permeance according to the Knudsen mechanism
based on SF6, as indicated by the broken lines. The experimentally obtained values for
He and H2 permeance far exceeded those calculated via the Knudsen mechanism,
which suggests the existence of pores so small that only H2 and He could permeate
them, which indicates a bimodal pore distribution in C-SiO2-ZrO2 membranes. It
should be noted that pressures are presented in absolute values. For the measurement
made at 200 kPa of upstream pressure, the gas permeance followed a trend such thatpermeance was decreased with an increase in the kinetic diameter, which generallyindicates a molecular sieving property. CO2 is the exception, and deviates from thetrend with a permeance lower than other gases irrespective of the kinetic diameter. At
500 kPa of upstream pressure, however, gas permeance follows a trend whereby it
decreases with an increase in the kinetic diameter without exception as opposed to gas
permeance at 200 kPa. It was noted that values for gas permeance at 200 kPa
remained either approximately the same (He, H2, CH4, CF4, SF6) at 500 kPa or
somewhat less (N2), but CO2 permeance at 500 kPa (1.5 x 10-9 mol m-2 s-1 Pa-1) was
increased by almost one order compared with that at 200 kPa (1.7 x 10-10 mol m-2 s-1
Pa-1). To further investigate this CO2-specific trend, the pressure-dependence of CO2
permeance was investigated using a fresh C-SiO2-ZrO2 membrane calcined at 550 oC,
152

Hiroshima University Doctoral Dissertation

as shown in Fig. 5-11 (b). The measurement was made multiple times to ensure
steady-state values and to verify accuracy. The data presented in Fig. 5-11 (b) are data
recorded as soon as CO2 reached steady state at 200 kPa. At subsequent pressures, thesystem response was quick and steady state was reached in a short time. In the initialfeed of CO2 gas to the upstream of the membrane at 200 kPa and 300 oC, it was
required to wait approximately 2 hours for steady state to be reached before recording
the results. The results obtained reveal that CO2 permeance showed almost no
dependence on upstream pressures that reached approximately 350 kPa. At 400 kPa
and above, however, a positive slope in the permeance of CO2 was almost linear withthe slope of further increases in the pressure.

N2
-9

10

-10

0.2

-7 -7

CH4

CF4

SF6

CO2

0.3

0.4

0.5

Kinetic diameter [nm]0.6Open: Pu 500 kPa
H2Knudsen
200 kPa
500 kPa

HHe
2
H2

-8

10108

-9 -9

CO2

CON22

CH4

CO

N2
CF4

SF6

4
2
1010-10200500
02 0.3400
4 0.4
6 600
8 700
100.6
0.2 300
0.5
Time
[h]
Upstream
Kinetic pressure
diameter [kPa]
[nm]
-10

-1

-1

Closed: Pu 200 kPa

-6

10

-1

16
1010

12-8
1010

-8

10

10

-6 -6

-2

Open: Pu 500 kPa
Knudsen
200 kPa
500 kPa

H2

-1

He

-2

-7

10

c)

20
1010

Permeance
Permeance[mol ms Pa ]-2 -1 -1 -1
-10
-2 -1
Permeances Pa[10 mol[mol
m m
s Pa
] ]

Closed: Pu 200 kPa

Permeance [mol m s Pa ]

b)
c)a)

-6

10

-2

-1

-1

Permeance [mol m s Pa ]

a)

H2

-7

10

H2

-8

10

CO2

N2

-9

10

-10

10

0

2

4

6

8

10

Time [h]Fig. 5-11 (a) Single-gas permeance at 300 oC as a function of kinetic diameter at 200
and 500 kPa upstream pressures; (b) CO2 permeance as a function of upstreampressure measured at 300 oC; (c) time course for the single-gas permeance of H2, CO2
and N2 at 300 oC measured at upstream pressures of 200 and 500 kPa (closed
symbols: 200 kPa; open symbols: 500 kPa).
This is consistent with the kinetic diameter dependence of single gas permeancemeasurements made at 200 and 500 kPa. In some cases, this type of trend can beassociated with the development of defects in a membrane [41, 42]. In the works of
DeSisto et al [41] and Barma and Mandal [42], defects appeared on SiO2 membrane
layers as thick as 7 µm and 6 µm respectively which happened to cause viscous flow.
As shown in Fig. 5-11 (c), however, after CO2 permeation at 200 and 500 kPa, thepermeance of H2 recovered as much as 99% of its levels before CO2 flow (afterdesorption of adsorbed CO2 in vacuum at 350 oC). Furthermore, the values for N2
permeation at 200 and 500 kPa showed similar values following CO2 permeation,
153

Hiroshima University Doctoral Dissertation

which is a sign of the absence of any defects. This suggests that the integrity of themembrane matrix was retained and was unaffected by the flow of CO2. Therefore, it
can be safely assumed that CO2 only displayed a switching of flow regimes in the CSiO2-ZrO2 membrane with 400 kPa being the transition pressure. Neither
plasticization nor defects can be suggested in this case. It should be noted that since
the measurement temperature (300 oC) was well above the critical temperature of CO2(31 oC), this phenomenon should not be ascribed to the flow of capillary condensate.
We therefore suggest a pressure-induced multilayer diffusion mechanism. It has
already been established in this section that adsorbed and immobilized CO2 molecules
prevent further permeation of gas-phase CO2 molecules. This trend persists at low
pressures. When the upstream pressure increases and the concentration of CO2
exceeds a certain threshold where adsorbent-adsorbate interaction dominates, multilayer adsorption may occur and thus lead to a greater mobility of gas-phase CO2
molecules [43].
Table 5-1 lists the calculations of activation energy for the temperaturedependence of the gas permeation for different gases (H2, CO2, N2 and SF6), as
regressed against Equation 2-1 and plotted in Fig. 5-12. Based on these calculations,H2 exhibited an activated diffusion mechanism with activation energy of 5.96 kJ mol1

, while other gases showed Knudsen permeation with activation energies of less than

1 kJ mol-1. In Fig. 5-12, it is evident that the blocking effect of adsorbed CO2 greatly
contributes to a reduction in the apparent CO2 permeance beyond the expected
Knudsen value across the tested range of temperatures.
Table 5-1 Calculated activation energy for the permeation of different gaseous species
through a C-SiO2-ZrO2 membrane regressed against Equation 2-1.
Gas

H2

CO2

N2

SF6

Activation energy [kJ mol-1]

5.96

0.55

-0.88

0.88

154

Hiroshima University Doctoral DissertationoTemperature [ C]
-6

300

-7

He

100

200

50

10

-2

-1

-1

Permeance [mol m s Pa ]

10

-8

10

-9

10

H2
N2
SF6
CO2

-10

10

1.6

2.0

2.4

2.8

3.2

-1

1000/T [K ]Fig. 5-12 Temperature dependence of single-gas permeance and permeance ratios attemperatures ranging from 50 to 300 oC for a C-SiO2-ZrO2 membrane prepared at 550oC.
Establishing the flow transition of CO2 presents an interesting phenomenon

whereby a C-SiO2-ZrO2 membrane could switch the flow of CO2 based on a predesigned threshold pressure. It is rational to assume that the threshold pressure for thetransition of CO2 flow would depend on the average pore size of the C-SiO2-ZrO2
membrane. Fine-tuning the average pore size to pre-design a threshold switchingpressure, however, is a subject for future investigation. To establish the recyclabilityof CO2-flow switching and the reliability of a C-SiO2-ZrO2 membrane over time, 3
cycles of the sequential pressure dependence of CO2 permeance at 200 and 500 kPa
upstream pressures were carried out at 300 oC for 16 hours, as shown in Fig. 5-13.
The two regimes of CO2 flow before and after transition were perfectly repeatable
over the 3 cycles without a significant difference in permeance. Also, N2 permeation
was carried out to further stress the fact that the flow-switching behavior is specific to
CO2. The permeance of N2 remained the same at 200 and 500 kPa, which indicated nopressure dependence for N2 permeation through the C-SiO2-ZrO2 membrane. Bothexperiments further verified that no defects were formed during this switching of theCO2 flow.
155

CO2
Closed: 200 kPa
Open: 500 kPa

-9

-1

100

-2

-1

Permeance [10 mol m s Pa ]Hiroshima University Doctoral Dissertation

N2

10

1

0

4

8

12

16

20

Time [h]Fig. 5-13 Pressure-induced switching cycles of CO2 flow represented by the timecourse of CO2 and N2 permeance under sequential pressurization at 200 and 500 kPa.
5.3.3 Binary H2/CO2 mixture separation performance of a C-SiO2-ZrO2derived membrane
5.3.3.1 Effects of temperature, pressure, and CO2 feed mole fraction on the binary
mixture separation performance.
Fig. 5-14 (a) shows the values for the permeance of H2 and CO2 through a C-SiO2ZrO2 membrane as a function of the CO2 mole fraction in the feed (and by extension
CO2 upstream partial pressure) measured at 50 oC. At 50 oC and 200 kPa, H2 and CO2
permeance was decreased as the CO2 mole fraction in the feed increased. This trend
resembles that of silica membranes where the permeance of the adsorptive gas is
expected to slightly increase as its mole fraction in the feed decreases [44]. The same
trend was observed at 50 oC and 500 kPa. Fig. 5-14 (b) illustrates the separation of abinary gas mixture of H2 and CO2 at 300 oC. At 200 kPa, the permeance of H2
maintained a high value of 7 x 10-8 mol m-2 s-1 Pa-1 up to a CO2 mole faction of 0.5
after which the permeance was reduced drastically to 4 x 10-9 mol m-2 s-1 Pa-1 at a CO2mole faction of 0.8. At 500 kPa, however, only a slight decrease in H2 permeanceoccurred even up to a CO2 feed mole fraction of 0.8. This underscores the fact that at
high values of both temperature and pressure, less CO2 is adsorbed, and the pressureinduced multilayer diffusion of CO2 presents less of a hindrance to the permeation of
H2, and, thereby, high H2/CO2 selectivity is maintained across all levels of CO2
concentration.
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Hiroshima University Doctoral Dissertation

CO2 feed fraction [-]CO2 feed fraction [-]-1

Permeance [mol m s Pa ]

a)

0 0.2

-7

0.5

0.8 1.0

0 0.2

0.5

0.8 1.0

-2

-1

10

-8

10

-9

10

CO2

CO2

Pu= 200 kPa

-10

10

H2

H2

0

Pu= 500 kPa

50 100 150 200 250

0

200

400

600

CO2 partial pressure [kPa] CO2 partial pressure [kPa]

CO2 feed fraction [-]
-7

10

0 0.2

-2

-1

-1

Permeance [mol m s Pa ]

b)

-8

0.5

CO2 feed fraction [-]0 0.2

0.8 1.0

0.5

H2

H2

CO2

CO2

0.8 1.0

10

-9

10

Pu= 500 kPa

Pu= 200 kPa

-10

10

0

50 100 150 200 250

CO2partial pressure [kPa]

0

200

400

600

CO2partial pressure [kPa]

Fig. 5-14 Gas permeance as a function of CO2 feed mole fraction and feed partialpressure at 50 oC (a) and 300 oC (b). Closed symbols: 200 kPa total upstreampressure; Open symbols: 500 kPa total upstream pressure.
As mentioned earlier in this section, the permeance increase as the concentrationof CO2 in the feed decreased was similar to that observed for silica membranes in asurface diffusion mechanism. The increases in CO2 permeance shown in Fig. 5-14 (a)and (b), however, were more substantial. In addition, the permeance of the CO2 in abinary mixture through the membrane was greater than that observed during single
permeation at the same feed pressure. This can be explained by the fact that H2 has amuch higher flux than the slower permeating component CO2, which exerts a‘sweeping’ effect, and CO2 is ‘swept’ along in the permeation paths. The same
phenomenon was observed for a polypyrrolone membrane where the much slower
permeating CH4 molecules were ‘swept’ along by the much faster CO2 molecules
[45].
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Hiroshima University Doctoral Dissertation

Sweep

Sweep

H2 + CO2 mixture flow

Pure CO2 flow

Fig. 5-15 Schematic illustration of the H2 ‘sweeping’ effect on CO2 permeation.
Such a ‘sweeping’ effect can be explained by considering the co-existence of two
gas species with different molecular sizes and adsorptive tendencies moving through acommon permeation path. The molecules of the fast and less-adsorptive permeating
gas continually collide with the slow permeating and more-adsorptive gas molecules
and therefore displace them from the adsorbent surface thereby increasing the gas
phase flow of the adsorbate gas molecules. Fig. 5-15 features a schematic illustrationof the effect of H2 presence on the adsorption-flux balance of CO2 through C-SiO2ZrO2 membrane pores. In the case of a binary mixture of H2 and CO2 gases, faster H2molecules sweep slower CO2 molecules and therefore there is a considerable
contribution of CO2 to the overall bulk phase transport of the gas mixture. The
fraction of the bulk mass flux contributed by each component in the gas mixture canbe calculated based on a set of equations presented by Kamaruddin and Koros [45],which compares to the gas-phase flux in single-gas permeation systems. When theseequations were applied to single and binary H2/CO2 systems at 300 oC and 500 kPa,
the results revealed that the fraction of the bulk-phase flux contribution by CO2 ismuch higher in the permeation of a gas mixture than in that of pure CO2 as shown in
Fig. 5-16, which was predicted by Kamaruddin and Koros [45].

158

Fraction of bulk flux contribution [-]Hiroshima University Doctoral Dissertation

0.6
BinarySingle

0.5
0.4
0.3
0.2
0.1
0.0

0

100 200 300 400 500 600
CO2 feed pressure [kPa]

Fig. 5-16 Calculated fraction of the CO2 bulk flux contribution at 300 oC as a functionof the CO2 feed pressure in single and binary systems (at total Pu = 500 kPa).
5.3.4 Outlook on the practical application of a C-SiO2-ZrO2 membrane for
H2/CO2 separation.
There are reports of membranes that separate H2 and CO2 by utilizing strong CO2
adsorption to create reverse CO2/H2 permselectivity at low temperature [46]. In thisprocess, the more adsorptive CO2 molecules permeate at the expense of blocked nonadsorptive H2 molecules. In these membranes, as permeation temperature increases,
molecular sieving properties dominate and achieve a level of H2 permselectivity thatis suitable for high-temperature applications such as steam methane reforming/watergas shift reactions (SMR/WGS) and the integrated gasification combined cycle(IGCC). In the present work, however, the molecular sieving properties at 300 oC with200 kPa of upstream pressure seemed inadequate to achieve sufficient H2
permselectivity across all CO2 concentrations due to persistent CO2 blocking. This isa drawback for practical and versatile high-temperature applications. This drawback iseliminated with a high pressure of 500 kPa where high H2 permselectivity can beachieved across various CO2 concentrations, as Fig. 5-14 (b) shows. For practicalhigh-temperature, high-pressure H2/CO2 separation purposes where membraneversatility may be of paramount importance as a result of fluctuating CO2
concentrations, a C-SiO2-ZrO2 membrane could be useful. Fig. 5-17 compares the H2
permeance-H2/CO2 mixture selectivity trade-off of high-temperature separation
membranes (details in Table 5-2). Most of the high-temperature H2/CO2 separation
159

Hiroshima University Doctoral Dissertation

membranes include a range of versions from zeolitic to ceramic that achieve
separation by means of molecular sieving (size exclusion of CO2) of a 1:1 binary H2CO2 mixture. Thus, it would be a novel occurrence to achieve comparable separation(good H2 permeance-H2/CO2 selectivity trade-off) with a CO2 adsorptive-type
membrane across several H2-CO2 mixture ratios. Fig. 5-17 clearly shows that C-SiO2ZrO2 membranes can achieve a separation performance that is comparable to many
molecular sieving membranes at high temperatures even with most CO2-concentrated
feed stock. This ability is a result of the pressure-induced transition of CO2 flow in CSiO2-ZrO2 membranes (as previously established), which results in less hindrance tothe flow of H2 and promotes activated H2 permeation.

H2/CO2 permeance ratio [-]

3

10

(8)

2

10

(3)
(4)

1

(7)

(6)
(2)

10

o

200-300 Co>300 CoThis work (50 C)oThis work (300 C)

(9)

(1)

(5)

0

10



Expected direction of reverse selectivity(High temperature)

-1

10

-9

10

-8

10

-7

10

-6

10

-1

-1

10
-2

-5

H2 permeance [mol m s Pa ]Fig. 5-17 Trade-off plot of H2/CO2 mixture selectivity versus H2 permeance values fordifferent high-temperature separation membranes

160

Hiroshima University Doctoral Dissertation

Table 5-2. Comparison of H2 permeance-H2/CO2 mixture selectivity trade-off of hightemperature H2/CO2 separation membranes
H2 permeanceNo.

Membrane type[10-8 mol m-2s-1Pa1]

α (H2/CO2)[-]

Reference

200-300 oC
1

Zeolite (SAPO34)

2
3

ZIF-7
Zeolite (B-ZSM5)

7

23

[47]

4.5

13.6

[48]

0.6

33

[49]

4

SiC

0.89

49

[50]

5

Si

50

6.8

[51]

6

Pd-organosilica

5.8

33.5

[52]

>300 oC
7

SiO2-TiO2

24.2

13.6

[53]

8

MFI-type Zeolite

39.6

141

[54]

10

47

[49]

20-30

This work

9

Zeolite (B-ZSM5)300 oC, 500 kPa
C-SiO2-ZrO2

7

α- selectivity

5.4 Conclusions
In this work we set out to investigate the possible unique properties that
could evolve during the in-situ carbonization of acetylacetonate ligands in a
SiO2-ZrO2 network and the resultant contribution to H2/CO2 separation. The insitu carbonization of SiO2-ZrO2-acac to C-SiO2-ZrO2 was confirmed by TG161

Hiroshima University Doctoral Dissertation

MS, DTG-TGA, FT-IR, CP/DD-MAS-13C-NMR, and TEM. The carbon
content in C-SiO2-ZrO2 promoted the adsorption of CO2, which resulted in a
maximum heat of adsorption at 290 kJ mol-1 and provided energy sufficient toallow the carbon layers in C-SiO2-ZrO2 to conform to the CO2 molecules. Thisdevelopment revealed that CO2 could be trapped. The CO2 permeance of the CSiO2-ZrO2 membrane demonstrated a significant level of dependence on thefeed pressure, but little dependence on either temperature or the CO 2 feed
concentration under conditions of 300 oC and 500 kPa, which allowed for aninteresting separation performance when used in binary H2/CO2 systems. The
C-SiO2-ZrO2

membraneachievedaH2/CO2

separation

performancecomparable to molecular sieving membranes irrespective of the CO 2 feed
content. In this study, we fabricated a two-parameter CO2-switchable ceramic
membrane. The switching parameters were the feed side pressure and the use ofa fast permeating gas component.
Cited literature
1.

D. S. Sholl, R. P. Lively, Seven chemical separations to change the world,
Nature, 532 (2016) 435-437

2.

R. W. Rousseau, Handbook of separation process technology, John Wiley andSons, New York, 1987

3.

S. Adhikari, S. Fernando, Hydrogen membrane separation techniques, Ind.
Eng. Chem. Res., 45 (2006) 875-881

4.

Y. Chang, J. Lee, H. Yoon, Alternative projection of the world energy
consumption-in comparison with the 2010 international energy outlook,
Energ. Policy, 50 (2012) 154-160

5.

E. A. Abdelaziz, R. Saidur, S. Mekhilef, A review on energy saving strategies
in industrial sector, Renew. Sust. Energ. Rev., 15 (2015) 150-168

6.

N. W. Ockwig, T. M. Nenoff, Membranes for hydrogen separation, Chem.
Rev. 107 (2007) 4078-4110

7.

H. Li, Z. Song, X. Zhang, Y. Huang, S. Li, Y. Mao, H. J. Ploehn, Y. Bao, M.
Yu, Ultrathin, molecular-sieving graphene oxide membranes for selectivehydrogen separation, Science, 342 (2013) 95-98

8.

T. Ji, Y. Sun, Y. Liu, M. Li, F. Wei, L. Liu G. He, Y. Liu, Facile in situ
hydrothermal synthesis of layered zirconium phenylphosphonate molecular162

Hiroshima University Doctoral Dissertation

sieve membranes with optimized microstructure and superb H2/CO2
selectivity, Acs. Appl. Mater. Inter., 2020, DOI: 10.1021/acsami.0c02789
9.

M. Kanezashi, D. Fuchigami, T. Yoshioka, T. Tsuru, Control of Pd dispersion
in sol-gel-derived amorphous silica membranes for hydrogen separation athigh temperatures, J. Membr. Sci. 439 (2013) 78-86

10.

S-J. Ahn, A. Takagaki, T. Sugawara, R. Kikuchi, S. T. Oyama, Permeation
properties of silica-zirconia composite membranes supported on porousalumina substrates, J. Membr. Sci. 526 (2017) 409-416

11.

E. Shoko, B. McLellan, A. L. Dicks D. da Costa, Hydrogen from coal:production and utilization technologies, Int. J. Coal Geo., 65 (2006) 213-222

12.

P. P. Edwards, V. L. Kuznetsov, W. I. F. David, N. P. Brandon, Hydrogen andfuel cells: towards a sustainable energy future, Energ. Policy, 36 (2008) 43564362

13.

C. Acar and I. Dincer, Comparative assessment of hydrogen production
methods from renewable and non-renewable sources, Int. J. Hydrogen Energ.,
39 (2014) 1-12

14.

C. Acar and I. Dincer, I., Review and evaluation of hydrogen production
methods for better sustainability Int. J. Hydrogen Energ., 40 (2015) 1109411111

15.

D. R. Simbeck, Hydrogen costs with CO2 capture. Presented at the
7thinternational conference on greenhouse gas control technologies (GHGT7), Vancouver, British Columbia, Canada, September 6–10; 2004

16.

A. P. Simpson and A. E. Lutz, Exergy analysis of hydrogen production via
steam methane reforming, Int. J. Hydrogen Energ., 32 (2007) 4811-4820

17.

L. Barelli, G. Bidini, F. Gallorini and S. Servili, Hydrogen production throughsorption-enhanced steam methane reforming and membrane technology: areview, Energ., 33 (2008) 554-570

18.

A. K. Ku, P. Kulkarni, R. Shisler and W. Wei, Membrane performancerequirements for carbon dioxide capture using hydrogen-selective membranes
in integrated gasification combined cycle (IGCC) power plants, J. Membr.
Sci., 367 (2011) 233-239

19.

J. J. Stanislowski, M. J. Holmes, A. C. Snyder, S. G. Tolbert and T. J. Curran,
Advanced CO2 separation technologies: coal gasification, warm-gas cleanup,and hydrogen separation membranes, Energ. Procedia, 37 (2013) 2316-2326
163

Hiroshima University Doctoral Dissertation

20.

S. C. A. Kluiters, presented under the 5th Research Framework Programme ofEuropean Union, Intermediate report EU project MIGREYD NNE5-2001-670,
Petten, December 2004

21.

S. Kitagawa and K. Uemura, Dynamic porous properties of coordination
polymers inspired by hydrogen bonds, Chem. Soc. Rev., 34 (2005) 109-119

22.

J. T. Culp, M. R. Smith, E. Bittner and B. Bockrath, Hysteresis in thephysisorption of CO2 and N2 in a flexible pillared layer nickel cyanide, J. Am.
Chem. Soc., 130 (2008) 12427-12434

23.

C. A. Fernandez, P. K. Thallapally, R. K. Motkuri, S. K. Nune, J. C. Sumrak,
J. Tian and J. Liu, Gas-induced expansion and contraction of a fluorinatedmeta-organic framework, Crys. Growth Des., 10 (2010) 1037-1039

24.

A. Boutin, F.-X. Coudert, M.-A. Springuel-Huet, A. V. Neimark, G. Férey and
A. H. Fuchs, The behavior of flexible MIL-53(Al) upon CH4 and CO2
adsorption, J. Phys. Chem. C, 114 (2010) 22237-22244

25.

J. van den Bergh, C. Gücüyener, E. A. Pidko, E. J. M. Hensen, J. Gascon and
F. Kapteijn, Understanding the anomalous alkane selectivity of ZIF-7 in theseparation of light alkane/alkene mixtures, Chem. Eur. J., 17 (2011) 88328840

26.

J. Wang, J. Luo, J. Zhao, D.-S. Li, G. Li, Q. Huo and Y. Liu, Assembly of two
flexible metal-organic frameworks with stepwise gas adsorption and highly
selective CO2 adsorption, Crys. Growth Des., 14 (2014) 2375-2380

27.

Z.-J. Li, J. Lu, M. Hong and R. Cao, Metal-organic frameworks based on
flexible ligands (FL-MOFs): structures and applications, Chem. Soc. Rev., 43(2014) 5867-5895

28.

G. A. Craig, P. Larpent, S. Kusaka, R. Matsuda, S. Kitagawa and S.
Furukawa, Switchable gate-opening effect in metal-organic polyhedral
assemblies through solution processing, Chem. Sci., 9 (2018) 6463-6469

29.

K. Yoshida, Y. Hirano, H. Fujii, T. Tsuru, M. Asaeda, Hydrothermal stability
and performance of silica-zirconia for hydrogen separation in hydrothermal
conditions, J. Chem. Eng. Jpn. 34 (2001) 523-530.

30.

S. Lawal, M. Kanezashi, H. Nagasawa, T. Tsuru, Development of anacetylacetonate-modified

silica-zirconia

membraneseoaration, J. Membr. Sci. 599 (2020) 117844
164

applicable

to

gas

Hiroshima University Doctoral Dissertation

31.

R. R. Dagastine, D. C. Prieve and L. R. White, Calculations of van der Waals
forces in 2-dimensionally anisotropic materials and its application to carbon
black, J. Colloid Interf. Sci., 249 (2002) 78-83

32.

M. C. Duke, J. C. Diniz da Costa, D. D. Do, P. G. Gray, G. Q. Lu,
Hydrothermally robust molecular sieve silica for wet gas separation, Adv.
Funct. Mater. 16 (2006) 1215-1220

33.

M. C. Duke, J. C. Diniz da Costa, G. Q. Lu, M. Petch, P. G. Gray, Carbonizedtemplate molecular sieve silica membranes in fuel processing systems:
permeation, hydrostability and regeneration, J. Membr. Sci. 241 (2004) 325333

34.

W. Puthai, M. Kanezashi, H. Nagasawa, K. Wakamura, H. Ohnishi, T. Tsuru,
Effect of firing temperature on water permeability of SiO2-ZrO2 membranes
for nanofiltration, J. Membr. Sci. 497 (2016) 348-356.

35.

X. Tao, W. Qiu, H. Li, T. Zhao, Synthesis of nanosized zirconium carbide
from preceramic polymers by the facile one-pot reaction, Polym. Adv.
Technol., 21 (2010) 300-304

36.

K. Ikeya and A. Watanabe, Application of 13C ramp CPMAS NMR withphase-adjusted spinning sidebands (PASS) for the quantitative estimation ofcarbon functional groups in natural organic matter, Anal. Bioanal. Chem., 408(2016) 651-655

37.

J. C. C. Freitas, F. G. Emmerich, G. R. C. Cernicchiaro, L. C. Sampaio and T.
J. Bonagamba, Magnetic susceptibility effects on 13C MAS NMR spectra ofcarbon materials and graphite, Solid State Nucl. Mag., 20 (2001) 61-73

38.

D. Pan, J. Zhang, Z. Li, C. Wu, X. Yan and M. Wu, Observation pf pH-,solvent-, spin-, and excitation-dependent blue photoluminescence from carbon
nanoparticles, Chem. Commun., 46 (2010) 3681-3683

39.

S. C. Xu, S. Irle, D. G. Musaev and M. C. Lin, Quantum chemical predictionof reaction pathways and rate constants for dissociative adsorption of COx andNOx on the graphite (0001) surface, J. Phys. Chem. B, 110 (2006) 2113521144

40.

D. D. L. Chung, Graphite, J. Mat. Sci., 37 (2002) 1475 – 1489

41.

B. A. McCool, N. Hill, J. DiCarlo and W. J. DeSisto, Synthesis andcharacterization of mesoporous silica membranes via dip-coating andhydrothermal deposition techniques, J. Membr. Sci., 218 (2003) 55–67
165

Hiroshima University Doctoral Dissertation

42.

S. Barma, and B. Mandal, Synthesis and characterization of orderedmesoporous silica membrane: role of porous support and gas permeation
study, Micr. Meso. Mater., 210 (2015) 10-19

43.

J.-G. Choi, D. D Do and H. D. Do, Surface diffusion of adsorbed molecules inporous media: monolayer, multilayer, and capillary condensation regimes, Ind.
Eng. Chem. Res., 40 (2001) 4005-4031

44.

M. Kanezashi, W. N. Shazwani, T. Yoshioka and T. Tsuru, Separation of
propylene/propane binary mixtures by bis(triethoxylsilyl) methane (BTESM)derived silica membranes fabricated at different calcination temperatures, J.
Membr. Sci., 415-416 (2012) 478-485

45.

H. D. Kamaruddin, and W. J. Koros, Some observations about the applicationof Fick’s first law for membrane separation of multicomponent mixtures, J.
Membr. Sci., 135 (1997) 147-159

46.

M. Hong, S. Li, J. L. Falconer and R. D. Noble, Modification of zeolite
membranes for H2 separation by catalytic cracking of methyldiethoxysilane, J.
Membr. Sc., 307 (2008) 277–283

47.

M. Yu, H. Funke, R. D. Noble and J. L. Falconer, H2 separation using defectfree, inorganic composite membranes, J. Am. Chem. Soc., 133 (2011) 17481750

48.

A. Huang, Y. Chen, Q, Liu, N. Wang, J. Jiang and J. Caro, Synthesis of highlyhydrophobic and permselective metal-organic-framework Zn(BDC)(TED)0.5
membranes for H2/CO2 separation, J. Membr. Sci., 454 (2014) 126-132

49.

M. Hong, J. L. Falconer and R. D. Noble, Modification of zeolite membranesfor H2 separation by catalytic cracking of methyldiethoxysilane, Ind. Eng.
Chem. Res., 44 (2005) 4035-4041

50.

B. Elyassi, M. Sahimi and T. T. Tsotsis, Silicon carbide membranes for gas
separation applications, J. Membr. Sci., 288 (2007) 290-297

51.

R. M. de Vos and H. Verweij, High-selectivity, high-flux silica membranes forgas separation, Science, 279 (1998) 1710-1712

52.

J. Lei, H. Song, Y. Wei, S. Zhao and H. Qi, A novel strategy to enhancehydrothermal stability of Pd-doped organosilica membrane for hydrogen
separation, Micr. Meso. Mater., 253 (2017) 55-63

166

Hiroshima University Doctoral Dissertation

53.

Y. Gu and S. T. Oyama, Permeation properties and hydrothermal stabiltity ofsilica-titania membranes supported on porous alumina substrates, J. Membr.
Sci., 345 (2009) 267-275

54.

Z. Tang, J. Dong and T. M. Nenoff, Internal surface modification of MFI-typezeolite membranes for high selectivity and high flux of hydrogen, Langmuir,
25 (2009) 4848-4852

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Chapter 6
Conclusions and considerations for further studies
6.1 Summary and conclusions
To summarize, the microporous structure and properties of a SiO2-ZrO2 composite
has been tailored in 3 major ways:
1. The chemical modification of the transition metal component, Zr, by chelatingwith acetylacetone ligand. This ligand performed two functions of arresting thehydrolysis rate of the Zr precursor and modifying the network of the resulting
SiO2-ZrO2 composite.
2. By utilizing unsaturated Si precursor (3-(trimethoxysilyl)propyl methacrylateMAPTMS) in combination with acetylacetone (ACAC) or allyl acetoacetate(AAA), it was possible to crosslink the unsaturated moieties via a thermal radical
polymerization.
3. The carbonization of the SiO2-ZrO2 composite by pyrolyzing acetylacetonatemodified SiO2-ZrO2. The effect of acetylacetone concentration and Si/Zr ratio
were investigated.
Thus, the major conclusions derived from these studies can be summarized into 4
items:(i)

Acetylacetonate-modified

SiO2-ZrO2-derived

membrane

showedH2permeance of 9.9 x 10 -7 with a H2/SF6 permeance ratio of 7,600, which was better
than the results when using a pure SiO2-ZrO2-derived membrane (H2 permeance: 1.4
x 10 -6 mol m-2 s-1 Pa-1, H2/SF6 permeance ratio: 11). When an acac--modified SiO2ZrO2-derived membrane prepared at 250 oC (H2 permeance: 1.8 x 10-7 mol m-2 s-1 Pa-1,
H2/CH4: 100, CO2/CH4: 60 at 50 oC) was heat-treated at 550 oC, the result was
improved H2 permeance of 5.2 x 10-6 mol m-2 s-1 Pa-1 with low permeance ratios(H2/CH4: 3, H2/SF6: 9) at 50 oC. The 250 oC-fired acac--modified SiO2-ZrO2-derived
membranes exhibited a gas permeation mechanism similar to polymer membranes,
and the 300 oC-heat-treated version showed the best trade-off between H2 permeanceand the H2/SF6 permeance ratio.

168

Hiroshima University Doctoral Dissertation(ii)

SiO2-ZrO2-based organically crosslinked hybrid membranes were fabricated

via the sol-gel crosslinking of ligand (AAA or ACAC)-chelated zirconium oxide
networks and MAPTMS derived siloxane networks. The AAA-chelated systemgenerated a copolymerization crosslinking with the Si precursor while the ACACchelated system generated only a homopolymerization within MAPTMS. New pore
structures were formed as organic crosslinked materials showed a significant increase
in measured BET surface areas (OCL-AAA: ≈320 m2 g-1, OCL-ACAC: ≈40 m2 g-1)
compared with non-crosslinked samples (< 1 m2 g-1). The organic crosslink-derived
membranes showed better molecular sieving properties with higher H2/N2 selectivity(OCL-AAA: 33, OCL-ACAC: 20) compared with non-crosslink-derived membranes(H2/N2 selectivity; NCL-AAA: 11, NCL-ACAC: 7). Finally, organic crosslinkderived membranes showed a change in microstructural properties from a flexible to arigid structure that ranged between that of silica-rich materials and polymer materials.
Considering that a facile method was established to fabricate and tune themicrostructure of supported class II hybrid inorganic-organic highly permeable
membranes, this study should serve as a template for future work where tailored
network tuning and functionalization can be achieved for specific applications.

(iii)

An attempt was made to derive the best sol synthesis and chemo-structuralparameters necessary for the fabrication of carbon-SiO2-ZrO2 (CSZ) composite
membranes by pyrolysis of a SiO2-ZrO2-ACAC (SZA) composite at 550 oC in N2
atmosphere. The hydrolysis water/alkoxide molar ratio (WR) and relative Si/Zr
atomic ratios were chosen as study parameters. The effects of water/alkoxide molar
ratios of 4, 60 and 240 as well as the Si/Zr atomic ratios were studied and optimized.
The most suitable WR to maintain the balance between carbon nanoparticles
retention, amorphousness and microporosity is water/alkoxide molar ratio of 60. Alsousing Si/Zr ratios of 9/1, 7/3, 5/5 and 3/7, CSZ9/1 membrane showed the best H2permselectivity with H2 permeance of 16 x 10-7 mol m-2 s-1 Pa-1 and H2/N2 andH2/CH4 ideal selectivities of 75 and 148, respectively. Therefore, the best combinationof parameters is a WR of 60 applied to prepare a carbon-SiO2-ZrO2 composite
containing a relative Si/Zr of 9/1which showed improved H2-permselectivity over apure unmodified SiO2-ZrO2 containing a similar Si/Zr ratio and similar distribution of
the small pores. This was attributed to the increased pore volume afforded by the169

Hiroshima University Doctoral Dissertation

carbon nanoparticles through ultra-micropores inherent in them. It is therefore
resolved that the densification of silica at elevated temperatures can be ameliorated by
a combination of ZrO2 composition and a carbonization of the resulting composite.

(iv)

A carbon-SiO2-ZrO2 (CSZ) with Si/Zr ratio of 5/5 was further investigated.

The carbon content in C-SiO2-ZrO2 promoted the adsorption of CO2, which resulted ina maximum heat of adsorption at 290 kJ mol-1 and provided energy sufficient to allow
the carbon layers in C-SiO2-ZrO2 to conform to the CO2 molecules. This development
revealed that CO2 could be trapped. The CO2 permeance of the C-SiO2-ZrO2
membrane demonstrated a significant level of dependence on the feed pressure, butlittle dependence on either temperature or the CO2 feed concentration underconditions of 300 oC and 500 kPa, which allowed for an interesting separation
performance when used in binary H2/CO2 systems. The C-SiO2-ZrO2 membrane
achieved a H2/CO2 separation performance comparable to molecular sieving
membranes irrespective of the CO2 feed content.

6.2 Considerations for further studies
So far, the studies undertaken during this research period, as pointed out earlier,have focused on ligand modification of ceramic composite, microstructural
engineering of such modified composites by choosing specific network forming
ligands and carbonization of the composite by pyrolysis. Based on the results and theconclusions drawn therefrom, the following considerations are put forward as future
study subjects.
1. Ligand-ceramic composites from network-forming and functional moieties
In chapter 3, allyl acetoacetate as a network forming ligand was used inconjunction with an unsaturated Si precursor for the microstructural engineering of
the ligand modified SiO2-ZrO2 composite. This microstructural tuning improved themolecular sieving and predictability of the network pore distribution. For futurestudies, the selection of a functional group bearing Si precursor is suggested inaddition to the network forming Si precursor (MAPTMS) so that a combination of
network engineering and functional interaction with permeating species could be
explored. Suggested precursors and their potential functional behavior include
170

Hiroshima University Doctoral Dissertation(i) Phenyltriethoxysilane- PTEOS (hydrophobicity)(ii) Aminopropyltrimethoxysilane- APTMOS (hydrophilicity or CO2 affinity)(iii) Tridekafluoro-tetrahydrooctyltriethoxysilane-TDFTHOTEOS (oleophobicity)
2. Formation of carbon-ceramic composites from thermosetting organic modifiers
The carbonization of SiO2-ZrO2 composites and the application to the formationof gas separation membranes were explored in chapters 4 and 5. One major challengewith carbonization by pyrolysis was the difficulty in preventing defect formation
during pyrolysis. The thermal stability of the ligands employed were low allowing thematerial to lose a lot of mass at the final carbonization temperatures. Therefore, it issuggested that thermally stable and thermosetting organic ligands be employed for themodification of ceramic composites. Based on chapter 5, by retaining a very highamount of residual carbon in the carbon-ceramic composite, high-pressure CO2/N2
and CO2/CH4 separations could be possible.
In conclusion, ligand-ceramic combinations offer a lot of opportunities for thedevelopment of new membrane types that can be applied to a variety of needs.

171

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172

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List of publications
Journal articles
1. Sulaiman Lawal, Masakoto Kanezashi, Hiroki Nagasawa, Toshinori Tsuru,
Development of an acetylacetonate-modified silica-zirconia composite membrane
applicable to gas separation, J. Membr. Sci., 599 (2020) 117844 (p.1-11)
2. Sulaiman Lawal, Liang Yu, Hiroki Nagasawa, Toshinori Tsuru, Masakoto
Kanezashi, A carbon-silica-zirconia ceramic membrane with CO2 flow switching
behavior promising versatile high-temperature H2/CO2 separation, J. Mater.
Chem. A, 8 (2020) 23563-23573
3. Sulaiman Lawal, Hiroki Nagasawa, Toshinori Tsuru, Masakoto Kanezashi,
Facile development of microstructure-engineered, ligand-chelated SiO2-ZrO2
composite membranes for molecular separations, Mol. Syst. Des. Eng., 6 (2021)
429-444
4. Sulaiman Lawal, Hiroki Nagasawa, Toshinori Tsuru, Masakoto Kanezashi,
Enhancement of the H2-permselectivity of a silica-zirconia composite membrane
enabled by ligand-ceramic to carbon-ceramic transformation, J. Membr. Sci.,
Under review, (July 2021)
Conferences
1. Sulaiman Lawal, Masakoto Kanezashi, Hiroki Nagasawa, Toshinori Tsuru,
Effect of zirconium-acetylacetonate chelation on the molecular sieving propertiesof a SiO2-ZrO2 composite membrane, The 16th International Conference on
Inorganic membranes (ICIM16), Taipei, Taiwan (Notification of acceptance on
2/27/2020), 2020/06/29–07/02
2. Sulaiman Lawal, Liang Yu, Hiroki Nagasawa, Toshinori Tsuru, Masakoto
Kanezashi, Fabrication of a carbon-functionalized silica-zirconia ceramic
membrane with pressure-induced CO2 flow switching, The Membrane Society ofJapan Membrane Symposium 2020 (Online Zoom platform), 2020/11/12-13

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174

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Acknowledgements
AlhamduliLlah! I must first and foremost thank Allah, the Almighty, the Wise.
Despite my many shortcomings in His worship and commands, He (Allah) has seenme through 3 years of rigorous studies and has held me up at trying times. There is noamount of gratitude I can express that comes close to appreciating His Mercies on me.
I must also thank my parents (Alhaji Abdurrauf and Mrs. Kifayat Lawal) who
have shown me great understanding, love and support throughout my life. Between us
they know when the goings were tough, and they were there to encourage andmotivate me to keep pushing for my life goals. I am very glad that they are both alive
to see me complete a major milestone in my life and see me accomplish the goals they
have dreamt for me in their daily struggles. I cannot acknowledge them enough and I
love them both for the sake of Allah.
Next, to my teachers, many of whom I have had since coming to Japan. I mustfirst thank immensely Professor Masakoto Kanezashi who showed belief by
accepting me to undertake a doctorate under his supervision despite having noconcrete research background in the field of membrane science and technology. I
would also like to appreciate the very kind, jovial and astute guidance of our head,
Professor Toshinori Tsuru. He always has a positive outlook towards everything,
and this quality is what drives the Separation and Technology laboratory at Hiroshima
University. Professor Hiroki Nagasawa has also been very helpful for my daily life.
He is down-to-earth and always ready to talk and be of help. I definitely could nothave asked for any better research group to work with.
It would be very criminal to forget to mention the family of students in the
Separation and Technology laboratory at Hiroshima University both past and present.
I would like to appreciate the help of Dr. Liang Yu (now back in China at the Beijing
Institute of Technology) who was always ready to discuss the progress of my work.
He always created time to listen to me and give very enlightening advise and insights.
He certainly is one of the smartest people I have met. I would also like to appreciate
Dr. Sofiatun Anisah who was the person that taught me my first experimental works
and showed patience towards me. The Japanese students have been of immense help
to me throughout my stay in this laboratory and for my daily life. I would like to175

Hiroshima University Doctoral Dissertation

particularly mention Norihiro Moriyama, Naoya Hataoka (my student guide), Mari
Takenaka (she helped me so much sometimes I am ashamed of asking for morehelp), Kenta Nakahiro, Kenta Watanabe, Takahiro Nagaoka, Mitsugu Kawasaki
and others that have shown me friendship over the years.
I also want to appreciate the Japanese government through the Ministry of
Education, Culture, Sports, Science and Technology (MEXT) for providing mewith the scholarship to study and live in Japan for three years and a half. I also wantto mention the kind help of the Japanese Embassy and Ambassador in Nigeria forproviding the required guidance during the scholarship application processes.
Last but definitely not the least, I want to appreciate and acknowledge my darlingwife Basirat Mojisola Ayinde for being my companion through this journey. You
have certainly made living in Japan a whole lot easier for me, and I would not haveachieved many things if not for your understanding, strength and love. I have many athing to apologize for, but I hope that you would overlook my faults. You are verystrong, dedicated and caring. You gave birth to our son Hudhayfah with no family toassist us, you bathed him and fed him everyday for the past 18 months while also
pursuing your doctorate degree. If that is not strength, I do not know what is. Thankyou very much. JazaakiLlahu khayr! I love you very much.
Thank you all.

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