Monash University Synthesis of Zeolite Nanocrystals and Their Application for Mixed Matrix Membranes by Dan Li July 2009 A dissertation submitted for the degree of Doctor of Philosophy in the Department of Chemical Engineering at Monash University Supervisor: Dr Huanting Wang Department of Chemical Engineering
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Monash University
Synthesis of Zeolite Nanocrystals and
Their Application for Mixed Matrix
Membranes by
Dan Li
July 2009
A dissertation submitted for the degree of Doctor of Philosophy in the Department of
Chemical Engineering at Monash University
Supervisor: Dr Huanting Wang
Department of Chemical Engineering
Declaration
i
Declaration
I hereby declare that this thesis contains no material which has been accepted for the
award of any other degree or diploma at any university or equivalent institution and that,
or the best of my knowledge and belief, this thesis contains no materials previously
published or written by another person, excepted where due reference is made in the
text of the thesis.
Signed: _________________
Date: ___________________
Acknowledgements
ii
Acknowledgements I would like to take this opportunity to thank my supervisor Dr Huanting Wang for his
willingness to accept me into his research team and his unwavering support throughout
my PhD study. His insight and suggestions have helped me progress so far.
I extend my appreciation to all of the people working in our lab, including Dr Xinyi
Zhang, Dr Jianfeng Yao, Dr Wei Zhu, Dr Jingping Wei, Ms Zhanli Chai, and Dr Dehua
Dong for their nice help during my experiments. I have enjoyed so much in working
with all of them.
Special thanks go to Dr Yunxia Yang and Ms Na Hao for their assistance on gas
adsorption-desorption test and 29Si MAS NMR analysis. Also thanks Yuan Fang, Nicky
Eshtiaghi, and Hue-chen Au Yong for their encouragement and kind care.
Many thanks go to the Australian research council discovery for funding this research
project.
Lastly, I would like to thank my family members, especially my parents and my
boyfriend Weihan Wang, for all of their love, support and encouragement in all that I
set out to achieve.
在我博士毕业在即，我希望借此机会感谢我的父母对我一直以来的教育和关怀。
回首过往的生活和学习中，每当我遇到艰辛和困惑的时候，他们总是给予我无限
的支持，信任和鼓励。没有他们多年来对我付出的心血，也不会有今天的我。亲
爱的爸爸妈妈，在我的心中，我永远爱你们。
Thank you all.
Dan Li
July 2009
Table of Contents
iii
Table of Contents
Declaration.................................................................................................. i
Acknowledgements.................................................................................... ii
Table of Contents ..................................................................................... iii
Abstract..................................................................................................... vi
List of Publications................................................................................... ix
List of Figures........................................................................................... xi
List of Tables ........................................................................................... xv
List of Schemes ....................................................................................... xvi
Abbreviations........................................................................................ xviii
Chapter 1 Introduction............................................................................. 1
Chapter 2 Literature Review.................................................................... 3
2.1 Overview.............................................................................................. 3
2.2 Introduction to zeolite.......................................................................... 3
2.3 Application of zeolite nanocrystals ..................................................... 7
2.4 Hydrothermal synthesis of zeolite nanocrystals .................................. 9
2.5 Use of polymers in zeolite synthesis ................................................. 17
2.5.1 Confined-space synthesis of zeolite nanocrystals..........................................17
2.5.2 Effect of added polymers on zeolite structure ...............................................21
2.5.3 Chitosan hydrogels.........................................................................................24
2.6 Application of mixed matrix membranes (MMMs) to hydrogen
separation .......................................................................................... 27
2.6.1 Hydrogen separation ......................................................................................27
2.6.2 Polymer membranes for hydrogen separation ...............................................28
2.6.3 MMMs for hydrogen separation ....................................................................33
Table of Contents
iv
2.6.4 Gas transport through membranes .................................................................39
2.7 Organic functionalization of zeolite nanocrystals and membrane
fabrication.......................................................................................... 47
2.8 Summary and Aims ........................................................................... 50
Chapter 3 Growth of Zeolite in Chitosan Hydrogels............................ 52
3.1 Overview............................................................................................ 52
3.2 Zeolite crystallization in crosslinked chitosan hydrogels.................. 52
3.2.1 Experimental ..................................................................................................52
3.2.1.1 Synthesis of zeolite LTA (NaA)................................................................52
3.2.1.2 Synthesis of zeolite FAU (NaY)...............................................................53
3.2.1.3 Removal of crosslinked chitosan hydrogels ............................................55
3.2.1.4 Characterization .....................................................................................55
3.2.2 Results and Discussion...................................................................................56
3.2.2.1 Effect of the amount of SiO2 ....................................................................57
3.2.2.2 Effect of the amount of chitosan..............................................................60
3.2.2.3 Effect of the amount of glutaraldehyde (GA) ..........................................61
3.2.2.4 Effect of aging time .................................................................................63
3.2.2.5 Effect of heating time ..............................................................................66
3.2.2.6 Comparison between the treatment of H2O2 and conventional
calcination .............................................................................................68
3.2.2.7 Synthesis of FAU nanocrystals................................................................69
3.2.3 Summary ........................................................................................................71
3.3 Formation of cubic zeolite A with an amorphous core in
uncrosslinked chitosan hydrogels ..................................................... 72
3.3.1 Experimental ..................................................................................................72
3.3.1.1 Synthesis of cubes of zeolite A with an amorphous core.........................72
3.3.1.2 Characterization .....................................................................................72
3.3.2 Results and Discussion...................................................................................73
3.3.3 Summary ........................................................................................................79
3.4 Comparisons of zeolite formation mechanisms ................................ 80
3.5 Conclusions........................................................................................ 82
Table of Contents
v
Chapter 4 Organic-functionalized Sodalite Nanocrystals.................... 84
4.1 Overview............................................................................................ 84
4.2 Experimental...................................................................................... 84
4.2.1 Synthesis of organic-functionalized silicalite nanocrystals ...........................84
4.2.2 Synthesis of organic-functionalized sodalite nanocrystals ............................85
4.2.3 Characterization .............................................................................................85
4.3 Results and Discussion ...................................................................... 86
4.3.1 Transformation of silicalite............................................................................86
4.3.2 Evidence of organic functionalization of sodalite..........................................89
4.3.3 Gas adsorption and pore structures ................................................................91
4.3.4 Surface modification: dispersion in solvents .................................................92
4.4 Conclusion ......................................................................................... 93
Chapter 5 Preparation & Characterization of Mixed Matrix
Membranes .............................................................................. 95
5.1 Overview............................................................................................ 95
5.2 Experimental...................................................................................... 95
5.2.1 Membrane fabrication ....................................................................................95
5.2.2 Characterization .............................................................................................97
5.3 Results and Discussion ...................................................................... 99
5.3.1 Membrane characterization............................................................................99
5.3.2 H2 sorption of sodalite nanocrystals and gas permeation of membranes.....105
5.4 Conclusion ....................................................................................... 111
Chapter 6 Conclusions & Recommendations for Future Work ........ 112
6.1 Conclusions...................................................................................... 112
6.2 Recommendations for future work .................................................. 113
References .............................................................................................. 116
Appendix-Relevant Publications.......................................................... 137
Abstract
vi
Abstract Zeolites are a class of microporous solids with well-defined crystalline structures. Due
to their ability to distinguish molecules on the basis of size and shape, zeolites are often
referred to as molecular sieves. There has been considerable interest in the synthesis of
zeolite nanocrystals (or nanozeolites), because they can serve as a model system for
fundamental understanding of zeolite nucleation and growth mechanisms, as seeds for
secondary growth of zeolite films and membranes, as building blocks for construction
of hierarchical porous nanostructures, and for preparation of mixed matrix membranes
(MMMs).
Presently, in the synthesis of zeolite crystals, it is well-known that the addition of
organic additives or polymers has an effect on the zeolite nucleation and crystallization.
Despite several synthetic strategies of using some polymers (e.g. polyacrylamide and
methylcellulose) have been developed to grow zeolites, there has been no research on
the synthesis with crosslinked chitosan hydrogels or uncrosslinked chitosan polymers,
which is one primary goal of this thesis. Therefore, chitosan (crosslinked or
uncrosslinked polymers) was introduced into zeolite crystallization process. The results
showed that glutaraldehyde-crosslinked chitosan (GA-CS) hydrogels were effective
three-dimensional network structures for controlling the growth of zeolite NaA and
NaY. The zeolite crystal sizes were significantly affected by formulation of silica-
containing GA-CS hydrogels and alkaline solution, and by aging and heating
conditions. Importantly, a novel method of using hydrogen peroxide solution was
developed to remove GA-CS hydrogels after zeolite synthesis, which was considered as
an effective way for removal of GA-CS hydrogels. The zeolite NaA nanocrystals with
an average size of 148 nm and NaY with an average size of 192 nm were synthesized in
this research. In addition, the resultant zeolite NaA and NaY nanocrystals were readily
redispersed in deionized water and some other solvents, and therefore they may be
useful for some applications, e.g. in the fabrication of zeolite-polymer mixed matrix
membranes (MMMs) and hierarchical porous zeolitic structures. In this thesis, the effect
of uncrosslinked chitosan hydrogels on zeolite nucleation and crystallization was also
studied. Cubic zeolite A with a single crystalline shell and an amorphous core was
prepared for the first time by in-situ crystallization of sodium aluminosilicate gel inside
Abstract
vii
the chitosan polymer networks. The TEM characterization further revealed that this
formation process of cube-like or rectangular core-shell structures involved particle
aggregation and surface-to-core crystallization induced by chitosan networks. It is
expected that this work would provide a new model system for understanding and
studying complex zeolite nucleation and growth mechanisms.
To date, the research into efficient separation of hydrogen has been driven by its
potential as an essential component of future energy economies. Despite some materials
emerging for this purpose, it is believed that there should be plenty of room to develop
mixed matrix membranes (MMMs) with an addition of inorganic particles, such as
zeolites, for hydrogen separation or purification. In order to reduce the phase separation
between organic and inorganic phases, organic functionalization is suggested as an
effective way, which has been applied in my study. Sodalite, whose framework consists
of a six-membered ring aperture with a pore size of 2.8 Å, was selected as inorganic
fillers in MMMs. Because of its small pore size and high ion-exchange capacity,
sodalite has been considered as a good candidate material for a wide range of
applications such as optical materials, waste management, hydrogen storage, and
hydrogen separation. To functionalize sodalite nanocrystals, organic functional groups
(i.e.,–(CH3)(CH2)3NH2 and –CH3) were successfully attached to sodalite nanocrystals
by the newly developed method – the direct transformation of organic-functionalized
silicalite nanocrystals. Gas sorption results showed that the organic-functionalized
sodalite nanocrystals contained uniform pore channels that were accessible to hydrogen
molecules at 77 K, but inaccessible to nitrogen, as expected. In addition, the dispersion
of organic-functionalized sodalite nanocrystals in organic solvents was favoured by the
presence of organic functional groups.
Organic-functionalized sodalite nanocrystals with –(CH3)(CH2)3NH2 functional groups
(denoted Sod-N) were incorporated into polyimide membranes to form sodalite-
polyimide mixed matrix membranes (MMMs) for hydrogen separation. Characterization
by SEM showed that Sod-N can be well distributed with polyimide phase, even when
their loadings reached 35 wt% in the hybrid, as confirmed by the FTIR spectroscopy
and XRD results. TG results revealed the temperatures for corresponding major mass
loss increased with the increasing Sod-N content of MMMs. This was attributed to the
interaction between the amino moieties from inorganic nanoparticles (Sod-N) and
Abstract
viii
polymer matrix, which restricted the movement of the main chains. The gas permeation
results exhibited the significantly improved hydrogen separation property. It was found
that H2 permeability was improved; and the MMMs with a loading of 35 wt% Sod-N
had the highest selectivity ( )( 22 NHα = 277) and a good permeability (8.04 Barrer) at
room temperature.
List of Publications
ix
List of Publications Publications related to thesis:
Journal Papers:
1) D. Li, H. Y. Zhu, K. R. Ratinac, S. P. Ringer, H. T. Wang, Synthesis and
Characterization of Sodalite–polyimide Nanocomposite Membranes,
Microporous and Mesoporous Materials 2009, doi:10.1016/j.micromeso.
2009.05.014.
2) J. Yao, D. Li, X. Zhang, C-H Kong, W. Yue, W. Zhou, H. T. Wang,
Cubes of Zeolite A with an Amorphous Core, Angewandte Chemie,
International Edition 2008, 120, 8525-8527.
3) D. Li, Y. Huang, K. R. Ratinac and S. P. Ringer, H. T. Wang, Zeolite
Crystallization in Crosslinked Chitosan Hydrogels: Crystal Size Control
and Chitosan Removal, Microporous and Mesoporous Materials
2008, 116, 416-423.
4) D. Li, J. F. Yao, H. T. Wang, N. Hao, D. Y. Zhao, K. R. Ratinac, S. P.
Ringer, Organic-functionalized Sodalite Nanocrystals and Their
Dispersion in Solvents, Microporous and Mesoporous Materials
2007,106, 262-267.
Refereed conference papers:
1) D. Li, H. T. Wang, Fabrication of Sodalite-polymer Nanocomposite
Membranes, Chemeca 2008, Newcastle, proceedings of Chemeca 2008,
993-1000.
2) D. Li, J. F. Yao, H. T. Wang, N. Hao, D. Zhao, K. R. Ratinac, S. P.
Ringer, Synthesis and Organic-functionalization Sodalite Nanocrystals,
Chemeca 2007, Melbourne, proceedings of Chemeca 2007, 117-123.
Oral Presentations:
1) D. Li, H. T. Wang, Fabrication and Characterization of Sodalite-Polymer
Nanocomposite Membranes, AIChE 2008, Philadelphia.
List of Publications
x
2) D. Li, H. T. Wang, Zeolite Crystallization in Crosslinked Chitosan
Hydrogels: Crystal Size Control and Chitosan Removal, AIChE 2008,
Philadelphia.
3) D. Li, H. T. Wang, Fabrication of Sodalite-polymer Nanocomposite
Membranes, Chemeca 2008, Newcastle.
4) D. Li, H. T. Wang, Fabrication and Characterization of Sodalite-
polyimide Nanocomposite Membranes, IMSTEC 2007, Sydney.
5) D. Li, J. F. Yao, H. T. Wang, N. Hao, D. Zhao, K. R. Ratinac, S. P.
Ringer, Synthesis and Organic-functionalization Sodalite Nanocrystals,
Chemeca 2007, Melbourne.
Other journal publications:
1) X. Zhang, D. Li, L. Bourgeois, H. T. Wang, P. A. Webley, Direct
Electrodeposition of Porous Gold Nanowire Arrays for Biosensing
Applications, Chemphyschem 2009, 10, 436-441.
2) X. Zhang, D. Dong, D. Li, T. Williams, H. T. Wang, P. A. Webley,
Direct Electrodeposition of Pt Nanotube Arrays and Their Enhanced
Electrocatalytic Activities, Electrochemistry Communications 2009, 11,
190-193.
3) H. Li, J. Yao, D. Li, J. Ho, X. Zhang, C.-H. Kong, Z.-M. Zong, X.-Y.
Wei, H.T. Wang, Hollow Zeolite Structures Formed by Crystallization in
Crosslinked Polyacrylamide Hydrogels, Journal of Materials Chemistry
2008, 18, 3337-3341.
4) J.-H. Hong, D. Li, H. T. Wang, Weak-base Anion Exchange Membranes
by Amination of Chlorinated Polypropylene with Polyethyleneimine at
Low Temperatures, Journal of Membrane Science 2008, 318, 441-444.
List of Figures
xi
List of Figures Figure 2.1. Framework structures of zeolites: (a) representations of [SiO4]4- or
[AlO4]5- and some possible connections; (b) structures of SOD,
FAU, LTA, MFI and*BEA [27, 28]............................................................. 4
Figure 2.2. Sodalite: (a) unit cell containing β-cage (b) framework [85]. ...................... 17
Figure 2.3. Chemical structures of chitin polymer (a) and chitosan polymer
(b) [104]...................................................................................................... 25
Figure 2.4. Upper bound correlation for H2/N2 separation (Prior upper bound
was firstly published in 1991 and the present upper bound was
updated in 2008.) [149]. ............................................................................. 30
Figure 3.1. The reaction between hydroxyl radical from hydrogen peroxide
and carbohydrates [116-118]...................................................................... 57
Figure 3.2. XRD patterns of the samples prepared with molar compositions of
0.01CS:ySiO2:12.5GA:21HAc:(1166+8y)H2O (y = 10-20) under
the same aging (36 h) and heating (90 °C for 3h) conditions: (a)
A-10SiO2; (b) A-12.5SiO2; (c) A-15SiO2; (d) A-17.5SiO2 and
(e) A-20SiO2. All of the samples were collected after H2O2
treatment..................................................................................................... 59
Figure 3.3. (a) SEM image, (b) particle size distribution determined by SEM,
(c) particle size distribution measured by light scattering, and (d)
N2 sorption isotherm of the sample A-17.5SiO2. ....................................... 59
Figure 3.4. XRD patterns of the samples prepared with molar compositions of
xCS:17.5SiO2:1250xGA:21HAc:(1233+6944x)H2O (x = 0.005-
0.0125) under the same aging (36 h) and heating (90 °C for 3 h)
conditions: (a) A-0.0125CS; (b) A-0.01CS; (c) A-0.0075CS; (d)
A-0.005CS. All of the samples were collected after H2O2
treatment..................................................................................................... 60
Figure 3.5. SEM images of the particles produced with different amounts of
added GA: (a) A-0.3GA; (b) A-1.0GA; (c) A-2.0GA; (d) A-
4.0GA. ........................................................................................................ 61
List of Figures
xii
Figure 3.6. Particle size distributions determined by SEM images for A-
0.3GA and A-1.0GA. ................................................................................. 62
Figure 3.7. Particle size distributions determined by SEM (a and c) and by
light scattering (b and d) for A-2.0GA (a and b) and A-4.0GA (c
and d).......................................................................................................... 62
Figure 3.8. XRD patterns of the samples obtained from the crosslinked
chitosan gels after (a) 12 h aging (A-12h), (b) 36 h aging (A-
36h), and (c) 72 h aging (A-72h)................................................................ 64
Figure 3.9. SEM images of the particles obtained from the chitosan gels after
(a) 36 h aging (A-36h), and (b) 72 h aging (A-72h)................................... 64
Figure 3.10. XRD patterns of the crystals in samples (a) A-0h (with alkaline
solution) and (b) A-36h (without alkaline solution)................................... 65
Figure 3.11. SEM images of the crystals in samples (a) A-0h and (b) A-36h. ............... 66
Figure 3.12. XRD patterns of the samples (a) A-1h, (b) A-3h, and (c) A-6h. ................ 67
Figure 3.13. SEM images of the particles produced after heating for (a) 3 h
(A-3h) and (b) 6 h (A-6h)........................................................................... 67
Figure 3.14. XRD patterns of (a) A-H2O2 and (b) A-cal................................................. 68
Figure 3.15. TG curves of the samples after the treatment of hydrogen
peroxide: (a) plain crosslinked chitosan (GA-CS), (b) A-cal, and
(c) A-H2O2.................................................................................................. 69
Figure 3.16. (a) XRD pattern, (b) SEM image, and (c) particle size
distribution by SEM and (d) particle size distribution by light
scattering of zeolite FAU (NaY) nanocrystals. .......................................... 70
Figure 3.17. (a) TG curve and (b) nitrogen adsorption-desorption isotherm of
zeolite Y samples ....................................................................................... 71
Figure 3.18. (a) XRD pattern and (b) SEM image of the as-synthesized
sample......................................................................................................... 73
Figure 3.19. TEM images of a typical zeolite A particle with a cube-like
morphology and the corresponding SAED patterns obtained
from the entire particle. (a) Original particle and (b) the same
particle after beam irradiation for a few minutes. ...................................... 74
Figure 3.20. Dark field TEM images of the cross sections of cubes of zeolite
A with an amorphous core. These images indicate that the shell
thickness varies in individual crystals. ....................................................... 75
List of Figures
xiii
Figure 3.21. (a) SEM image and (b) XRD pattern of the sample treated with
0.35 M acetic acid solution for 4 h. The insert in (a) is a TEM
image of boxes. .......................................................................................... 75
Figure 3.22. SEM images of sample prepared without addition of chitosan,
before (a) and after (b) treatment in 0.35 M acetic acid solution
for 4 h. ........................................................................................................ 76
Figure 3.23. XRD of the samples prepared with different hydrothermal times.
(a) 1 h, (b) 2 h, (c) 3 h, (d) 4 h, and (e) 6 h (The peak labeled by
asterisk is one of sodalite characteristic peaks.)......................................... 77
Figure 3.24. SEM images of samples prepared with different hydrothermal
times and then treated with 0.35 M acetic acid solution for 4h.
(a) 2 h, (b) 3 h, (c) 4 h and (d) 6 h. ............................................................. 77
Figure 4.1. XRD patterns of samples prepared with dried organic-
functionalized silicalites by hydrothermal treatment at 80 °C for
different times. (a)Sil-N to Sod-N, (b) Sil-C to Sod-C. ............................. 87
Figure 4.2. FT-IR spectra of samples (a) Sil-N and Sod-N, obtained after 3 h
hydrothermal reaction, (b) Sil-C and Sod-C, obtained after 4 h
hydrothermal reaction. ............................................................................... 87
Figure 4.3. SEM images (a, b, d, c and e) and particle size distributions (c, f)
of organic-functionalized silicalites and organic-functionalized
sodalites. SEM images: (a) dried silicalite Sil-N, (b) sodalite
Sod-N obtained after 3 h hydrothermal reaction, (d) dried
silicalite Sil-C, and (e) sodalite Sod-C obtained after 4 h
hydrothermal reaction. Particle size distributions: (c) Sil-N and
Sod-N obtained after 3 h hydrothermal reaction, and (f) Sil-C
and Sod-C obtained after 4 h hydrothermal reaction. ................................ 88
Figure 4.4. (a) 29Si-NMR of organic-functionalized sodalite nanocrystals and
(b) the bonding scheme for organic-functionalized sodalite
nanocrystals. ............................................................................................... 89
Figure 4.5. TGA curves of organic-functionalized sodalite nanocrystals and
hydroxyl-sodalite nanocrystals. (a) Sod-N, (b) Sod-C and (c)
Sod.............................................................................................................. 90
List of Figures
xiv
Figure 4.6. (a) Nitrogen and (b) hydrogen adsorption-desorption isothermals
of plain sodalites (Sod) and organic-functionalized sodalites
(Sod-N and Sod-C)..................................................................................... 92
Figure 4.7. Particle size distributions of organic-functionalized sodalite
nanocrystals and plain sodalite nanocrystals in different solvents:
(a) deionized water, (b) dimethylformamide (DMF), (c)
isopropanol and (d) dichloromethane (DCM). ........................................... 93
Figure 5.1. Digital photos of PI-0, PI-15, PI-25 and PI-35 showing the change
in transparency with increasing Sod-N content........................................ 101
Figure 5.2. XRD patterns for samples Sod-N, PI-0, PI-15, PI-25, and PI-35.
The peaks labeled with asterisks arise from Sod-N. ................................ 101
Figure 5.3. IR spectra of samples Sod-N, PI-0 and PI-35............................................. 102
Figure 5.4. SEM images for PI-0, PI-15, PI-25 and PI-35............................................ 103
Figure 5.5. TGA curves for samples PI-0, PI-15, PI-25, and PI-35.............................. 104
Figure 5.6. H2 adsorption-desorption isotherms of amino-functionalized
sodalite nanocrystals at 77 K and 298 K. ................................................. 106
Figure 5.7. Selectivity )( 22 NHα for PI-0, PI-15, PI-25, and PI-35 at different
temperatures (25 °C, 60 °C and 100 °C) .................................................. 109
Figure 5.8. Apparent activation energy (Ep) for PI-0, PI-15, PI-25 and PI-35. ............ 110
List of Tables
xv
List of Tables Table 2.1. Application and properties of zeolites [4, 43, 44]............................................6
Table 2.2. Annual global hydrogen production from different sources [127]. ...............27
Table 3.1. Experimental design for the zeolite A synthesis in crosslinked
chitosan gel.................................................................................................54
Table 5.1. DTG and TG results of PI-0, PI-15, PI-25 and PI-35. .................................108
Table 5.2. Gas permeation results of the PI-0, PI-15, PI-25, and PI-35
membranes. ..............................................................................................108
List of Schemes
xvi
List of Schemes Scheme 2.1. Schematic representation of zeolite formation [62]. (*SDA is
structure-directing agent). SEM-scanning electron microscopy
(a, b) and TEM-transmission electron microscopy (c, d) images
show zeolite FAU (a), LTA (b), MFI (c) and SOD (d)
nanocrystals cited from Ref [63]. ...............................................................10
Scheme 2.2. Mechanism of structure-directing crystal growth in the silicalite-
1 synthesis by using tetrapropylammonium hydroxide (TPAOH)
as structure-directing agent (SDA) [71, 73]. ..............................................13
Scheme 2.3. Generalized scheme for the synthesis of porous ZSM-5 via ion-
exchange: (1) synthesis of ZSM-5 with SDAs; (2) cleavage of
the SDAs inside the ZSM-5 pores; (3) removal of the organic
fragments; (4) recombination of the fragments into the original
SDAs [81]...................................................................................................16
Scheme 2.4. Schematic representation of the confined-space synthesis of
zeolite nanocrystals [89].............................................................................18
Scheme 2.5. The procedure for preparing hollow zeolite spheres (a); and
SEM image of hollow zeolite beta spheres (the inset shows the
hollow structure) [101]..............................................................................22
Scheme 2.6. Scheme of crosslinking reaction between chitosan (CS) and
glutaraldehyde (GA) [104]. ........................................................................26
Scheme 2.7. Reaction mechanism of imide formation [153]..........................................32
Scheme 2.8. Schematic representation of a mixed matrix membrane (MMM)
[132]. ..........................................................................................................35
Scheme 2.9. Mechanism for permeation of gases through porous and dense
gas separation membranes [138, 205]. .......................................................40
Scheme 2.10 The nano-gaps were hypothesized to exist in the silica-BPPOdp
nanocomposite membranes by Cong et al. [191]. ......................................44
Scheme 2.11. Gas permeation through mixed-matrix membranes containing
dispersed zeolite particles [138]. ................................................................45
Scheme 2.12. Summary of the relationship between MMMs morphologies
and transport properties. Solid circles represent calculated
List of Schemes
xvii
values for MMMs with an incorporation of 35 vol% zeolite 4A
and Ultem as polymer matrix.(Modified from Ref.[194]) .........................46
Scheme 2.13. Principle of coupling of organofunctional silanes onto zeolite
surface [192]...............................................................................................48
Scheme 2.14. Reaction between an amine-functionalized zeolite and the
polyimide to create a covalent amide linkage during annealing
[231, 232]. ..................................................................................................49
Scheme 3.1. Synthesis of zeolite crystals within crosslinked chitosan
hydrogels (GA-CS). ...................................................................................56
Scheme 3.2. Representation of the formation of cubes of zeolite A, with an
amorphous core (e). The rounded boxes in a) — d) are about 4
μm × 4 μm in dimension. ...........................................................................78
Scheme 3.3. (a) Crosslinking reaction between glutaraldehyde (GA) and
chitosan molecules; (b) pH dependent protonation/deprotonation
of the chitosan molecule [246]. ..................................................................81
Scheme 5.1. Apparatus for measuring gas permeance through the membrane. .............97
Scheme 5.2. Preparation of sodalite-N/PI nanocomposite membranes. .........................99
Scheme 5.3. Fabrication of BTDA-MDA polyimide by two-step method. ..................100
Scheme 5.4. Schematic representation of sodalite-polyimide interfacial
structure (a) and covalent linker between Sod-N and polyamide
(b). ............................................................................................................107
Abbreviations
xviii
Abbreviations
ADMS: 3-aminopropyl(diethoxy) methylsilane
AM: acrylamide CHCONH2
BET: Brunauer-Emmett-Teller
BTDA: benzophenone-3,3’,4,4’-tetracarboxylic dianhydride
C-PAM: crosslinked polyacrylamide hydrogels
CS: Chitosan
D: diffusivity coefficient (kinetic parameter)
DCM: dichloromethane
DMF: dimethlformamide
DTG: differential thermogravimetric
FT-IR: Fourier transform infrared spectra
GA: glutaraldehyde
GA-CS: glutaraldehyde-crosslinked chitosan
HAc: acetic acid
HRTEM: high-resolution transmission electron microscopy
J: gas flux
MBAM: N,N -methylenebisacrylamide (CH2CHCONH2)2CH2
MDA: 4,4’-diaminodiphenylmethane
MMMs: mixed matrix membranes
MTMS: methyltrimethoxysilane
MAS NMR: magic angle spinning nuclear magnetic resonance
Pi: gas permeability
PAA: poly(amic acid)
PDMS: polymer polydimethyl siloxane
PES: polyethersulfone
PI: polyimide
PSF: polysulfone
S: solubility coefficient (thermodynamic factor)
SDAs: structure-directing agents
SEM: scanning electron microscopy
Sil: silicalite
Abbreviations
xix
Sil-C: silicalite with methyl organic groups
Sil-N: silicalite with amino organic groups
SOD or Sod: sodalite
Sod-C: sodalite with methyl organic groups
Sod-N: sodalite with amino organic groups
TEM: transmission electron microscopy
TEMED: N,N,N ,N -tetramethylethy lenediamine
TEOS: tetraethyl orthosilicate
TGA: thermogravimetric analysis
TPAOH: tetrapropylammonium hydroxide
XRD: X-ray diffraction
)( 22 NHα : selectivity hydrogen over nitrogen
Chapter 1
1
Chapter 1 Introduction
Zeolites are crystalline microporous materials consisting of tetrahedral units and
producing open framework structures; which generate a system of pores and cavities
having molecular dimensions. To date, more than 40 naturally occurring zeolites are
known and around 200 synthetic zeolite types have been already reported [1-3].
Because of their high surface area (up to 1000 m2/g), high void volume (ca. 30% of the
total volume of zeolite) and uniform pore size distribution, zeolites have been widely
applied in different areas, especially gas separation or purification [3, 4].
One would expect a significant change in properties for the zeolite nanocrystals in
comparison with those of conventional micro-sized crystals [5, 6]. For example, the
decrease of the particle size causes a strong increase in the surface to volume ratios, and
this is expected to be of importance in catalytic reactions [7, 8]. Hence, zeolite
nanocrystals have been synthesized and used for applications such as in sensors,
membranes, and microelectronics [9]. Despite the definition for the exact size range of
zeolite nanocrystals differs largely according to various research, most recent studies
refer ideal zeolite nanocrystals as discrete, uniform crystals with dimensions of around
100 nm [5, 6, 10]. There have been a wide range of methods developed to synthesize
zeolite nanocrystals. The synthesis is often undertaken at hydrothermal condition from
clear precursor solutions or gels, where an organic structure-directing agent (SDA) is
used. However, nanoparticles may aggregate or disperse poorly in solvents or water
after high-temperature calcination for template removal. Therefore, a SDA-free method
has recently been considered. Three-dimensional polymer networks may be applied to
restrict the zeolite growth and lead to nanocrystal formation [11, 12]. Furthermore,
some previous research has suggested that the roles of polymer hydrogels in zeolite
synthesis are complex, and found that the addition of organic polymers may have an
effect on the zeolite nucleation and crystallization process [13, 14]. Hence, it would be
of considerable interest to further explore the feasibility and mechanisms of zeolite
crystallization in polymers or polymer hydrogels.
There is a growing interest in the further development on combining zeolite
nanocrystals and polymeric membranes for gas separation or purification. Polymeric
Chapter 1
2
membranes are considered to be an important medium for gas separation such as
hydrogen purification, nitrogen recovery from air, and natural gas purification, since
they are relatively inexpensive and can be fabricated into compact hollow fiber and flat
sheet modules with a high separation area to volume ratio [15]. To fabricate practical
zeolite-polymer composite membranes (also known as mixed matrix membranes
(MMMs)), zeolite nanocrystals are required instead of the micro-sized ones [15].
Research has proven that such membranes exhibit better performance and selectivity
than some traditional pure polymeric membranes in gas (e.g. H2) separation and
purification [15-20].
Furthermore, zeolite nanocrystals with good interfacial compatibility with the chosen
polymers are needed to fabricate zeolite-polymer mixed matrix membranes (MMMs)
[15]. The strategy of attaching organic functional groups to zeolites is one of the most
effective ways for modifying surface properties or adding surface reactivity to zeolite
nanocrystals [6, 21-26]. Thus, there needs to be further research into the organic
functionalization of zeolite nanocrystals.
The goal of this research is to develop novel methods for synthesizing, functionalizing
zeolite nanocrystals, and fabricating zeolite-polymer MMMs for gas separation. Chapter
2 of this thesis reviews the relevant literature about zeolites, zeolite nanocrystals, zeolite
synthesis and their possible applications. Chapter 3 describes the zeolite crystallization
and growth in glutaraldehyde-crosslinked chitosan (GA-CS) and uncrosslinked chitosan
hydrogels. Chapter 4 presents the results for the synthesis of organic-functionalization
of sodalite nanocrystals by applying direct transformation of silicalite nanocrystals. The
fabrication and characterization of zeolite-polymer mixed matrix membranes (MMMs)
by applying the produced organic-functionalized sodalite nanocrystals are presented and
discussed in Chapter 5. Finally, the conclusions arising from this study and
recommendations for future work are summarized in Chapter 6, respectively. My
publications relevant to this thesis are attached in Appendix section.
Chapter 2
3
Chapter 2 Literature Review
2.1 Overview The objective of the literature review is to provide a summary of the current
development of zeolites, zeolite nanocrystals and their different applications. This
chapter reviews the literature relevant to the hydrothermal synthesis of zeolites or
zeolite nanocrystals, especially by applying polymer hydrogels. Previous research about
the application of zeolites for mixed matrix membranes (MMMs) to gas separation and
some theories about gas transport through membrane modules are also reviewed.
Current research about organic functionalization of zeolites, which is suggested as an
effective way to fabricate defect-free mixed matrix membranes (MMMs), is
summarized and discussed at the end of this literature review.
2.2 Introduction to zeolite In the 18th century, a Swedish mineralogist, Axel Fredrik Cronstedt, discovered a new
mineral species. He observed that the stones began to dance about as the water
evaporated, upon rapidly heating a natural mineral. By using the Greek words which
mean “stone that boils”, it was named “zeolite”.
Zeolite is a class of microporous crystalline materials, composed of TO4 tetrahedra (T =
Si, Al) with O atoms connecting neighboring tetrahedra (Figure 2.1a) [1-4]. These
tetrahedras then link together by their corners to form a rich variety of beautiful
structures. More than 130 framework types with numerous compositional variations are
known. Each framework type is assigned with a unique three-letter code by the
International Zeolite Association [27]. Figure 2.1b shows the several framework
structures of common zeolites, including SOD, LTA, FAU, MFI, etc.
Chapter 2
4
(a) (b)
Figure 2.1. Framework structures of zeolites: (a) representations of [SiO4]4- or [AlO4]5-
and some possible connections; (b) structures of SOD, FAU, LTA, MFI and*BEA [27,
28].
The frameworks in zeolites are generally very open. They contain channels and cavities
in which water molecules and a wide variety of balance framework charges, such as Na+,
K+, Ca2+, and Mg2+, are located. Normally, the zeolite composition can be best
described as having three components: extraframework cations, framework, and sorbed
phase [4]:
+mmnM / . ][ 21 OAlSi nn− . OnH 2
extraframework cations framework sorbed phase
In Figure 2.1b, sodalite (SOD) has a general composition represented by
Na8−y[T2O4]X2−y·nH2O (T = tetrahedral framework cation, usually Si and Al; X =
monovalent ‘guest’ anion, 0≤ y ≤2, 0≤ n ≤8). It is formed by TO4 tetrahedra as
elementary building units, which are connected via corners to form six-membered rings
SOD FAU LTA
MFI *BEA
Chapter 2
5
resulting in the cubic sodalite framework. This structure contains isolated cavities,
called the β-cages [29, 30]. The pore size of sodalite is around 2.8 Å.
Zeolite A exhibits the Linde Type A (LTA) structure with a composition
Na12[Al12Si12O48]·27H2O. Its framework can be regarded as consisting of two types of
cages: sodalite cages (β-cages) and α-cages. α-cages, which are also called supercages
of zeolite A, are formed by β-cages interconnected by oxygen bridges at the double
four-membered rings. Thus, the zeolite A framework consists of α-cages by sharing the
eight-membered rings, which are considered as the windows of the α-cage with
effective pore sizes about 4 Å [31, 32].
Faujasite (FAU) is an aluminosilicate zeolite whose structure contains three-
dimensional pores; their framework consists of sodalite cages (β-cage) which are
connected through hexagonal prisms. The pores are arranged perpendicularly to each
other. Zeolite X and Y exhibit the faujasite (FAU) structure, whose chemical
composition can vary according to the silicon and aluminum content. Zeolites X has a
Si/Al ratio varying from 1 to 1.5, whilst zeolite Y has a Si/Al ratio raging from 1.5 to 3
[33, 34]. The pore size of FAU-type zeolite is around 7.3 Å [35-37].
Both silicalite-1 and ZSM-5 has a MFI structure. The MFI-type zeolite is made of a
three-dimensional network of interconnected pores. The pore network is composed of
straight channels (pore size: 0.53 x 0.56 nm) that are intercepted by zig-zag channels
(pore size: 0.51 x 0.55 nm) [38, 39]. Silicalite-1 is a pure silica zeolite with the unit cell
stoechiometry defined by the formula [SiO2]48. ZSM-5 is an aluminosilicate zeolite and
its chemical formula is Nan[AlnSi96-nO192]·16H2O (0< n <27).
Zeolite beta is a large-pore and high-silica microporous material with a three-
dimensional intersecting channel system showing *BEA structure. Zeolite *BEA is a
complex intergrowth family, which consists of two (polymorph A and B) or more
polymorphs [40]. Two mutually perpendicular straight channels, each with a cross
section of 0.76 x 0.64 nm, run in the n- and h-directions. A sinusoidal channel of 0.55 x
0.55 nm runs parallel to the c-direction [41]. Zeolite beta can be represented by a
composition [SiO2]64 [27].
Chapter 2
6
Typically, in the zeolite compositions, the extraframework cations often have a high
degree of mobility giving rise to facile ion exchange. The water molecules in sorbed
phase are readily lost and regained by the zeolite framework. Some specific properties
of zeolites stem from their microporosity and are a result of the framework topology,
which include [42]:
High degree of hydration and the behavior of “zeolitic water”;
Low density and large void volume when dehydrated;
Stability of the crystal structure when dehydrated, and as much as 50
vol% of the dehydrated crystals are void;
Cation or anion exchange properties;
Uniform molecular-sized channels in the dehydrated crystals;
Various physical properties like electrical conductivity;
Adsorption ability of gases and vapors.
At present, zeolites are commercially used in the chemical industry as filters, adsorbents
and catalysts for structured catalytic reactors taken of their specific properties, which
have been summarized in Table 2.1.
Table 2.1. Application and properties of zeolites [4, 43, 44].
Application in: Taking advantage of :
Catalysis Acidity, porosity, high surface area
Detergents Ion exchange capability
Dessicants Microporosity (adsorption), polarity, and molecular
sieve effect
Gas separation Microporosity (adsorption)
Zeolites have the ability to act as catalysts in chemical reaction, which take place within
the internal cavities. Underpinning all these types of reaction is the unique microporous
nature of zeolites, where the shape and size of a particular pore system exert a steric
Chapter 2
7
influence on the reaction, controlling the access of reactants and products. Thus a
zeolite is also named as a “shape-selective catalyst” [3].
Zeolites can also be used in detergents in industry, which make good use of their ion-
exchange property. The loosely bound nature of extraframework metal ions means that
they are often readily exchanged with other types of metals in aqueous solution. This
has been exploited in water softening, where alkali metals such as Na+ or K+ prefer to
exchange out of the zeolites, being replaced by the "hard" Ca2+ and Mg2+ from the water
[43]. From the framework of zeolites, it has also been suggested that the water
molecules are readily lost and regained. Hence, this type of porous materials can be
used as drying agents.
Their application for molecular adsorption and desorption is an attractive area of
research. The ability preferentially to adsorb certain molecules, while excluding others,
has opened up a wide range of molecular sieving applications. Sometimes it is simply a
matter of the size and shape of pores controlling molecular species to access into the
zeolite. For example, FAU-type zeolite with Si/Al ranging from 1 to 1.5 is NaX, and
that with Si/Al ranging from 1.5 to 3 is NaY. This type of zeolite materials has large
pores (ca. 0.73 nm), presenting no steric hindrance for small molecules, e.g. CO2
(kinetic diameter 0.33 nm) or N2 (kinetic diameter 0.37 nm). Therefore, FAU-type
zeolite or their zeolitic membranes have been intensively studied, such as for the CO2
capture from exhaust gases in combustion process, for the purification of natural gas,
for the CO2 separation from synthesis gas mixture, etc [35-37]. Among the zeolites,
Linde Type A (LTA) has been the objective of many applications due to its long history.
Zeolite A with sodium cations, denoted NaA contains cages with orthogonal 3-D
oriented apertures of approximately 0.4 nm [45]. As a molecular sieve, NaA can be used
in air and natural gas purification [46].
2.3 Application of zeolite nanocrystals The major interest in zeolite nanocrystals is due to their use for the construction of
structured materials as well as the preparation of zeolitic films and membranes [47-52].
The reduction of particle size from the micrometer to the nanometer scale leads to
Chapter 2
8
substantial changes in the properties of zeolites, which also have an impact on the
performance of zeolites in some traditional applications, e.g. separation or catalysis.
Zeolite nanocrystals and colloidal zeolite suspension in particular are recognized as a
very convenient source to prepare structured materials. The term “structured materials”
denotes polycrystalline extended zeolitic structures, which have a certain level of
organization. By the spatial arrangement of the nanocrystals, it can provide materials
with bi- or tri-modal pore organization, such as hollow zeolite spheres or ordered
macroporous zeolite macrostructures [10]. The produced structured materials have been
explored for their possibilities in various new application fields, such as chemical
sensors, shape-selective adsorbents and catalysts [51].
Furthermore, it is well known that zeolite nanocrystals can be used as seeds for the
tailored synthesis of porous zeolites, because of their discrete nature and homogeneous
distribution of particles in the colloid. During the synthesis, these nanocrystals can
effectively control the particle size of products, avoid the need for an organic template
and accelerate the synthesis rate [10]. Presently, this synthesis by seeding zeolite
nanocrystal gels has been applied in industrial production of zeolites.
The use of preformed zeolite nanocrystals for the preparation of zeolitic films and
membranes is another major application for zeolite nanocrystals. Generally, the certain
quantity of zeolites is coated on supports with different configurations (e.g. with flat,
tubular, fibrous, or spherical shape). When the support is removed by combustion or
dissolution after coating, the zeolite structures obtained are self-standing [10]. However,
the self-standing zeolite structures are often characterized with a poor mechanical
stability. Therefore, research has been focused on the application of zeolite-polymer
mixed matrix membranes (MMMs), which is an important research objective in this
thesis.
To date, zeolite nanocrystals have been combined with flexible polymer membranes,
which would be fabricated into thin zeolite-polymer MMMs, e.g. with 100 nm in
thickness [11, 15, 53-55]. Previous studies have suggested that the incorporation of
inorganic particles (e.g. zeolites) into the polymer matrix can improve the poor
separation ability of some traditional polymeric membranes and significantly increase
gas separation efficiency by enhancing selective gas adsorption and diffusion through
Chapter 2
9
the membranes [15, 19, 21, 56]. The development of this type of MMMs will be further
reviewed in Section 2.6.
Apart from the applications discussed above, there are some other emerging uses. For
example, the decrease of particle sizes to nano-range leads to a significant increase of
the surface to volume ratios, and this is expected to be of significant importance in
catalytic reactions [7, 8, 57]. Research also has indicated that the use of zeolite
nanocrystals as catalysts could reduce the mass transport limitations since the diffusion
path is relatively short and the accessibility of the catalytic sites through the external
surface is high [10].
2.4 Hydrothermal synthesis of zeolite nanocrystals Because of their specific characteristics and applications, the synthesis of zeolites and
zeolite nanocrystals has been extensively reviewed in several books and literature [49,
58-61].
The synthesis of zeolite is typically carried out under hydrothermal conditions, which
has also been widely applied to the synthesis of zeolite nanocrystals. A schematic
representation of zeolite formation is given in Scheme 2.1. The zeolite hydrothermal
synthesis involves two main steps: nucleation and crystallization. Nucleation is defined
as a process where the small aggregates of precursors give rise to nuclei (or called
embryos). With increasing time, the nuclei become larger and zeolite crystals form,
which is called “crystallization” [62].
Generally, the hydrothermal synthesis requires several basic conditions as follows [42],
Reactive starting materials, such as freshly coprecipitated gel, or
amorphous solid;
Relatively high pH introduced in the form of an alkaline metal hydroxide
or other strong base;
A high degree of supersaturation of the components of the gel leading to
the nucleation of a large number of crystals.
Chapter 2
10
Scheme 2.1. Schematic representation of zeolite formation [62]. (*SDA is structure-
directing agent). SEM-scanning electron microscopy (a, b) and TEM-transmission
electron microscopy (c, d) images show zeolite FAU (a), LTA (b), MFI (c) and SOD (d)
nanocrystals cited from Ref [63].
A gel which is defined as a hydrous metal aluminnosilicate, is prepared from aqueous
solutions, reactive solids, colloidal solution, or reactive aluminosilicates. The utilization
of initially clear homogeneous solutions or gels, where only sub-colloidal or discrete
amorphous particles are present, is most widely adopted. Synthesis proceeds at elevated
temperatures in a closed nutrient pool, where crystals form through a nucleation step. It
is found that the increased number of nuclei leads to a reduction in the ultimate crystal
size. Thus, the formation of zeolite nanocrystals requires the condition that favors
nucleation over crystal growth in the system. Moreover, zeolite nanocrystals have to be
Chapter 2
11
recovered by avoiding aggregation from the produced stable colloidal suspension. The
synthesized zeolites, in particular zeolite nanocrystal suspensions, are usually purified
by repeated high-speed centrifugation and redispersed in a liquid (e.g. water or ethanol)
under ultrasonication [28].
Research has shown that the one of the factors in the synthesis of non-aggregated
zeolite nanocrystals is high supersaturation, since it tends to result in high nucleation
rates, a large number of nuclei, and thus producing the smallest particle sizes [9, 64]. In
aluminosilicate gels, the supersaturation is strongly influenced by the pH of the solution
[10]. High supersaturation is achieved by high alkalinity of gel solution, which would
also permit a decrease of the synthesis temperature, favoring nucleation and minimizing
the ultimate crystal size [65].
The addition of inorganic base (e.g. NaOH) is one of the ways to achieve high alkalinity
in the initial gel solution, which has also been applied to synthesize zeolite nanocrystals.
For example, Zhan et al. prepared NaX (FAU) zeolite nanocrystals with controlled sizes
and surface properties by using inorganic base NaOH and different silicate sources (e.g.
silica colloid, fumed silica, and tetraethyl orthosilicate). The final products, ultra-fine
NaX zeolite (20−100 nm), were synthesized at 60 °C for 4 days [66]. Valentin et al.
used a Na2O:Al2O3:SiO2:H2O gel system for the synthesis of a FAU-type zeolite under
room temperature condition. A well-crystallized material (zeolite X) containing
100−300 nm spherical aggregates built of 10−20 nm nanocrystals was obtained after 3
weeks of synthesis [67]. However, in the above-mentioned synthesis, the products
obtained by using inorganic base in the synthesis gel are found highly aggregated.
Moreover, the formation of nanocrystals is achieved by adjusting crystallization
temperature or crystallization time. At low temperature which favors the preparation of
zeolite nanocrystals, long synthesis duration is required, lasting to several days or even
weeks. Furthermore, only some zeolites with specific compositions can be synthesized
by using inorganic base in zeolite synthesis gel [68].
Apart from the use of inorganic base, the condition of high alkalinity is possibly
achieved by utilizing abundant amount of organic cations and decreasing alkali cations
in the synthesis gel or completely substituting the use of inorganic base. Research found
that low content of alkali cations limits the aggregation of negatively charged
Chapter 2
12
subcolloidal particles in solution [10], thus dispersible zeolite nanocrystals are expected
to be produced. Organic structure-directing agents (SDAs), which are normally a type
of alkali-free organic basics, can be utilized by introducing organic cations into zeolite
synthesis gel. More importantly, organic structure-directing agents (SDAs) can play a
role of a pore filling agent during zeolite synthesis and structurally direct the
crystallization towards the formation of specific zeolitic structures. A large amount of
zeolite must be synthesized with the addition of SDAs, which can not be obtained by
simply using inorganic alkali solution and varying the composition of zeolite synthesis
gel. Furthermore, the choice of SDAs has been proven to affect the zeolite synthesis rate
[69, 70].
Alkali-free tetrapropylammonium hydroxide (TPAOH), which is regarded as one of the
effective templates (SDAs) in silicalite-1 (MFI) synthesis, has been widely studied.
Scheme 2.2 illustrates the mechanism of structure direction and crystals growth in the
silicalite-1 synthesis by using TPAOH [71]. It was found that the addition of TPAOH
indeed enhanced zeolite nucleation rate and fastened the silicalite-1 crystallization [72].
As shown in Scheme 2.2, the initial formation of the inorganic-organic composite is
initiated by the overlapping of the hydrophobic hydration spheres TPA+-H2O, followed
by a subsequent release of ordered water (H2O) to establish favourable interactions
between TPA+ and silicate species. Thereafter, the aggregation of these composite
species results in the silicalite-1 nucleation. Crystal growth occurs through the diffusion
of the same species to the surface of the growing crystallites, giving a layer-by-layer
growth mechanism. In this process, TPA+ molecules are located at the channel
intersections with their propyl arms extending into the linear and zig-zag channels. They
are tightly encapsulated in siliclalite-1 pores so that the calcination is normally required
to remove TPA+. In other words, this tight entrapment suggests that TPA+ molecules are
actively involved in the nucleation period and crystal growth [71, 73].
Chapter 2
13
Scheme 2.2. Mechanism of structure-directing crystal growth in the silicalite-1
synthesis by using tetrapropylammonium hydroxide (TPAOH) as structure-directing
agent (SDA) [71, 73].
Chapter 2
14
The application of SDAs in the synthesis of zeolite nanocrystals has been investigated
and reported in a wide range of studies [10, 74]. However, there may be some
drawbacks for the use of organic structure-directing agents (SDAs) in the synthesis of
zeolite nanocrystals. High-temperature calcination of products is often required to
remove SDAs and then open zeolite porosity, which leads to the aggregation of
individual particles or low crystallization [15, 75]. Wang et al. have developed a
calcination procedure. An organic polymer network formed after a reaction of
acrylamide, N,N -methylenebisacrylamide and the initiator (NH4)2S2O8, which was used
as a temporary barrier during calcination in order to prevent the aggregation of zeolite
nanocrystals [15, 75]. This method has been successfully used to produce silicalite-1
and zeolite A nanocrystals with good redispersibility, which have been applied in the
preparation of zeolite-polymer mixed matrix membranes (MMMs). Despite both SDAs
and temporary organic barrier can be burned off (>500 ºC), NOx may be released which
is produced by the combustion of organic SDAs, since most of them are quaternary
ammonium cations or amine. The required calcination process may consume more
energy, thus resulting in the increase of operating cost for the industrial-scale
production of zeolites or zeolite nanocrystals.
To avoid the high-temperature calcination for the removal of organic SDAs within
zeolite channels, solvent extraction was firstly applied to remove SDA
(tetraethylammonium hydroxide-TEAOH) from the synthesized zeolite beta (*BEA-
type zeolite) in refluxing nitric acid solutions at 80 ºC by Fajula and co-workers in 1993
[76]. However, the removal of SDA was not complete, and also a concomitant
formation of significant mesoporous volume and loss of some zeolite microporosity
occurred. Davis’s group used tetraethylammonium fluoride (TEAF) as SDA, and their
experimental results showed that tetraethylammonium fluoride (TEAF) was more easily
extracted out of the zeolite beta molecular sieves by using heating acetic acid than the
removal of TEAOH [77]. They attributed this to the weak interaction between zeolite
beta and TEA+ in the form of TEAF. Similar work has also been conducted by
Takewaki et al. and other groups [25, 78, 79]. It was concluded that a complete solvent
extraction process required structure-directing agents with a smaller molecular size than
the pore opening of zeolites and the weak interactions with the zeolite framework.
Therefore, only a few successful solvent extractions applied for SDAs removal have
Chapter 2
15
been reported, all of which are limited to synthesize the molecular sieves of *BEA
topology involving tetraethylammonium ions (TEA+)-based organic SDAs [80].
In 2003, Davis et al. developed a novel method to cleave SDAs within zeolite pore
spaces via ion-exchange process, which is shown in Scheme 2.3 [81, 82]. The key
feature of this method is that the organic SDAs can be disassembled by changing pH in
the reaction condition (Scheme 2.3-2), to allow the removal of their fragments from the
zeolite pore spaces (Scheme 2.3-3). The organic molecules from SDAs can be
recombined into the original SDAs for further zeolite synthesis (Scheme 2.3-4). The
possible recycling of the SDAs may significantly decrease the overall costs in the
zeolite synthesis. In their experiment, a commercially available SDA, 1,4-dioxa-8-
azaspiro [4,5] decane, was selected in the synthesis of ZSM-5 (MFI). Holmberg et al.
also applied a similar method to remove tetramethylammonium bromide (TMABr)
from the synthesized zeolite Y [83]. The produced zeolite with SDA was added to an
ion-exchange solution, sodium nitrate (NaNO3), which was then sealed in a
polypropylene bottle and reacted at 90 ºC for 12 h. After ion exchange, almost all the
TMA+ ions located within the supercages of the zeolite Y structure were removed and
replaced with Na+ ions. It is noted that the ion-exchange method requires structure-
directing agents (SDAs) with weak interactions with the zeolite framework.
Furthermore, the synthesized zeolites may need to have large channels or pores, which
ions can enter and exchange with SDA cations. Hence, this method may not be suited to
removing the SDAs in the zeolites with small channels, such as sodalite.
Recently, Wang’s group also reported a novel method to obtain colloidal sodalite
nanocrystals, free of SDAs by the direct transformation of silicalite nanocrystals,
especially without high-temperature treatment [84]. Sodalite (SOD) is a small-pore
zeolite whose framework consists of a six-membered ring aperture with a pore size of
2.8 Å (Figure 2.2) [85]. Because of the unique pore size, only some particular molecules,
such as helium (2.58 Å) and water (2.64 Å) can enter the pores of sodalites [86, 87]. To
prepare dispersible sodalite nanocrystals, the silicalite nanocrystals were used as the
silica source for this process because silicalite nanoparticles made with
tetrapropylammonium hydroxide (TPAOH) as a template have a high resistance against
dissolution in alkaline solution. The as-synthesized colloidal silicalite nanocrystals were
directly dried so that the excess TPAOH molecules and a small amount of silica species
Chapter 2
16
coated around the nanocrystals [88]. XRD patterns showed that the directly dried
silicalite nanocrystals retained the MFI structure when they were in contact with the
alkaline solution at room temperature. During subsequent treatment at 80 °C, the
alkaline solution attacked the silicalite structures, allowing sodium and aluminum ions
to enter the zeolitic lattice. Meanwhile, the TPA+ ions in the silicalite zeolitic channels
were driven out of the nanocrystal frameworks by exchanging with the high
concentration of Na+ ions. After the incorporation of sodium and aluminum and the
transformation of the silicalite crystal structure, the resulting sodalite nanocrystals
possess small sizes — averagely 60 nm, which morphologies were similar to those of
the precursor silicalite nanocrystals. Additionally, the sodalite nanocrystals can be
readily dispersed in water or ethanol under mild ultrasonication, and the colloidal
sodalite suspensions thus formed were stable for weeks [84].
Scheme 2.3. Generalized scheme for the synthesis of porous ZSM-5 via ion-exchange:
(1) synthesis of ZSM-5 with SDAs; (2) cleavage of the SDAs inside the ZSM-5 pores;
(3) removal of the organic fragments; (4) recombination of the fragments into the
original SDAs [81].
Chapter 2
17
(a) (b)
Figure 2.2. Sodalite: (a) unit cell containing β-cage (b) framework [85].
2.5 Use of polymers in zeolite synthesis
2.5.1 Confined-space synthesis of zeolite nanocrystals As mentioned earlier, polymers have been applied as temporary barriers to prevent the
aggregation of nanocrystals under high-temperature calcination for template removal
[15, 75]. The use of polymers in confined-space synthesis of zeolite nanocrystals has
also attracted great interest. Confined-space synthesis, in which an inert matrix is used
to provide a steric hindered space for zeolite nanocrystal growth has been widely
studied [68, 89-92], since the first publication by Madsen and Jacobsen in 1990s [89]. A
schematic illustration for this method is shown in Scheme 2.4.
Typically, an inert matrix, such as carbon black, may be impregnated with clear zeolite
precursor solution [89]. The impregnated matrix can then be transferred into a porcelain
cup, which is then treated in an autoclave with sufficient water providing saturated
steam at high temperature. The zeolite nanocrystals are formed and confined by the pore
spaces of inert matrix. Generally, the crystal size distributions of the zeolites obtained
are governed by the pore sizes of inert matrix. As mentioned in Section 2.4, zeolite
nanocrystals prepared from the conventional hydrothermal synthesis method normally
requires high-speed centrifugation for recovery due to their colloidal aspect. The
recovery of the zeolite nanocrystals prepared by confined-space synthesis can be easily
achieved by simple calcination, during which both the inert matrix and the structure-
directing agents (SDAs) can be removed. Some of the large aggregates caused by, e.g.
calcination, may be removed from zeolite nanocrystal suspensions by filtration [10, 68].
Chapter 2
18
Scheme 2.4. Schematic representation of the confined-space synthesis of zeolite
nanocrystals [89].
In confined-space method, it is essential to apply incipient wetness impregnation
method and load the porous matrices with zeolite synthesis gel, followed by a
hydrothermal reaction within the pores of the matrices. To successfully produce zeolite
nanocrystals, there are two basic requirements for the matrices: (1) confined-space
matrices must be inert and stable in hydrothermal synthesis conditions; (2) the matrices
should possess a narrow pore size distribution, which can help yield the products with
uniform particle sizes [10].
Currently, carbon is one of the most commonly applied inert matrices for the confined-
space synthesis of zeolite nanocrystals. Schmidt et al. adopted this method to prepare
nanosized silicalite-1 (20−75 nm), zeolite X (22−60 nm), and zeolite A (25−37 nm) by
using mesoporous carbon blacks as an inert matrix [68]. The carbon matrix was finally
removed by combustion, recovering the pure and highly crystalline zeolite products.
However, the resulting nanocrystal size distribution is relatively broad, which may be a
result of the wide pore size distribution of the carbon matrix. Kim et al. used a different
type of mesoporous carbon with much more uniform pore size distributions which was
formed through the carbon replication of meso-structured silica and through the colloid
imprinting of pitch [93]. The synthesized ZSM-5 (MFI) zeolite nanocrystals had highly
uniform crystal size distributions and average sizes of 13, 22, 42, and 90 nm, which
Chapter 2
19
were close to that of colloid-imprinted carbon templates with average pore sizes of 12,
22, 45 and 85 nm, respectively.
The carbon nanotubes were also reported and used for the synthesis of zeolite
nanocrystals. For instance, Pham-Huu et al. synthesized nano-sized zeolite beta, which
belongs to an interesting class of inorganic heterogeneous catalysts for petrochemical
reactions, inside the multi-walled carbon nanotubes (MWNTs) [94]. The unique
characteristic of the MWNTs: high volume-to-weight ratio, allowed for synthesizing a
large amount of zeolites with only adding a small amount of template. Using MWNTs
as confined-space matrix also made the recovery of nanocrystals easy. The final product
obtained was washed with deionized water and then carbonized at 550 °C to decompose
the SDA (tetraethylammonium hydroxide-TEAOH), followed by a further calcination at
650 °C to completely remove the carbon MWNTs. The produced zeolites had an
average diameter between 50 and 80 nm in the form of zeolite nanowires composed of
10 nm particles. However, the expense for fabricating MWNTs is much higher than that
of zeolites, which makes such zeolite production unsuitable in practice.
Naik et al. have reported the use of CTAMeBr (cetyltrimethylammonium bromide)
surfactant as an inert matrix for the silicalite-1 crystal growth [95]. The precursor
nanoparticles formed in a clear and low-alkalinity TPA-silicate synthesis solution were
collected and protected by the addition of CTAMeBr before the crystallization. After
steamed at 150 °C, the precursor particles were converted into silicalite nanocrystals.
Because of the interruption and dilution with CTAMeBr and ethanol during the
hydrothermal process, the induction time was extended and the nanocrystals with
smaller than 30 nm were produced. The final pure silicalite-1 sample was obtained by
undergoing a conventional calcination, instead of using high-speed centrifugation to
collect samples. Furthermore, compared with the application of carbon black, it can
avoid applying a large amount of inert matrices used in the confined-space synthesis.
However, the obtained nanocrystals had poor redispersibility in water, and their thermal
stability was not satisfactory [95].
Zeolite NaY with a size in the range from 50 to 100 nm has been synthesized by using
starch as a matrix [91]. This synthesis was performed without adding SDAs and yielded
smaller zeolite crystals compared with the synthesis conducted in the absence of
Chapter 2
20
confined-space additives. However, the resultant zeolites had a broad crystal size
distribution.
In all confined-space synthesis by using inert matrices summarized-above, calcination is
commonly required to remove the porous carbon black, carbon nanotube or starch
templates. The calcination temperature varies from 300 ºC to 550 ºC based on the
thermal properties of the applied inert matrices. Some studies have found that the
produced zeolite nanocrystals after high-temperature combustion have a poor
redispersibility in water or organic solvents [75]. This limits the application of zeolite
nanocrystals, such as in fabricating zeolite-polymer nanocomposite membranes and
hierarchical porous zeolitic structures.
Therefore, polymers or polymer hydrogels have been considered to be a possible
candidate matrix for the confined-space synthesis. Wang et al. developed a confined-
space synthesis of zeolite nanocrystals by applying polymer hydrogels. As a class of
soft space-confinement additives, polymer hydrogels comprise three-dimensional
networks that are created via physical or chemical crosslinking [96, 97], which can be
readily introduced into zeolite synthesis due to good compatibility between zeolite
precursors and polymer gels [12]. They have also demonstrated the controlled synthesis
of zeolite nanocrystals in chemically crosslinked polyacrylamide hydrogel and
physically crosslinked thermoreversible methyl cellulose hydrogels [11, 12].
Polyacrylamide hydrogel (C-PMA) was prepared by the water soluble organic
monomers acrylamide CH2 CHCONH2 (AM), and N,N -methylenebisacrylamide,
(CH2CHCONH2)2CH2 (MBAM), and the initiator (NH4)2S2O8. The monomers can
polymerize and crosslink via a free-radical polymerization into an elastic hydrogel once
the temperature is increased to 50 °C or a catalyst [N,N,N ,N -tetramethylethy
lenediamine (TEMED)] is added at room temperature [98]. The crystal sizes of
produced SAPO-34 molecular sieves were substantially reduced in the crosslinked
polymer hydrogels, followed by a vapor phase transport process [12]. However, the
synthesized SAPO-34 nanocrystals exhibited a very poor dispersibility in solvents.
Similarly, NaA (20-180 nm in size) and NaX (10-100 nm in size) nanocrystals were
synthesized by employing methylcellulose hydrogels to confine crystal growth [11].
Chapter 2
21
Methylcellulose is a type of thermoreversible polymer hydrogel, and used as space-
confinement additive because of their specific gelation behavior, which is reversibly
responsive to temperature. In particular, this polymer that gels at elevated temperatures
and turns back to solution at room temperature is attractive, since the temperature
profile of their solution-gel transition nicely fits that of hydrothermal synthesis of
zeolites. By using this polymer hydrogel, no organic templates (SDAs) were needed
during the synthesis process. Hence, by the restriction of confined-space pores from this
thermoreversible polymer hydrogel and avoiding high-temperature burning of
templates, zeolite nanocrystals were produced with high dispersibility in both water and
ethanol [11].
2.5.2 Effect of added polymers on zeolite structure Polymers have been developed as additives in confined-space synthesis. Actually,
organic additives are also known to possibly affect zeolite nucleation and growth [13,
14]. As early as 1990, Dutta et al. found that cosolvents: dimethyl sulfoxide (DMSO)
and hexamethylphosphoramide (HMPA) had an important effect on the nucleation of
zeolite A and X. The addition of cosolvents speeded up the zeolite crystallization
process [13]. Myatt et al. reported their results about the crystallization of NaA zeolite
in the presence of various water-soluble surfactants (sodium dodecyl sulfate, sodium
dioctylsulfosuccinate, cetyltrimethylammonium bromide) and of organic polymer
(poly(ethylene glycol)) compared with a crystallization in the absence of additives [14].
The additives were shown to dramatically shorten prenucleation and nucleation periods
and accelerate crystal growth [14]. The addition of all surfactants or polymer, except
sodium dioctylsulfosuccinate, increased the total number of nuclei produced, giving
crystals with a reduced mean size and narrower size distribution. This result was
attributed to an effective reduction in the water content of the hydrothermal system by
adding organic additives. The addition of sodium dioctylsulfosuccinate was different
that it reduced the number of nuclei and produced larger crystals, ascribed to a specific
interaction between sodium dioctylsulfosuccinate and aluminum species. Also as
discussed in Section 2.5.1, in Naik’s study, by the interruption and dilution with
CTAMeBr (cetyltrimethylammonium bromide) surfactant and ethanol during the
Chapter 2
22
hydrothermal process, the zeolite crystallization was extended and the nanocrystals with
sizes smaller than 30 nm were finally produced [95].
Recently, the formation of unique hollow zeolite structures, which would be induced by
the addition of polymers or hydrogels in zeolite synthesis, has attracted much attention
[99, 100]. This is because that the hollow structures may exhibit more attractive
properties for their applications ranging from catalysis to electronic devices in the areas
of chemistry, biotechnology and materials science [101, 102].
(a)
(b)
Scheme 2.5. The procedure for preparing hollow zeolite spheres (a); and SEM image of
hollow zeolite beta spheres (the inset shows the hollow structure) [101].
One of the traditional polymer materials for the preparation of hollow zeolites is
polystyrene (PS) microsphere, which is used as a template. This is mainly attributed to
Chapter 2
23
the relatively easy preparation of PS microspheres with mono-dispersibility and
adjustable diameters (ranging from nanometers to micrometers) [103]. As shown in
Scheme 2.5a, typically, the PS spheres are firstly charged by sequentially depositing
several layers of cationic and anionic polyelectrolytes (Scheme 2.5a-i). The pre-
synthesized zeolite nanocrystals (Scheme 2.5a-ii) and oppositely charged
polyelectrolytes (e.g. poly(diallyldimethylammonium chloride) (PDDA)) (Scheme 2.5a-
iii) are then alternately deposited onto the charged PS substrates to form zeolite
nanocrystals/PDDA multi-layers. The seeded PS spheres subsequently undergo a
hydrothermal reaction, thus a zeolite layer can form on PS surface. The PS template is
finally removed by calcination (Scheme 2.5a-iv), resulting in a final product with
hollow structures. Until now, by changing the composition of zeolite synthesis gel and
thickness of deposited layers, a series of zeolite-type hollow spheres have been reported,
including silicalite-1 and zeolite beta (Scheme 2.5b) [63, 101].
Wang’s group reported hollow zeolite structures including sodalite spheres and hollow
zeolite NaA crystals which were synthesized by introducing crosslinked polyacrylamide
(C-PAM) hydrogels into zeolite synthesis gels [99]. From their experimental results, the
formation of hollow sodalite spheres and zeolite A crystals was explained by the
facilitation of zeolite nucleation and crystallization by polymer hydrogel networks.
After free-radical polymerization at the hydrothermal synthesis temperature (e.g. 90
°C), the zeolite synthesis gels were entrapped in individual micro-sized three-
dimensional (3-D) crosslinked polymer pores. Hydrophilic polyacrylamide (C-PAM)
possessing abundant amide groups are highly compatible with zeolite synthesis gels [12,
75], and substantially affect zeolite nucleation and crystallization. As a result, the
interfaces between the swollen polymer networks in solution and zeolite gel presumably
served as ideal nucleation sites, where nuclei would form rapidly. Zeolite nanoparticles
subsequently grew and aggregated within C-PAM hydrogel networks. Consequently,
hollow sodalite spheres or zeolite A crystals developed via consumption of zeolite
synthesis gels, which were located in the centre of aggregates through the solution-
mediated process. Based on the experimental results, the formation of these hollow
structures was attributed to the surface-to-core crystallization mechanism, which was
also reported by Chen et al. in their synthesis of hollow zeolite analcime structures
[100].
Chapter 2
24
Zeolite analcime with a core-shell and hollow icositetrahedron architecture was
prepared by a one-pot hydrothermal route in the presence of ethylamine and Raney Ni
[100]. Chen et al. explained that the formation of core-shell crystal morphology may
depend on two factors: intrinsic crystal structure and synthesis conditions. In particular,
the presence of organic additive — ethylamine, possibly interacting with the {111}
planes of zeolite analcime, inhibited crystal growth in the [111] direction at an early
stage of the formation of the nanoplatelets and reduced the potential of possible
aggregation on the (111) surface. Following that, the formed [111]-oriented
nanoplatelets stacked into discus-shaped aggregates and further self-assembled into
polycrystalline microspheres. Those nanoplatelets composing microspheres were
observed to continuously grow into large nanorods; meanwhile, the surface of the
microspheres recrystallized into a single crystalline thin shell and the icositetrahedral
morphology gradually developed, resulting in hollow-structure analcime [100].
2.5.3 Chitosan hydrogels
Given the success of those studies mentioned above, it would be of considerable interest
to further explore and investigate the zeolite nucleation and crystallization in other types
of polymers or polymer hydrogels.
Chitosan (CS) is a linear polysaccharide composed of randomly distributed β-(1-4)-
linked D-glucosamine (deacetylated unit) and N-acetyl-D-glucosamine (acetylated unit)
(Figure 2.3b). It can be obtained by extensive deacetylation of chitin (Figure 2.3a),
which is found in a wide range of natural sources, such as crab, lobster and shrimp
shells [104]. The amino group in chitosan has a pKa value of ~6.5, thus, chitosan is
positively charged and soluble in acidic to neutral solution with a charge density
dependence on pH. Moreover, because of its aliphatic primary amino groups being
regularly distributed along the polymer backbones, some chemicals, such as glyoxal
[105], glutaraldehyde (GA) [106-108], formaldehyde [109, 110], and epichlorohydrin
[111] can be applied to crosslink chitosan, forming more rigid polymer networks.
Chapter 2
25
Figure 2.3. Chemical structures of chitin polymer (a) and chitosan polymer (b) [104].
Scheme 2.6 shows the crosslinking reaction between chitosan polymer and
glutaraldehyde, which is the most widely studied crosslinker, forming three-
dimensional hydrogel networks [104]. This crosslinking occurs among one
glutaraldehyde molecule and two chitosan unities, involving the formation of two Schiff
bases. The crosslinking reaction is quite quick, and sometimes takes less than 1 h [112].
The produced hydrogel can swell in aqueous solution, and may degrade in some
solutions by changing pH, temperature and salt composition [112, 113]. Furthermore,
the addition of hydrogen peroxide may affect the stability of the crosslinked chitosan
hydrogels [114-118].
The glutaraldehyde-crosslinked chitosan (GA-CS) hydrogels have been studied for
various applications such as in pervaporation separation through chitosan [119] or
chitosan-zeolite membranes [120], enzyme immobilization [121], cationic specimen
transportation [122], controlled ingredient-release [123, 124], environmental
applications [125] and fuel cells [126]. However, no work has been focused on the
zeolite crystallization and growth in, either crosslinked or uncrosslinked chitosan
hydrogels.
(a)
(b)
Chapter 2
26
HOCH2
NH2
O
HO
CH2OH
O
NH2
O
HO
O
HO
NH2
HOCH2NH2
O
HO
CH2OH
O O
n
HO O H
OOO
HO
CH2OH
HOCH2
O
HO
O
HOO
O
HO
CH2OH
HOCH2NH2
NH2
NH2
NH2
OH OH
n
O O
H HHH
OO
Scheme 2.6. Scheme of crosslinking reaction between chitosan (CS) and glutaraldehyde
(GA) [104].
OHO
n
OO
O
HO
CH2OH
NHOCH2
N
O
HO
O
HO
N
OO
HO
CH2OHN
HOCH2
H
n
OHH O
N
N
N
NHOCH2
O
HO
CH2OH
OO
HO
O
HO
HOCH2
O
HO
CH2OH
O O
HH HH
Chapter 2
27
2.6 Application of mixed matrix membranes (MMMs) to hydrogen separation
2.6.1 Hydrogen separation
Table 2.2. Annual global hydrogen production from different sources [127].
Source Natural gas Oil Coal Electrolysis Total
Billion cubic meters/yr 240 150 90 20 500
Share (%) 48 30 18 4 100
Hydrogen is used commercially in petroleum and chemical processing to produce
syngas, ammonia, methanol, higher alcohols, urea, and hydrochloric acid [128]. It is
also used as a reducing agent in metallurgy, and in Fischer Tropsch reactions to upgrade
petroleum products and oils by hydrogenation and hydrocracking [129, 130]. Presently,
with interest in the exploration and development of clean energy, hydrogen is expected
to have a major role as an energy carrier in future energy supply [127]. Because one
important advantage of hydrogen over other fuels is that its only major oxidation
product is water vapor; its use produces no CO2 and no toxicity or ozone-forming
potential. Therefore, there is a growing demand for hydrogen in chemical
manufacturing, petroleum refining, and the newly emerging clean energy concepts will
be placing even greater demand on hydrogen supply. H2 production from a variety of
primary sources worldwide is shown in Table 2.2. Almost half of industrial hydrogen is
currently produced from natural gas by steam reforming, partial oxidation and auto-
thermal reforming. This process provides over 240 billion m3/yr hydrogen gas, holding
48% of global hydrogen production. As seen, in the production of hydrogen from
natural gas, there is a need to purify or separate hydrogen from other gases before use.
Clearly, an improved H2 separation technology can offer substantial benefit [131]. The
gas separation by membranes is a dynamic and rapidly booming field [132], since
membrane separation has significant advantages over other processes, including low
capital and operating costs, lower energy requirements [133, 134]. Therefore, the
membrane-based gas separation is an exclusively employed device in current
Chapter 2
28
commercial process [135-137], which also has attracted great interest for its application
in hydrogen separation or purification area.
2.6.2 Polymer membranes for hydrogen separation
The choice of membrane materials for gas separation applications is based on their
specific physical and chemical properties, since these materials should be designed and
tailored in an efficient way to separate the gas product with high purity from mixtures.
Moreover, the materials with long-term stability are required in the industrial membrane
separation. Previous research has concluded that the gas separation performance of
membranes basically depends upon [137]:
Membrane permeability and separation factor/selectivity;
Membrane structure and thickness;
Membrane module configuration, e.g. flat sheet or hollow fiber;
and the membrane system design.
To date, there have been a large number of polymeric materials investigated and
developed for gas separation, hydrogen separation or purification in particular.
Generally, polymers can be divided into two broad categories: rubbery and glassy. In a
rubbery polymer, segments of the polymer backbones can rotate freely around their axis,
making this polymer soft and elastic. Thermal motion of these segments also leads to
high permeant diffusion coefficients (D), which will be further discussed in Section
2.6.4. The glassy polymer is relatively tougher, more rigid, and exhibits better impact
resistance than rubbery polymer. This is caused by steric hindrance along the polymer
backbones prohibiting rotation of polymer segments. In other words, thermal motion in
this type of material is limited, thus low permeant diffusion coefficients (D) are
obtained. If the temperature is elevated, the glass transition of a glassy polymer occurs
as the increase in vibrational (thermal) energy is sufficient to overcome the steric
hindrance restricting rotation of polymer backbone segments. This temperature is called
the glass transition temperature (Tg), defining the polymer changes from a glassy to a
rubbery state. At this point, the mechanical behavior of the polymer changes from rigid
and brittle to tough and leathery, which is defined as “plastic behavior” [138].
Chapter 2
29
A useful theory describing glass transition was developed by Fox and Flory based on
the “free volume” model reported by Cohen and Turnbull in 1959 [139]. In this theory,
the extent of molecular motion depends on the membrane free volume. It is assumed
that the total volume inside polymers can be divided as occupied and free volume.
When the temperature is decreased to the glass transition temperature (Tg), the free
volume reaches a critical value, which is not sufficient for molecules to adjust and glass
transition occurs. Below the Tg, both the quantity and the spatial arrangement of free
volume remain unchanged [140]. According to the free volume theory, a molecule does
not need to obtain specific energy to overcome an activation energy barrier, but it can
undergo translational motion by simply jumping into free-volume holes arising from the
continuous redistribution of free volume within the material [141]. A concept of
fractional free volume (FFV) defined by the occupied volume and free volume is shown
in Equation 2-1 [142, 143]. The glass transition from glassy to rubbery polymers occurs
when the fractional free volume (FFV) reaches the standard value of 0.025 ± 0.003
[144-147]. Several works have shown a good correlation between FFV and the gas
permeability coefficient of various polymers [143, 148]. Thus, FFV can be used to
evaluate the diffusivity and diffusity selectivity, and partly affects solubility, which are
important properties for gas separation membranes.
T
T
VVVFFV 0−
= ……………………….Equation 2-1
where VT is the specific volume at temperature T, and V0 is the volume occupied by the
molecules at 0 K per mole of repeated unit of the polymer.
In a gas separation process, polymeric membranes generally undergo a trade-off
limitation between the permeability and selectivity, in other words, the permeability
decreases as the selectivity increases. In 1991, Robeson collected a large number of
permeation data for different polymeric membranes, and summarized so-called Robeson
trade-off limits in a number of gas pairs, including H2/N2, He/N2, O2/N2, and H2/CH4
[150]. In Robeson trade-off limits, an upper bound exists in each log-log plot of
selectivity versus permeability, which can be used to evaluate the gas separation
performance of polymeric membranes. Figure 2.4 shows an example of the Robeson
Chapter 2
30
upper bound relationship for H2 to N2 gas separation, including the initially published
data in 1991 and updated data in 2008 [149-151].
Figure 2.4. Upper bound correlation for H2/N2 separation (Prior upper bound was firstly
published in 1991 and the present upper bound was updated in 2008.) [149].
There have been a large number of polymeric materials investigated for hydrogen gas
separation, including cellulose acetate, polyimide, and polysulfone, some of which has
been commercialized [137, 152]. Polyimide is selected as a polymer candidate in my
study. During the past three decades since the commercialization of polyimide —
Kapton, an impressive variety of polyimides have been synthesized, because of both
scientific and commercial interest. It is known that polyimides possess outstanding
properties, such as thermoxidative stability, high mechanical strength, high modulus,
excellent electrical properties, and superior chemical resistance [153-157]. Therefore,
they can be used as insulation layers for semiconductor devices or substrates for flexible
printed circuits. However, more attention has been paid to the use of polyimide
Chapter 2
31
polymers as gas separation membrane materials. Various gas separation performances
have been observed in the polyimides, which depend on their different molecular
structures [149, 156-158]. Generally, gas selectivity and permeability may be controlled
by some factors, such as stiffness of polymer backbone; the free volume and its
distribution in polymer; and the penetrant-polymer interaction; etc. Other factors
including the curing and casting procedure employed to make polyimides may also have
an effect on their gas separation performance [159]. For instance, hydrogen can be
separated efficiently from the gaseous mixture using polyimides because this type of
membranes allows hydrogen to permeate faster than other gases excluding water vapor
[160]. Based on the currently reported results, fluorinated polyimides usually possess
higher H2 permeability and lower selectivity over other gases as compared with non-
fluorinated polyimides [149, 156, 157]. 6FDA (2,2-bis(3,4-dicarboxyphenyl)
hexafluoropropane dianhydride)–DDBT (3,7-diamino-2,8(6)-dimethyldibenzothiophene
sulfone) derived polyimide exhibits a H2 permeability of 156 Barrers and a H2/CH4
selectivity of 78.8 whereas BPDA (3,3',4,4'-biphenyltetracarboxylic acid dianhydride)-
ODA (4, 4'-oxydianiline) derived polyimide has a H2 permeability of 1.33 Barrers and a
H2/N2 selectivity of 365 [149, 156]. Among different types of polyimides, a polyimide
with a moderate selectivity for hydrogen gas made by benzophenone-3,3’,4,4’-
tetracarboxylic dianhydride and 4,4’-diaminodiphenylmethane was chosen in my thesis
to fabricate zeolite-polyimide nanocomposite membranes.
Apart from the selection of polymeric membrane materials with the appropriate
monomers, the fabrication process is another key issue in the successful fabrication of a
polymeric membrane with high gas separation performance [153]. Dense and flat sheet
polyimide membranes can be prepared by casting viscous polymeric solution on a flat
plate, followed by solvent evaporation to produce a flat and uniform polymer film. The
membranes fabricated by this method have been widely used in laboratory work to
characterize membrane properties. It is known that polyimides are particularly good
materials with extremely high hydrogen permselectivity among the many other types of
investigated polymers, however they have a low permeability problem, especially for
non-fluorinated polyimides. Therefore, the dense and flat membranes actually have
limited use in the industrial gas separation because of low transmembrane flux. A
possible way to overcome this drawback is to design the active layers, which serve to
separate gas mixture, to be ultrathin (with ~100 nm thickness) on the porous support or
Chapter 2
32
gutter in order to obtain a practical gas permeation rate. The support only reinforces the
surface thin layer and hardly affects gas permeation [160]. Currently, hollow fiber
polyimide membranes have been widely investigated and applied in hydrogen industrial
separation, offering numerous advantages, e.g. orders of magnitude larger surface area
packaged in a given volume of modules [156, 160]. The fabrication technique of hollow
fiber membrane includes fiber spinning, sometimes followed by thin film casting and
polymer coating [160].
O
O
O
NH2
O
COOH
NH
N
O
O
-H2O
+
Scheme 2.7. Reaction mechanism of imide formation [153]. In this thesis, to characterize the properties of zeolite-polyimide nanocomposite
membrane, dense and flat membranes are prepared by solution casting method.
Chapter 2
33
Typically, the polyimide membranes can be formed by “two-step” polymerization. The
monomers include diamine and dianhydride. When a diamine and a dianhydride are
added into dipolar aprotic solvent such as N, N-dimethylacetamide, poly(amic acid) is
rapidly formed at ambient temperature. The reaction mechanism involves the
nucleophilic attack of the amino groups on the carbonyl carbon of anhydride groups,
followed by the opening of the anhydride ring to form amic acid groups as illustrated in
Scheme 2.7. The most important aspect of this process is that it is an equilibrium
reaction. Often it appears to be an irreversible reaction because a high-molecular-weight
poly(amic acid) is readily formed in most cases as long as pure reagents are used. It
should be also noted that the acylation reaction of amines is an exothermic reaction and
that the equilibrium is favored at lower temperature, thus ice bath is commonly applied
[161]. After elevating temperature, the poly(amic acid)s lose water thus forming the
polyimide product, which can be applied for the following membrane characterization
[138].
2.6.3 MMMs for hydrogen separation
As discussed in Section 2.6.2, for polymeric materials, a rather general trade-off limit
exists between permeability and selectivity, called Robeson “upper bound”. Previous
studies considered that the Robeson “upper bound” would represent the asymptotic end
point in the performance of polymeric membranes whose separation properties are
governed by solution-diffusion transport mechanisms [132]. Therefore, in order to
obtain a polymeric membrane with an improved performance beyond the Robeson
“upper bound”, some methods have been investigated. One way is to modify traditional
structures of polymeric membranes and develop a novel group of polymers. For
example, in 2004, a new class of microporous glassy polymers, called “polymers of
intrinsic microporosity” (PIMs) was introduced by Budd’s group [162-164]. Because of
no rotational freedom in the polymer backbones, this group of polymers has very high
free volume. Later studies, reported by Thomas et al. showed that the separation results
placed PIMs above the Robeson “upper bound” in the oxygen/nitrogen separation [165].
Another possible method to improve the gas separation performance of polymeric
membranes comes to consider the combination with inorganic materials [132],
including
Chapter 2
34
Carbon molecular sieves [166-169];
Nonporous silica [170-174];
Zeolites [17, 175, 176];
Others, e.g. C60 [177], graphite [178], activated carbons [179], etc.
In fact, those inorganic materials have been studied as inorganic membranes for gas
separation. Kusuki et al. prepared carbon molecular sieve hollow fiber membranes by
pyrolyzing an asymmetric 3,3’,4,4’-biphenyltetracarboxylic dianhydride (BPDA)-based
polyimide hollow fiber membranes in a nitrogen stream. The membranes which were
pyrolyzed over 700 °C displayed excellent H2/CH4 separation performances. Hydrogen
permeability of resultant carbon membranes ranged from 10−4 to 10−3 cm3
(STP)/(cm2·s·cmHg) and the ratios of hydrogen permeation rate to that of methane
varied from 100 to 630 at a feed gas composition of 50% hydrogen in methane [180].
Silica is a type of non-porous materials, which tends to be attractive and there has been
much advancement in controlling the structural formation of microporous silica
membranes to achieve high-purity H2 separation applications [181-183]. The silica
membranes prepared in the work reported by Kim et al. had H2/N2 selectivity at around
80 at 100 ºC with a hydrogen permeability 240×l0-9 mol·m-2·s-l·Pa–l [184]. Upon
controlling the micropore sizes of silica membranes, there have been different results
about hydrogen permeability and selectivity reported [183]. As reviewed, zeolite is a
class of microporous materials, which can be used as molecular sieves for gas
separation. For instance, Xu et al. reported the H2/n-C4H10 permselectivity of the
hydroxy-sodalite zeolite membrane was higher than those of the other types of zeolitic
membranes reported in the literature. The high H2/n-C4H10 permselectivity of hydroxy-
sodalite zeolitic membrane was greater than 1000, which had hydrogen permeance
1.14×10−7 mol·m−2·s−1·Pa−1 [185, 186]. The high selectivity of sodalite membranes was
attributed to the small channel sizes (0.28 nm) of hydroxy-sodalite zeolite, which has a
six-membered ring of Si–O–Si bond. A similar H2/N2 selectivity result was also
reported, which was higher than 1000 [185]. It has been found that many of the
inorganic membranes have excellent gas separation properties lying far beyond the
upper-bound limit for the organic polymers [132]. However, most of the inorganic
membranes are fragile. The difficulty in controlling preparation of defect-free inorganic
membranes, e.g. zeolitic membranes, makes their scale-up difficult, limiting their
Chapter 2
35
widespread practical applications. Therefore, an attempt to incorporate inorganic
particles into polymeric organic matrix, which has better flexibility and attractive
separation property, has attracted growing interest.
Scheme 2.8. Schematic representation of a mixed matrix membrane (MMM) [132].
A variety of inorganic fillers, including zeolites, porous carbon, and nonporous silica,
have been used to fabricate such type of “mixed” materials over the last two decades.
The resultant inorganic-organic polymer composite membrane, also known as mixed
matrix membrane (MMM), has shown high potential for superior gas separation
performance [132, 187]. The mixed matrix membranes (MMMs) consisting of organic
polymers and inorganic particles, as shown schematically in Scheme 2.8. The bulk
phase (phase A) is typically a polymer matrix. The dispersed phase (phase B) represents
the inorganic fillers, which may be zeolites [17, 175, 176], carbon molecular sieves
[166-169], activated carbons [179], non-porous silica [170-174], C60 [177], and graphite
[178]. To date, more research has shown that MMMs have the potential to achieve
higher selectivity and permeability than the existing polymeric membranes, resulting
from the addition of inorganic fillers. Some of them even have exceeded the Robeson’s
“upper bond” limit [132]. Another significant improvement from MMMs is that the
inherent fragility of the inorganic membranes, which limits their practical applications,
can be avoided as flexible polymers are the continuous phase [15]. It is apparent that
MMMs have potential for practical applications, thus some of research has been
devoted to develop the composite membranes with different configurations. As early as
in Kulprathipanja’s study, the flat sheet silicalite-cellulose acetate MMMs were
prepared by following the steps: evaporating solvent to form a initial MMM, immersing
the MMMs in an ice-water bath, treating membrane at 90 ºC and drying it by air [175].
Flat active membranes were also developed by forming a coating on porous stands or
porous polymeric support [132, 188]. As an effective polymeric membrane
A. Polymer phase
B. Inorganic particle phase
Chapter 2
36
configuration, hollow fiber MMMs have been studied. Bhardwaj et al. reported their
work on MMM hollow fibers for gas separation [189]. Different fillers, e.g. carbon
black, were dispersed in a polysulfone spinning solution to produce highly selective
membranes in the form of single-layer hollow fibers with good mechanical strength.
There are some requirements for successful large-scale fabrication of MMMs. For
example, sub-microsized particles were found to be better than micro-size ones, since
they can be fit well inside the ultra-thin skin layer on porous substrates. It is important
to avoid polymer/particle interface defects, and this can be achieved by modification of
the dispersed particles to increase surface hydrophobicity [132]. In Section 2.6.2, it is
mentioned that only dense and flat membranes are selected in this thesis to characterize
the nanocomposite membranes. Therefore, the fabrication process of dense and flat
sheet MMMs here is similar to that of polymeric membranes, which is called “solution
casting” method. Normally, the homogeneous mixture of polymers, inorganic fillers and
solvents is prepared and then casted on a smooth plate, followed by evaporating solvent
and annealing the membranes at elevated temperatures to remove the residual solvent
[132].
Until now, the many attempts to develop the composite membranes with improved
hydrogen permeability and selectivity have been reported. Nonporous fillers such as
silica nanoparticles have been incorporated into polymer matrix to yield silica-organic
polymer composite membranes. Joly et al. fabricated silica-polyimide (poly(4,4-
oxydiphenylene pyromellitimide) by adding silica source (tetramethoxysilane-TMOS)
to the polymeric acid solution. The addition of silane not only induced the formation of
silica particles in the organic matrix after elevating temperature, but also made a
significant change in the imidization degree of polymer phase and the morphological
modifications in the organic-inorganic interphases. They further investigated the
transport properties of a series of gases, including N2, CO2, H2, etc, by comparing
MMMs with plain polyimide membranes and inorganic silica membranes. The
selectivity of hydrogen over nitrogen for the composite membrane was 69.2. It was
almost 15-fold and 3-fold higher than that of plain polyimide and pure microporous
silica membrane, respectively. The observed hydrogen permeability for MMM was 9.0
Barrer, compared with 4.4 Barrer for polyimide H2 permeability [190]. Differently,
Merkel et al. physically dispersed fumed silica nanoparticles (~13 nm) in poly(4-
methyl-2-pentyne) (PMP) to form nanocomposite membranes. They found that the
Chapter 2
37
produced MMMs exhibited significantly enhanced membrane permeability and
selectivity for large organic molecules over small permanent gases. For example, the n-
butane permeability for the MMMs with adding 30 wt% silica was increased by a factor
of 3 higher than that of the pure PMP at 25 °C. The n-butane/CH4 selectivity for
composite membrane was doubled, compared with that of the pure PMP. This was
because physical dispersion of non-porous nanoparticles yielded polymer-particle
interfaces, disrupted polymer chain packing and thus affected molecular transport [171].
Cong et al. reported a similar work that the nanocomposite membranes were prepared
by using pure silica nanoparticles (~10 nm) and trimethylsilyl or triphenylsilyl-modified
silica nanoparticles. The composite membranes with modified silica particles had lower
gas permeability compared to the membranes with the unmodified silica [191]. Cong et
al. attributed this to the poor compatibility of the silica surface and the polymer, where
the polymer chains did not tightly contact the silica nanoparticle, thus forming a narrow
interface gap surrounding the silica particles.
Zeolite is another inorganic filler, and has been widely investigated in MMMs.
Theoretically, the incorporation of zeolites into polymer matrix can improve both of the
selectivity and permeability, compared with plain polymeric membranes, because of the
molecular sieving effect from zeolite particles. Compared with plain polymer, some
research has indicated that the selectivity decreased or remained the same with
increasing gas permeability, when zeolite particles were added. They attributed this to
the fact that the zeolite was less selective and more permeable than the polymeric phase
by providing low energy pathways for the movement of gas molecules [192].
The first investigation of MMMs for gas separation was reported in 1970s by Paul and
Kemp, who added 5A zeolite into rubbery polydimethyl siloxane (PDMS) [16]. Their
results showed that the addition of 5A zeolite into the polymer matrix caused a very
large increase of the diffusion time lag, but had only minor effects on the steady-state
permeation. After their work, Kulprathipanja’s group reported that that mixed matrix
membranes (MMMs) systems yielded superior separation performance to that of pure
polymeric system [175]. They observed an enhanced O2/N2 selectivity from 3.0 to 4.3
with increasing silicalite content in cellulose acetate (CA) matrix. By using silicalite/CA
MMMs for CO2/H2 separation, a feed mixture of 50/50 (mol%) CO2/H2 with a
differential pressure of 50 psi was used and the separation factor for CO2/H2 was
Chapter 2
38
calculated to be 5.15±2.2 [175]. This value was found to be much higher than that of
pure CA membranes, indicating that the presence of silicalites in the polymer phase can
efficiently improve the gas separation performance of polymeric membranes.
Most of the previous studies on inorganic filler-polymer MMMs use large particles,
with sizes in the micron range, for gas (e.g. hydrogen) separation. For instance, Şen et
al. developed polycarbonate-matrix membranes filled with highly crystalline zeolite-4A
with particles size of 3 μm [193]. At a zeolite loading of 30 wt%, the composite
membranes had an improved H2/N2 selectivity of 73.2 compared with 56.7 for plain
polycarbonate. However, they also found a decrease in hydrogen permeability, which
they attributed to increased rigidity of the polymer chains in the presence of the zeolite
particles [194, 195], the partial blockage of the zeolite pore by the polymer chains [195]
and/or the extended diffusion pathways of the hydrogen molecules through the
membrane [196, 197]. A similar trend was reported by Li et al. [196] and Huang et al.
[195]. Li et al. demonstrated that membranes of polyethersulfone and zeolite 5A (1-5
μm) exhibited about 25% higher H2/N2 selectivity than a plain polyethersulfone
membrane, but had a decrease in gas permeability of at least 25% [196]. Huang et al.
prepared their composite membranes by incorporating 20 wt% of micrometer-sized (1-5
μm) or nano-sized (50-140 nm) zeolite A in polyethersulfone (PES) [195]; the
hydrogen permeability of the PES membrane dropped from 8.96 Barrers to 8.3 Barrers
when filled with nano-zeolite, and further down to 4.94 Barrers with micro-zeolite.
Interestingly, the gas permselectivity enhancement was much more pronounced when
zeolite-4A nanocrystals were incorporated in a PES membrane. Indeed, nano-sized
zeolite are required for fabricating composite membranes because the polymeric
membranes are usually shaped into asymmetric hollow fibers or flat sheets with a thin
selective layer (e.g., <1 μm) for practical applications [15]. Moreover, nanoparticles
would be more suitable for the industrial fabrication of MMMs hollow fibers. This is
because there would be a fouling occurring when large zeolite particles are added into
polymeric solution, which then passes through the nozzle to form MMMs hollow fibers.
By using zeolite nanocrystals, this problem may possibly be minimized rather than
micro-size zeolites. Thus, zeolite nanocrystals are suggested to be more applicable in
zeolite-polymer MMMs.
Chapter 2
39
Golemme et al combined up to 40.2 wt% silicalite-1 (MFI) nanocrystals (80 nm) with
Telfon AF 1600 polymers [198]. The mixed matrix membranes (MMMs) had a
hydrogen permeability of 3580 Barrer, a 15-fold increase relative to the pure polymer
membranes. However, the H2/N2 selectivity of the composite membranes, at just 4.6,
was 50% less than the plain Telfon AF 1600 film. According to the study by Moore and
Koros [194], this result was due to interfacial voids between the zeolite and polymer,
which were probably formed because of low adhesion between the polymer matrix and
the zeolite crystals [193, 199-202]. The formed interfacial voids may have different
effects on the gas separation performance of the MMMs, which will be further
discussed in Section 2.6.4. Here, in Gloemme’s research, it was found that there was an
abrupt decrease of selectivity with increasing permeability. Several approaches have
been proposed to fabricate the mixed matrix membranes (MMMs) that are free of voids
and have enhanced selectivity. Previous research suggests that one of the most effective
ways is surface modification of the zeolite particles with silane-coupling agents [192,
194, 199, 201], which will be further discussed at section 2.7.
In my research, a type of zeolite nanocrystals, sodalite, is considered to be used as
inorganic filler. This type of zeolite has small pores, only permeating small molecules
(e.g. water or He) [84, 203, 204]. To date, there has been some research into making use
of pure sodalite or hydroxy-sodalite membranes for hydrogen separation [185, 186], but
it is still challenging to make defect-free pure zeolitic membranes, especially for large-
scale production. So far, no study has been conducted on the combination of sodalite
nanocrystals and polymeric membranes for this purpose of developing MMMs.
Therefore, one of the objectives in my research is to develop sodalite-polymer MMMs
and investigate their performance for hydrogen gas separation.
2.6.4 Gas transport through membranes
Both porous and dense membranes can be used as selective gas separation barriers;
Scheme 2.9 illustrates the mechanism of gas permeation. When the pore sizes of
membranes range from 0.1 μm to 10 μm, gas molecules travel through membranes by
convective flow. If the pore sizes of membranes are smaller than 0.1 μm, gas transport
is governed by Knudsen diffusion. Finally, if the membrane pores are extremely small,
of only 5–20 Å, then gases are separated by molecular sieving effect [138].
Chapter 2
40
Scheme 2.9. Mechanism for permeation of gases through porous and dense gas
separation membranes [138, 205].
In a gas separation process, components are separated from their mixtures by
differential permeation through dense polymeric membranes, which are governed by
solution-diffusion mechanism as shown in Scheme 2.9. The first person to use the term
“solution-diffusion mechanism” was Graham in 1866 [206]. He postulated that the
penetrant left the external phase by dissolving in the membrane. It then underwent
molecular diffusion in the membrane, driven towards the downstream by a
concentration or pressure gradient, after which it evaporated again in the external phase.
According to the solution-diffusion model, the permeation of molecules through
membranes is controlled by two major parameters: diffusivity coefficient (kinetic
parameter) (D) and solubility coefficient (thermodynamic factor) (S). The diffusivity is
a measurement of the mobility of individual molecule passing through the voids among
the polymeric chains in membrane materials. The solubility coefficient equals the ratio
of the dissolved penetrant concentration in the upstream side of the polymers to the
upstream penetrant partial pressure. The permeability (Pi) representing the ability of
molecules to pass through a membrane is defined in Equation 2-2 [207], by the ratio
between the flux J of the permeant species and its concentration gradient ∆pi over the
membrane thickness d:
Chapter 2
41
dp
PJ iiΔ
=…………………………Equation 2-2
Permeabilities are customarily given in Barrers, where 1 Barrer = 1×10-10
(cm3(STP).cm.cm-2.s-1.cmHg-1) = 3.35 ×10-16 (mol.m .m-2.s-1.Pa-1).
As shown in Equation 2-3, the permeability can be alternatively written by the product
of diffusion coefficient D and solubility coefficient S:
iii SDP ×= …………………………Equation 2-3
A second important relation is shown in Equation 2-4, where the ideal selectivity of the
membrane for component i relative to j is expressed as the ratio of the pure gas
permeabilities of the two penetrants in the membrane materials,
j
i
j
i
j
iij S
SDD
PP
×==α…………………..Equation 2-4
This factor, selectivity, provides a good measurement of the ability of a given polymeric
material to provide a permselective barrier to i relative to j. Also, from the Equation 2-4,
it can be seen that the difference in permeability is caused not only by the diffusivity
(mobility) difference of various gas species, but also by the difference in
physicochemical interactions of these species with the polymer which determine the
amount of gas that can be accommodated per unit volume of the polymer matrix [132].
Therefore, the balance between the solubility selectivity and the diffusivity selectivity
determines the selective transport of gas component through membranes.
The apparent activation energy is a factor defining the amount of energy required to
develop molecular mobility in the polymer chains. This molecular mobility includes
rotational carbon-carbon bonds, segmental chain bonds motion and intermolecular
separations between polymer chains (rises in the free volume) [208]. The apparent
Chapter 2
42
activation energy Ep is normally analyzed according to the Arrhenius equation based on
gas permeability [209-212],
⎟⎟⎠
⎞⎜⎜⎝
⎛ −=
RTE
PP pexp0
……………….Equation 2-5
where P0 the pre-exponential factor, R the ideal gas constant and T is the temperature
(T).
As discussed in Section 2.6.3, inorganic materials include porous and non-porous ones.
Silica is a characteristic non-porous material, which can be controlled to prepare as
ultra- (pore size <0.7 nm) and super-microporous (pore size 0.7-2 nm) amorphous thin
layer and thus used for molecular sieving application, which is similar to zeolite [213].
Zeolites are crystallized solids with structural ultramicroporosity. As microporous
materials, zeolites can be used to achieve gas separation because of their molecular
sieving effect. If the components of a gas mixture are smaller/larger than the pore size
of the zeolite membrane, these components can either pass through the membranes or
be retained by zeolites. Despite this concept is relatively simple, it is difficult to find
exact examples for this exclusion mechanism in the literature because of the defects
existing in zeolitic membranes [214]. A “perfect” membrane without defects is actually
difficult to achieve in practice. The detailed separation ability of a microporous
membrane can be described by the interplay between the mixture adsorption and
diffusion, which is suggested to be similar to the solubility-diffusivity model established
for describing the permeation behavior of polymeric membranes [214]. The forecast of
the separation ability of a given zeolitic membrane is principally possible on the basis of
separately measured mixture adsorption and mixture diffusion data. In 1994, Kapteijn et
al. reported their work on Maxwell-Stefan diffusion model on the permeation flux of n-
butane through a silicalite-1 (MFI) membrane as varying at different feed pressures and
temperatures [215]. Their results showed that the experimental data were excellently
described by a Maxwell-Stefan diffusion model, which was further developed and
universally adopted by other researchers [214-218].
Chapter 2
43
With the addition of inorganic filler, either porous (e.g. zeolite and carbon molecular
sieve) or non-porous (e.g. silica), the gas transport mechanisms in nanocomposite
membranes differs to the gas transport in plain polymeric membranes or in inorganic
membranes. To date, several mechanisms have been reported, including Maxwell’s
model, free-volume increase mechanism, and nano-gap hypothesis mechanism and so
on.
Maxwell’ model is an example of frequently used theoretical expressions to describe
transport behavior in composite polymer systems. Maxwell’s model was firstly
developed to analyze the steady-state dielectric properties of a diluted suspension of
spheres [219], which was further developed by Bruggemann [220], Higuchi and
Higuchi [221], and Davis [222], etc. This effective steady-state permeability of a
material can be given by the following expression,
fPC PP
φφ5.01
1 f
+−
=…………….Equation 2-6
where Pc and Pp are the permeability of the nanocomposite and the pure polymer matrix,
respectively, and Φf is the volume fraction of the nanofiller.
From Equation 2-6, it is clear that there is a loss of membrane permeability with
increasing volume fraction of the incorporated inorganic particle. Maxwell’s model was
suggested to describe permeability in membranes filled with roughly spherical
impermeable particles [223]. However, a different conclusion from Maxwell’s
prediction was obtained in experimental work. For instance, Merck et al. who prepared
fumed silica-poly(4-methyl-2-pentyne) nanocomposite membranes for gas separation.
By adding 30 wt% fumed silica, the silica nanocomposite permeability was 1.4-fold
greater than that of plain polymer, whereas Maxwell's equation predicted 35% reduction
in the permeability at the same filler loading [171]. Similar non-Maxwell effect also has
been investigated in other research [224-226]. Studies attributed this problem to the
disadvantage of Maxwell’s model, which neglected the interactions between nano-
fillers and polymer chains, and the relation between nano-fillers and penetrants [223]. In
Chapter 2
44
some nanocomposite membranes, these types of interactions are strong, and
significantly change the diffusivity and solubility of penetrants.
Scheme 2.10 The nano-gaps were hypothesized to exist in the silica-BPPOdp
nanocomposite membranes by Cong et al. [191].
As mentioned in Section 2.6.2, molecular diffusion through a dense polymer membrane
strongly depends on the amount of free volume existing in a material. The free volume
mechanism indicates that an increase in polymer free volumes is expected to enhance
penetrant diffusion. Moreover, free volume also affects gas solubility in a polymer, but
having much less extent than its influence on diffusion, which increases slightly with
increasing polymer free volume. Thus, permeability of a composite membrane, which is
governed by gas solubility and diffusivity, increases with polymer free volume in a
manner similar to that of diffusivity. The addition of inorganic fillers may affect the
polymer chains, thus increasing the free volume between polymer chains, and finally
enhancing the permeability of composite membranes [173, 223]. This mechanism is
consistent with a number of reported experimental works [173, 226-228].
Cong et al. found that the nanocomposite membranes fabricated from the modified
silica nanoparticles with trimethylsilyl and triphenylsilyl organic functional groups had
lower gas permeability than the membranes prepared by using the unmodified silica
[191]. They proposed a “nano-gap” hypothesis. The lower permeability of gases was
attributed to the poor compatibility of the silica surface and polymers, where the
polymer chains did not tightly contact the silica nanoparticle, thus a narrow gap
Chapter 2
45
surrounding the silica particles (as shown in Scheme 2.10). The permeability was
increased due to the short diffusion path, but the nano-gap had surprisingly no effect on
the gas selectively.
Koros et al. also reported these interfacial gaps/voids in their work by studying zeolite
4A-Udel polymer composite membranes [194]. They found that there seems to be a
much more complicated condition for the composite membranes with incorporated
zeolite crystals or nanocrystals than those with silica (non-porous) particles. It has been
known that the channels and pores existing in zeolites may provide low-energy
pathways for gas molecules, in zeolite-polymer composite membranes as shown in
Scheme 2.11. For the gas with molecular sizes larger than the pores of zeolite, gas
permeation will be hindered [138]. However, the actual gas transport theory may not be
so optimal, on which Koros’s group has conducted a systematic study.
Scheme 2.11. Gas permeation through mixed-matrix membranes containing dispersed
zeolite particles [138].
In Scheme 2.12, there several possible conditions summarized when the inorganic
particles, in particular, zeolite, exist in polymer matrix, including voids or high-free
volume phase (case II and case III), a rigidified or compressed region of the polymer
matrix (case I), a reduced permeability region in the outer layer of the zeolite (case IV
and case V). Case I represents a rigidified region in the polymer phase externally
surrounds zeolite, showing reducing permeability. Both cases II and III have voids at
Chapter 2
46
the inorganic-organic interfaces. However, case III has a thinner effective void than
case II, which is on the order of the sizes of the gas penetrants. In case IV, the region on
the zeolite surface completely prevents the gas penetrants from entering the zeolite. In
case V, the gas penetrants enter the channels of zeolite, but at a slower rate than normal
one when passing through pure zeolite. From the predicted O2 permeability-O2/N2
selectivity graph displayed in Scheme 2.12, there would be an increase of O2
permeability and selectivity in case II and III, whilst there may be an increase of
selectivity for oxygen over nitrogen with decreasing oxygen permeability for the other
cases [194].
Scheme 2.12. Summary of the relationship between MMMs morphologies and transport
properties. Solid circles represent calculated values for MMMs with an incorporation of
35 vol% zeolite 4A and Ultem as polymer matrix. (Modified from Ref. [194])
Chapter 2
47
2.7 Organic functionalization of zeolite nanocrystals and membrane fabrication
Zeolites can be functionalized with organic groups in order to modify the surface
properties of these materials. Many different applications are envisioned for these
resulting organic-inorganic hybrid materials including catalysis, environmental
protection and sensors [24-26]. Up to now, a number of studies have been focused on
the application of organic-functionalized zeolites to form zeolite-polymer mixed matrix
membranes (MMMs).
It is known that zeolite-polymer MMMs are a combination of inorganic zeolites and
organic polymeric membranes. The key challenge for the preparation of these hybrid
materials is to avoid phase separation between the organic and inorganic moieties [229].
It has been suggested this phase separation can be improved by incorporating organic-
functionalized zeolite instead of un-modified ones in polymer matrix.
Presently, the organic groups which have been used for zeolite functionalization can be
divided into two main groups, unreactive and reactive ones [6, 21-26]. Methyl moiety is
an example of unreactive organic ones. Some research has pointed out that the
attachment of methyl groups can improve the hydrophobicity of zeolites that are
relatively more hydrophilic as compared to polymeric membranes [22]. This
modification of crystal surface helps reduce the phase separation between organic or
inorganic phases when combining zeolite and polymers.
However, recent studies have found that covalent bonding more efficiently prevents
phase separation between zeolite and polymer matrix. Therefore, the effort is made to
the organic functionalization with reactive groups, which can add surface reactivity to
zeolite crystals and make the rigid inorganic particles react with the selected polymer
precursors. Some common functional groups attached on the surface of inorganic
particles, includes amino, vinyl, acryl, hydroxyl, and carboxyl [26, 230-233], can
produce radicals, cations or anions through high-energy radiation, plasma or other
means to graft to polymer matrix.
Chapter 2
48
Scheme 2.13. Principle of coupling of organofunctional silanes onto zeolite surface
[192].
In light of preparing polyimide MMMs for gas separation, the organic functionalization
of amino groups onto zeolites can be effectively used to form covalent bonds between
zeolites and polyimide. As shown in Scheme 2.13, Duval et al. promoted the adhesion
between zeolite particles and polymer matrices by modifying the zeolite surfaces with
silane-coupling agents (e.g. γ-aminopropyltriethoxy silane, N-p-(aminoethy1)-y-
aminopropyltrimethoxy silane and styryl amine functional silane) [192]. After
incorporating organic-functionalized zeolites, it was found that amino functional silanes
were efficient in improving the adhesion between organic and inorganic phases. Pechar
et al. developed mixed matrix membranes (MMMs) from polyimide and zeolite L or
ZSM-2 zeolite, which were functionalized by APTES (aminopropyl-triethoxysilane)
coupling agents (Scheme 2.14) [231, 232]. The prepared mixed matrix membranes
composed of 6FDA-6FpDA-DABA and amine-functionalized zeolite L or ZSM-2 were
fabricated without interfacial defects. Similar findings have also been reported in other
research [196, 230].
Chapter 2
49
Scheme 2.14. Reaction between an amine-functionalized zeolite and the polyimide to
create a covalent amide linkage during annealing [231, 232].
Because of its importance, there is an interest in the development of novel methods for
organic functionalization of zeolite, especially with amino organic groups. As
mentioned in section 2.4, hydrothermal synthesis with organic structure-directing agents
(SDA) is a commonly used method to produce zeolites or zeolite nanocrystals. Some
organosilanes, such as those with methyl groups, can be directly mixed and then react
with the synthesis gel to form the zeolites attached with organic moieties [22]. However,
usually an organic SDA remains in intracrystalline voids, and has to be removed to open
the zeolite micropores. Unfortunately, the template removal, normally done through
high-temperature calcination, has proven unsuitable for colloidal nanocrystals because it
leads to significantly irreversible aggregation [11, 75]. Furthermore, some active
organic groups (e.g. amino moiety) can not withstand high-temperature calcination [23].
Hence, Smaihi et al. developed a novel method with two consecutive grafting
procedures to prepare organic-functionalized zeolite nanocrystals [23]. Firstly, organic
ligands were grafted at template-containing zeolite beta nanocrystals inhibited
irreversible aggregation during the combustion of SDA. After removal of organic
templates at 550 °C, a re-grafting step was required to obtain colloidal functionalized
Chapter 2
50
zeolite nanocrystal suspension, with organic ligands attached on zeolitic surface silicon
or aluminum atoms. However, the repeated addition of organosilanes is costly.
Thus, an alternative synthetic method for zeolite nanocrystals attached with organic
groups, in particular reactive ones, is highly desirable. This process requires avoiding
high-temperature template removal, preserving good dispersion of formed nanoparticles
and efficiently organic-functionalizing nanoparticles.
As discussed in section 2.4, Yao et al. developed a novel method to prepare sodalite
nanocrystals by direct transforming silicalite nanocrystals [84]. It is suggested to be
useful for synthesizing organic-functionalized sodalite nanoparticles since the whole
preparation process can be undertaken at relatively low temperatures (80―90°C).
Furthermore, the prepared SDA-free sodalite can be readily redispersed, which should
help to form homogeneous MMMs for gas separation application.
2.8 Summary and Aims The review of relevant research has highlighted the need to investigate the effect of
polymer hydrogels on zeolite crystallization and growth. Currently, there has been no
study on the formation of zeolites in either crosslinked or uncrosslinked chitosan
hydrogel systems. The potential application of zeolite nanocrystals to mixed matrix
membranes (MMMs) for gas separation has been discussed in the literature review. No
relevant studies have been conducted on the combination of sodalite nanocrystals with
polymer matrix. To fabricate defect-free MMMs, organic-funtionalization of zeolite
nanocrystals is suggested as an effective way. Therefore, one of the main aims in this
study is to develop novel methods for synthesizing, functionalizing zeolite nanocrystals,
and fabricating zeolite-polymer MMMs for gas separation.
The specific aims of my PhD research are listed as follows:
o To investigate zeolite crystallization in chitosan hydrogels:
Chapter 2
51
Zeolite nanocrystals with controllable particle sizes, such as zeolite A and
Y, will be synthesized by using glutaraldehyde-crosslinked chitosan gel
(GA-CS);
The effects of crosslinking of chitosan polymer on the zeolite nucleation
and crystallization will be studied. This will likely lead to new approaches
to control the synthesis of zeolite with different morphologies and crystal
sizes.
o To develop sodalite-polymer mixed matrix membranes (MMMs) for gas
separation application:
Sodalite nanocrystals with organic functional groups (reactive or
unreactive ones) will be synthesized and characterized;
Sodalite-polyimide MMMs will be prepared by applying organic-
functionalized sodalite nanocrystals. The separation property of the
sodalite-polyimide MMMs will be determined, and their gas transport
mechanism will be discussed.
Chapter 3
52
Chapter 3 Growth of Zeolite in Chitosan Hydrogels
3.1 Overview Chapter 3 presents the investigation on zeolite nucleation, crystallization and growth in
chitosan hydrogels. Section 3.2 includes the results about the zeolite crystallization in
glutaraldehyde-crosslinked chitosan (GA-CS) hydrogels for the purpose of controlling
crystal size and synthesizing zeolite nanocrystals. The effect of the synthesis parameters
(e.g. the amounts of silica, chitosan, and glutaraldehyde and aging and heating times) on
the size, size distribution and crystallinity of the particles are systematically studied in
this section. Section 3.3 reports the effect of uncrosslinked chitosan hydrogels on the
nucleation and formation of zeolite. The zeolite A particles consisting of a thin
crystalline shell and anamorphous core has been proven to form within chitosan
hydrogels. Our results indicate that the formation of cube-like or rectangular core-shell
structures involves particle aggregation and surface-to-core crystallization, which may
be induced by chitosan polymer networks.
3.2 Zeolite crystallization in crosslinked chitosan hydrogels
3.2.1 Experimental
3.2.1.1 Synthesis of zeolite LTA (NaA)
Table 3.2 summarizes the experimental design for the synthesis of LTA (NaA) in
crosslinked chitosan gel. Firstly, 0.6-1.5 g of chitosan (average molecular weight
120,000 g/mol, ~80% deacetylation, Sigma-Aldrich, denoted CS) was dissolved in 21 g
of 1 M acetic acid (Sigma-Aldrich). The resulting solution was stirred at room
temperature for 1 h and then left overnight without stirring, after which 2.0-4.0 g of
colloidal silica (HS-30, 30%, Sigma-Aldrich) were added. A given amount (0.6-5.0 g)
Chapter 3
53
of glutaraldehyde (50%, Sigma-Aldrich, denoted GA) was added into the CS-silica
solution, and left undisturbed at room temperature for 2 h, resulting in crosslinked
chitosan hydrogel (denoted GA-CS). Therefore, the silica-filled crosslinked chitosan
(GA-CS) hydrogels were synthesized from molar compositions in the range 0.005-
0.0125CS:10-20SiO2:1.88-25 glutaraldehyde (GA):21acetic acid (HAc):1243-2499H2O,
corresponding to mass compositions of 0.6-1.5CS:0.6-1.2SiO2:0.4-5.0GA:1.3HAc:22.4
-24.7H2O. Secondly, an alkaline solution was prepared by dissolving 5.56 g of NaOH
(99%, Merck) in 20.00 g of deionized water, with subsequent addition of 2.45 g of
NaAlO2 (anhydrous, Sigma-Aldrich) during stirring. The molar composition of the
alkaline solution was 7.7Na2O:1.0Al2O3:111.0H2O. This alkaline solution was
introduced into the silica-filled crosslinked chitosan hydrogel with a final molar
composition of 000.5-0.0125CS:10-20SiO2:1.88-25GA:21HAc:80Na2O:10Al2O3:2396-
2455H2O, and aged for 12-72 h at room temperature. After aging, the gel was removed
from the alkaline solution, transferred to a sealed polypropylene bottle and then heated
at 90 °C for 1, 3 or 6 h to allow zeolite crystallization. To make a comparison, another
sample prepared without aging was heated at 90 °C for 3 h in the presence of the
alkaline solution.
3.2.1.2 Synthesis of zeolite FAU (NaY)
In this case, the synthesis of silica-filled hydrogels was performed from a system with a
molar composition of 0.01CS:17.5SiO2:12.5GA:21HAc:1302H2O. The same chemicals
were used for synthesis of NaY crystals as described earlier for the synthesis of NaA
crystals. Typically, 1.2 g of CS was dissolved into 21 g of 1 M HAc. As with NaA, the
solution was stirred at room temperature for 1 h, and then left overnight, after which 3.5
g of colloidal silica was added, and 2.5 g of GA was added to form GA-CS. The
alkaline solution was prepared as follows: 4.14 g of NaOH was dissolved in 25.83 g of
deionized water, with subsequent addition of 0.75 g of NaAlO2 during stirring. The
molar composition of alkaline solution was 17.3Na2O:1Al2O3:455.6H2O. The solution
was stirred for 0.5-1 h until it became clear and then it was introduced into the
crosslinked chitosan gel system with a molar composition of 0.01CS : 17.5SiO2 :
12.5GA : 21HAc : 55Na2O : 3.18Al2O3 : 2769H2O, and allowed to age at room
temperature for 36 h. After aging, the gel was removed from the alkaline solution,
transferred to a sealed polypropylene bottle, and then heated at 90 °C for 5 h.
Chapter 3
54
Table 3.2.
Exper
iment
al
desig
n for
the
zeolit
e A
synth
esis
in
crossl
inked
chitos
Tabl
e 3.
1. E
xper
imen
tal d
esig
n fo
r the
zeo
lite
A sy
nthe
sis i
n cr
ossl
inke
d ch
itosa
n ge
l.
Chapter 3
55
an gel
3.2.1.3 Removal of crosslinked chitosan hydrogels
The heat-treated gels, which contained zeolite, were repeatedly washed with deionized
water until a pH of less than 8 was attained. Approximately 3 g of hydrogel was stirred
into 150 mL of 10%-H2O2 solution and then heated at 80-90 °C for 1-2 h. The zeolite
crystals were retrieved by high-speed centrifugation and repeated washing with
deionized water; these were dried at 60 °C. For comparison, gels also were calcined to
remove the crosslinked CS. After washing, the zeolite-containing gels were dried at 80
°C overnight, ground by hand using a mortar and pestle, and calcined at 550 °C under
air for 2 h at an initial heating rate of 2 °C.min-1.
3.2.1.4 Characterization
Scanning electron microscopy (SEM) images were taken with a JSM-6300F microscope
(JEOL). The particle size distributions for zeolite crystals were determined by manual
measurement of 300 crystals for each sample from the SEM images with Adobe
Photoshop software. Elemental Si/Al ratios of samples were determined by energy
dispersive X-ray spectroscopy (EDXS) on the JSM-6300F microscope. X-ray
diffraction (XRD) patterns were recorded on a Philips PW1140/90 diffractometer with
Cu Kα radiation (25 mA and 40 kV) at a scan rate of 2 °/min and a step size of 0.02°.
Thermogravimetric analysis (TGA, Perkin Elmer, Pyris 1 analyzer) was performed at a
heating rate of 5 °C/min to 700 °C in oxygen with a flow rate of 15 cm3⋅min-1. Nitrogen
adsorption-desorption experiments were performed at 77 K with a Micrometritics
ASAP 2020MC analyzer. The samples were degassed at 673 K for 24 h, and 623 K for
4 h, respectively prior to analysis, and the specific surface areas were calculated
according to the Brunauer-Emmett-Teller (BET) method. To study the dispersibility of
zeolite nanocrystals, the particle size distributions of colloidal zeolite suspension were
analyzed by light scattering with a Malvern Mastersizer 2000 analyzer. Approximately
12-15 mL samples of colloidal zeolite suspension were prepared for this purpose by
dispersing 50 mg of each sample into 50 mL of deionized water during ultrasonication.
Chapter 3
56
3.2.2 Results and Discussion
Scheme 3.1. Synthesis of zeolite crystals within crosslinked chitosan hydrogels (GA-
CS).
A schematic diagram for the formation of zeolite nanocrystals in GA-CS hydrogels is
shown in Scheme 3.1. A colloidal silica solution is dispersed in the solution of CS and
acetic acid. When the crosslinker (GA) is added, the amino groups from the backbones
of chitosan are crosslinked [234], which causes chitosan solution to solidify into yellow
gels that contain colloidal silica. To form aluminosilicate zeolite gels, the alkaline
solution is added into the yellow gel. During the aging, alkaline solution penetrates into
the GA-CS gel and reacts with the entrapped silica. Aluminosilicate gels crystallize
during the subsequent heating, producing zeolite crystals within the crosslinked
Chapter 3
57
chitosan hydrogel networks. To remove the chitosan hydrogels, hydrogen peroxide is
added to solubilize crosslinked chitosan by degradation of the network structure. Then
the zeolite crystals are readily collected through high-speed centrifugation and repeated
washing (Scheme 3.1).
The possible mechanism for the decomposition of polymer hydrogels is shown in
Figure 3.1 based on some previous studies [114-118]. H2O2 can be easily decomposed
to form the highly reactive hydroxyl radical (HO·), in particular under the heating. The
hydroxyl radical can react with carbohydrates exceedingly rapidly, such as chitosan,
abstracting a C-bonded H atom into soluble small molecules (R1 and R2 in Figure 3.1).
Thus, chitosan can be decomposed under heating hydrogen peroxide, resulting in a
release of polymer-free zeolite nanocrystals.
The degradation of chitosan by the highly reactive hydroxyl radical (HO·) from
hydrogen peroxide:
H2O2 → H+ + HOO-
HOO-→ OH- + (O)
H2O2 + HOO- → OH- + O2- · + H2O
RH (from CS) +HO·→R·+H2O
R·→ R1+ R2
Figure 3.1. The reaction between hydroxyl radical from hydrogen peroxide and
carbohydrates [116-118].
3.2.2.1 Effect of the amount of SiO2
Figure 3.2 shows the XRD patterns of the samples prepared with molar compositions of
0.01CS:ySiO2:12.5GA:21HAc:(1166+8y)H2O (y = 10-20). They are denoted A-10SiO2,
A-12.5SiO2, A-15SiO2, A-17.5SiO2, and A-20SiO2, respectively. From Figure 3.2, A-
10SiO2 appears to be amorphous, and A-12.5SiO2 exhibits very low crystallinity,
whereas A-15SiO2, A-17.5SiO2, and A-20SiO2 are pure zeolite A. Interestingly, A-
17.5SiO2 exhibits the smallest average particle size (148 nm), which is much smaller
Chapter 3
58
than the mean particle sizes of 325 nm and 239 nm of A-20SiO2 and A-15SiO2,
respectively. It is well known that organic additives play an important role in zeolite
nucleation and growth [13, 14]. Previous studies have found that the addition of water-
soluble surfactants (e.g. sodium dodecyl sulfate, sodium dioctylsulfosuccinate and
cetyltrimethylammonium bromide) and organic polymers (e.g. poly(ethylene glycol)) in
the zeolite gel dramatically shortened prenucleation and nucleation periods and
accelerated crystal growth [14]. Some recent mathematical [235, 236] modeling and
experimental results [237] have shown that the zeolite nucleation takes place at the
interface between the solution and the gel solid by adsorption and rearrangement of
soluble precursor. Our experimental results could possibly be explained by the effect of
the ratio of SiO2 to CS on the rates of nucleation and growth. When the ratio of SiO2 to
CS is high at 2000:1, both initial nucleation rate and subsequent growth rate are
presumably high due to the high concentration of aluminosilicate gel, ultimately leading
to the larger crystal sizes. As the ratio of SiO2 to CS is lowered to 1750:1, the initial
nucleation rate might drop slightly, accounting for far smaller crystals. When the ratio
of SiO2 to CS decreases further to 1250:1, both nucleation rate and growth rate would
significantly decrease, resulting in decrease in the number of nuclei formed in the
system, but an overall increase in final particle size. As expected, if the ratio of SiO2 to
CS becomes too low (e.g. 1000), the amount of silica is insufficient for zeolite
crystallization.
The SEM image, particle size distributions, and nitrogen sorption isotherm of A-
17.5SiO2 are shown in Figure 3.3. The SEM image (Figure 3.3a) reveals that the
crystals exhibit irregular shapes. The particle sizes determined by SEM range from 100
nm to 200 nm, averaging 148 nm (Figure 3.3b), which is slightly smaller than their
mean particle size 149 nm measured by light scattering (Figure 3.3c). The similarity of
these size distributions and mean sizes suggests that the produced particles were well
dispersed in water. Furthermore, the suspension formed by dispersing the zeolite
particles in water is stable under lab condition for at least one week. In the nitrogen
adsorption-desorption isotherm (Figure 3.3d), the amount of nitrogen adsorbed in the
sample is very low at low relative pressures, and substantially increase at high relative
pressures (e.g. P/Po>0.8). It is clear that the nitrogen adsorption arises from the external
surfaces of nanocrystals because the micropores of well-crystallized zeolite LTA
crystals are inaccessible to nitrogen molecules at the liquid nitrogen temperature (77K)
Chapter 3
59
[238]. The BET surface area is calculated to be 26.2 m2/g, and this further supports a
high crystallinity of the sample.
10 20 30 40 50 60
(e)
(d)
(c)
(b)
(a)
In
tens
ity (a
.u.)
2θ (degrees)
Figure 3.2. XRD patterns of the samples prepared with molar compositions of
0.01CS:ySiO2:12.5GA:21HAc:(1166+8y)H2O (y = 10-20) under the same aging (36 h)
and heating (90 °C for 3h) conditions: (a) A-10SiO2; (b) A-12.5SiO2; (c) A-15SiO2; (d)
A-17.5SiO2 and (e) A-20SiO2. All of the samples were collected after H2O2 treatment.
Figure 3.3. (a) SEM image, (b) particle size distribution determined by SEM, (c)
particle size distribution measured by light scattering, and (d) N2 sorption isotherm of
the sample A-17.5SiO2.
Chapter 3
60
3.2.2.2 Effect of the amount of chitosan
10 20 30 40 50 60
2θ (degrees)
(d)
(c)
(b)
Inte
nsity
(a.u
.)
(a)
Figure 3.4. XRD patterns of the samples prepared with molar compositions of
xCS:17.5SiO2:1250xGA:21HAc:(1233+6944x)H2O (x = 0.005-0.0125) under the same
aging (36 h) and heating (90 °C for 3 h) conditions: (a) A-0.0125CS; (b) A-0.01CS; (c)
A-0.0075CS; (d) A-0.005CS. All of the samples were collected after H2O2 treatment.
Figure 3.4 shows the XRD patterns of the samples produced with molar compositions of
xCS:17.5SiO2:1250GA:21HAc:(1233+6944x)H2O (x = 0.005-0.0125), and clearly the
crystallinity of samples declines as the amount of chitosan increases. For instance, when
x = 0.0125 (Figure 3.4a), the sample (A-0.0125CS) exhibits very low crystallinity. This
is probably due to an increase in the density of the polymer networks with the CS
concentration, resulting in significant reduction of diffusion of zeolite precursors, and
hence slows crystallization. The same argument can be applied to the difference in
particle size. The SEM results indicate that the crystal size of zeolite A decreases with
increasing amounts of CS; average sizes are 399 nm, 341 nm and 148 nm for samples
prepared with x = 0.0050, 0.0075, and 0.01, respectively. Again this can be explained
by the decrease in local concentration of aluminosilicate available within the chitosan
hydrogels as the amount of chitosan increases.
Chapter 3
61
3.2.2.3 Effect of the amount of glutaraldehyde (GA)
The amount of GA was varied in the CS hydrogels to study the effect of crosslinking.
The samples were prepared from crosslinked CS hydrogels with a molar composition of
0.01CS:17.5SiO2:zGA:21HAc:(1233+6z)H2O (z = 1.88-25) under the same aging (36 h)
and heating (90 °C for 3 h) conditions, and H2O2 treatment. The GA concentration in
the crosslinking gel was expressed as the molar ratio “GA molecules: amino groups
from chitosan” (GA:NH2), which was varied at 0.3GA:1.0NH2 (z = 1.88),
1.0GA:1.0NH2 (z = 6.25), 2.0GA:1.0NH2 (z = 12.5), 4.0GA:1.0NH2. (z = 25). The
samples obtained were denoted A-0.3GA, A-1.0GA, A-2.0GA and A-4.0GA
respectively. Figure 3.5 displays the SEM images for A-0.3GA, A-1.0GA, A-2.0GA
and A-4.0GA. Figure 3.6 clearly shows that A-0.3GA has a wide size distribution
ranging from 150 nm to over 500 nm (Figure 3.5a), whereas A-1.0GA exhibits a
narrower size distribution ranging from 100 nm to approximately 340 nm (Figure 3.5b).
A-2.0GA and A-4.0GA exhibit still smaller sizes and narrower size distributions
(Figure 3.5c, d).
Figure 3.5. SEM images of the particles produced with different amounts of added GA:
(a) A-0.3GA; (b) A-1.0GA; (c) A-2.0GA; (d) A-4.0GA.
Chapter 3
62
100 200 300 400 5000
10
20
30
40
50
A-0.3GA A-1.0GA
Rel
ativ
e Fr
eque
ncy
(%)
Particle Size (nm)
(a)
Figure 3.6. Particle size distributions determined by SEM images for A-0.3GA and A-
1.0GA.
Figure 3.7. Particle size distributions determined by SEM (a and c) and by light
scattering (b and d) for A-2.0GA (a and b) and A-4.0GA (c and d).
Chapter 3
63
Figure 3.7 compares the particle size distributions measured by SEM and light
scattering for A-2.0GA and A-4.0GA. In terms of SEM-derived particle size
distributions, both A-2.0GA and A-4.0GA have similar particle size distributions, with
average sizes of 148 nm and 147 nm, respectively (Figure 3.7a and c). Light scattering
measurements indicate that both A-2.0GA and A-4.0GA are well dispersed in deionized
water, and that their mean particle size is 149 nm for A-2.0GA (Figure 3.7b) and 148
nm for A-4.0GA (Figure 3.7d).
These results suggest that the degree of crosslinking of chitosan hydrogel has a strong
effect on the crystal size. As the amount of GA increases, the three-dimensional 3D
networks of the hydrogels become denser and more uniform [107]. The higher degree of
crosslinking leads to lower degree of swelling, thereby decreasing the diffusion of
zeolite precursors in solution [112]. Therefore, crosslinked chitosan gels with more GA
provide more rigid, confined-spaces for zeolite crystallization and less material for
crystallization, resulting in smaller and more uniform zeolite nanocrystals. A-2GA and
A-4GA exhibit similar morphologies and particle size distributions because, as has been
pointed out in the literature [112], there is little change in the degree of crosslinking at
high concentrations of GA and, therefore, little change in the crystallization
environment.
3.2.2.4 Effect of aging time
The crosslinked chitosan gels with a molar ratio of 0.01CS:17.5SiO2:12.5GA:21HAc:
1302H2O were aged for different periods of 12 h, 36 h or 72 h, and then heated at 90 °C
for 3 h. The resulting samples were denoted A-12h, A-36h and A-72h, respectively.
Figure 3.8 shows the XRD patterns of A-12h, A-36h and A-72h. A-12h possesses a
LTA crystal phase with a very low crystallinity (Figure 3.8a). As the aging time
increases, the crystallinity of sample increases (Figure 3.8b and c). This is because the
penetration of the alkaline solution through the crosslinked chitosan hydrogels is
essential for producing aluminosilicate gels via reaction with silica and longer aging
times allow greater penetration. In addition, longer aging allows more uniform
nucleation to occur through the gel matrix, which assists the formation of zeolite
Chapter 3
64
crystals with small sizes and narrow size distributions. The SEM images of A-36h and
A-72h are shown in Figure 3.9.
10 20 30 40 50 60
(a)
(b)
2θ (degrees)
Inte
nsity
(a.u
.) (c)
Figure 3.8. XRD patterns of the samples obtained from the crosslinked chitosan gels
after (a) 12 h aging (A-12h), (b) 36 h aging (A-36h), and (c) 72 h aging (A-72h).
Figure 3.9. SEM images of the particles obtained from the chitosan gels after (a) 36 h
aging (A-36h), and (b) 72 h aging (A-72h).
Figure 3.9a exhibits that the sample A-36h has particle sizes from 100 to 200 nm,
averaging 148 nm as mentioned above. When the aging period is extended to 72 h
(sample A-72h), the resultant NaA particles have larger sizes ranging from
approximately 200 to 350 nm (Figure 3.9b). This is probably due to the swelling
Chapter 3
65
behavior of crosslinked chitosan, which is enhanced during the long exposure time to
alkaline solution [113]. Moreover, it is possible that the crosslinked chitosan may partly
degrade after days of exposure these highly alkaline conditions [112]. As a result, the
crosslinked chitosan can adsorb more precursor Na and Al species and provides larger
spaces for the further growth of zeolite particles.
10 20 30 40 50 60
(a)
Inte
nsity
(a.u
.)
2θ (degrees)
(b)
Figure 3.10. XRD patterns of the crystals in samples (a) A-0h (with alkaline solution)
and (b) A-36h (without alkaline solution).
For further comparison, sample A-0h was made by directly heating a gel after the
addition of alkaline solution without any aging period. Figure 3.10 and Figure 3.11
show the XRD patterns and SEM images, respectively, for samples A-0h and A-36h. A-
0h appears to have a higher degree of crystallinity (as seen in greater peak heights) than
A-36h (Figure 3.10b), as well as the presence of some amorphous material (Figure
3.10a). The SEM images of A-0h (Figure 3.11a) exhibit generally coarser particles with
a broad size distribution—some particles even exceed 500 nm—and a mean particle
size of 380 nm. In contrast, A-36h possesses a far more uniform particle size
distribution (Figure 3.11b) with a smaller average size of 148 nm.
The reasons for the difference in particle sizes can be explained by the extent and
uniformity of penetration of alkaline solutions. After the crosslinking reaction with GA,
the CS hydrogels entrap colloidal silica particles within their networks. To produce
Chapter 3
66
zeolite crystals, an aluminosilicate gel of the desired composition needs to be formed by
diffusion of the alkaline solution throughout the GA crosslinked CS hydrogel network,
and subsequent reaction with colloidal silica. Without aging (i.e., sample A-0h), the
large compositional gradients of aluminosilicate gel exist during heating, such that the
zeolite nucleation and growth can only start from the outside of the polymer hydrogel
surfaces, resulting in non-uniform growth. Thus, during the limited heating time of 3 h,
the poorly distributed precursor material deep within the polymer gel can not fully
crystallize into zeolite, and this unconverted is the amorphous phase found in the XRD
pattern (Figure 3.10a). However, aging at room temperature for 36 h (sample A-36h)
allows the alkaline solution to be evenly distributed throughout the crosslinked
hydrogels, leading to an aluminosilicate gels with a uniform composition. After aging,
the polymer hydrogels that incorporate the aluminosilicate gel are removed from the
solutions for heating, which helps prevent the polymer hydrogels from over-swelling.
On the other hand, during hydrothermal reaction, the wet gels may also slightly shrink
due to water evaporation; presumably this makes confined-space growth more
effectively [239].
Figure 3.11. SEM images of the crystals in samples (a) A-0h and (b) A-36h.
3.2.2.5 Effect of heating time
To study the effect of heating time, the molar composition of the gels was fixed at
0.01CS:17.5SiO2:12.5GA:21HAc:1302H2O, and all the synthesis gels were aged for 36
h and then heated at 90 °C. The heating time was varied from 1 h, 3 h or 6 h, and the
corresponding as-synthesized samples were denoted A-1h, A-3h, and A-6h, respectively.
Chapter 3
67
The XRD patterns (Figure 3.12) indicate that there is no crystalline material formed
after 1 h heating, whereas zeolite A crystals are produced once the heating period is
extended to 3 h (A-3h) or 6 h (A-6h). Figure 3.13 shows the SEM images of A-3h and
A-6h. It can be seen that the zeolite A produced under 6 h hydrothermal treatment has
larger particle sizes than those produced during 3 h of heating. This difference might be
attributed to a combination of the flexibility, interconnected pore channels, and large
pore sizes (e.g. a few microns) of polymer hydrogel networks. Therefore, optimized
hydrothermal conditions are required for controlled synthesis of zeolite nanocrystals
with a narrow size distribution.
10 20 30 40 50 60
(a)
(b)
2θ (degrees)
Inte
nsity
(a.u
.)
(c)
Figure 3.12. XRD patterns of the samples (a) A-1h, (b) A-3h, and (c) A-6h.
Figure 3.13. SEM images of the particles produced after heating for (a) 3 h (A-3h) and
(b) 6 h (A-6h).
Chapter 3
68
3.2.2.6 Comparison between the treatment of H2O2 and
conventional calcination
In this section, a novel method was applied to remove the confining polymer network
by degradation of crosslinked chitosan hydrogels in a hydrogen peroxide solution. H2O2
easily decomposes to form the highly reactive hydroxyl radical (HO·), especially under
heating. The hydroxyl radical attacks polymer hydrogels, degrading the crosslinked
structure and chitosan molecules [114-118]. For comparison, high-temperature
calcination, which is a conventional method for removing organic agents, was also used.
The sample for this comparison was crosslinked chitosan with zeolite A, which was
produced from gels with a molar ratio of 0.01CS:17.5SiO2:12.5GA:21HAc:1302H2O,
which were aged for 36 h in alkaline solution and heated for 3 h at 90 °C. The sample
treated by hydrogen peroxide is denoted as A-H2O2 and that treated by calcination as A-
cal. Figure 3.14 compares the XRD patterns of pure particles treated with H2O2 (Figure
3.14a) and calcination at 550 °C for 2 h (Figure 3.14b). Clearly, the zeolite A crystals
retain greater crystallinity under the hydrogen peroxide treatment than those obtained
from high-temperature calcination.
10 20 30 40 50 60
(a)Inte
nsity
(a.u
.)
2θ (degrees)
(b)
Figure 3.14. XRD patterns of (a) A-H2O2 and (b) A-cal.
Figure 3.15 shows the thermogravimetric (TG) curves of plain crosslinked chitosan, A-
cal and A-H2O2. Under flowing pure oxygen, there is a continuous mass loss from 100
°C to 540 °C for the plain crosslinked chitosan (Figure 3.15a). Its total mass loss
Chapter 3
69
reaches 100% after 540 °C. A-H2O2 (Figure 3.15c) has a total mass loss of
approximately 11% occurring, which is mainly attributed to loss of the structural water
from the zeolite as well as physically adsorbed water. This confirms that the polymer
hydrogels were completely removed by the hydrogen peroxide treatment method. For
A-cal, however, there is another mass loss after 400 °C (Figure 3.15b) in addition to the
loss of adsorbed and structural water at around 100 °C. This suggests that the
crosslinked chitosan was not completely burned off during calcination. Given the
greater crystallinity retained and cleaner removal of the hydrogel, the hydrogen
peroxide treatment is clearly the preferred method for removal of GA-CS after
hydrothermal synthesis.
100 200 300 400 500 600 7000
20
40
60
80
100
(a)
(b)
Temperature (°C)
W
eght
(%)
(c)
Figure 3.15. TG curves of the samples after the treatment of hydrogen peroxide: (a)
plain crosslinked chitosan (GA-CS), (b) A-cal, and (c) A-H2O2.
3.2.2.7 Synthesis of FAU nanocrystals
FAU nanocrystals were synthesized in GA-crosslinked CS hydrogels by simply varying
the compositions of the alkaline solution and heating times. Figure 3.16 displays the
XRD pattern, SEM image, and particle size distributions (from SEM and light
scattering) of the FAU nanocrystals. The Si/Al ratio of the synthesized crystals was
determined to be 2.07: 1.00 by the EDXS analysis, suggesting the nanocrystals are
Chapter 3
70
zeolite Y. The nanocrystals have a high crystallinity and a narrow size distribution.
Their average crystal size measured by SEM is 192 nm; the mean particle size from
light scattering shows at 171 nm. This confirms that the zeolite NaY produced in this
way can be well dispersed in deionized water.
Figure 3.16. (a) XRD pattern, (b) SEM image, and (c) particle size distribution by SEM
and (d) particle size distribution by light scattering of zeolite FAU (NaY) nanocrystals.
Figure 3.17 shows TG and nitrogen adsorption-desorption isotherm of the zeolite Y
nanocrystals. A mass loss of approximately 15% occurs, which is due mainly to loss of
structural water from the zeolite. This confirms that no polymer molecules remain in the
voids of zeolite, a conclusion that is supported by the nitrogen adsorption-desorption
isotherm in Figure 3.17b. The sample exhibits a much higher nitrogen adsorption
capacity than the NaA samples because NaY has larger, nitrogen accessible pores [240].
The BET specific surface area of zeolite Y nanocrystals is calculated to be 602.2 m2/g.
Chapter 3
71
Figure 3.17. (a) TG curve and (b) nitrogen adsorption-desorption isotherm of zeolite Y
samples
3.2.3 Summary
In section 3.2, the glutaraldehyde crosslinked chitosan (GA-CS) hydrogels with three-
dimensional network structures have been shown effective for controlling the growth of
zeolite NaA and NaY. The zeolite crystal sizes were significantly affected by the
formulation of silica-containing GA-CS hydrogels and alkaline solution, and by the
aging and heating conditions. The zeolite NaA nanocrystals with an average size of 148
nm and NaY with an average size of 192 nm were synthesized in this study. A novel
method of using hydrogen peroxide solution was developed to remove GA-CS
hydrogels after zeolite synthesis. TGA results confirmed that polymer hydrogels were
completely removed by this hydrogen peroxide treatment method. The NaA samples
obtained via this method exhibited much higher crystallinity than those obtained via
conventional calcination. This suggested that the hydrogen peroxide treatment method
be preferred for removal of GA-CS hydrogels. In addition, the zeolite NaA and NaY
nanocrystals produced here are readily dispersed in solvents such as deionized water,
and therefore they may be useful for applications such as in the fabrication of zeolite-
polymer composite membranes and hierarchical porous zeolitic structures.
Chapter 3
72
3.3 Formation of cubic zeolite A with an amorphous core in uncrosslinked chitosan hydrogels
3.3.1 Experimental
3.3.1.1 Synthesis of cubes of zeolite A with an amorphous core
Acetic acid (99%, Sigma–Aldrich; 7 g of 1M) was dissolved in deionized water (14 g)
in a polypropylene bottle. Chitosan (average molecular weight 120000 gmol-1, ca. 80%
deacetylated, Sigma–Aldrich; 1.2 g) was dissolved in the prepared acetic acid solution
under magnetic stirring for 1 h, followed by addition of the silica sol (HS-30 30 wt%,
Sigma–Aldrich; 3.38 g) to the chitosan/acetic acid solution. The alkaline solution was
prepared by mixing NaOH (99%, Merck; 5 g), and NaAlO2 (anhydrous, Sigma–Aldrich;
2.45 g) with deionized water (20 g). The solution was stirred for 0.5–1 h until it became
clear. The Na2O/Al2O3/H2O alkaline solution was added to the chitosan/acetic acid
solution without stirring, resulting in a sodium aluminosilicate gel entrapped inside the
chitosan hydrogel. The final molar composition of chitosan/SiO2 was 1.18:1. After
hydrothermal treatment at 90 °C for 3 h, the samples were washed with sufficient water
and dried at 80–100 °C overnight, followed by calcining the dried sample at 500 °C in
oxygen, or treating them with 10% hydrogen peroxide to remove chitosan. In addition,
samples were also synthesized at 90 °C with different hydrothermal reaction times (1, 2,
4, and 6 h).
3.3.1.2 Characterization
Scanning electron microscopy (SEM) images were taken with a JSM-6300 F
microscope (JEOL). Transmission electron microscopy (TEM) images and selected-area
electron diffraction (SAED) were taken with a JEOL JEM-2011 electron microscope
operated at 200 kV. X-ray diffraction (XRD) patterns were recorded on a Philips
PW1140/90 diffractometer with Cu Ka radiation at a scan rate of 2° min-1 and a step
size of 0.02°.
TEM observation of the cross sections of cubes of zeolite A with an amorphous core
Chapter 3
73
was carried by the following steps. The cross sections were prepared with a FEI xT
Nova Nanolab 200 DualBeam Focused Ion Beam (FIB) miller equipped with a LMIS of
Ga and a field emission electron gun. To avoid the aggregation of the powder, the
sample was dispersed in ethanol to form a suspension using an ultrasonic bath for 10
seconds. A drop of the solution was dropped onto a piece of single crystal Si and dried
in air, so a single layer of individual particles was settled on the surface of Si. The
sample was coated with 25 nm thick Au in EmiTech K550x sputter coater to ensure a
sound conductivity and to prevent the initial beam damage within the ion beam miller.
The specimens were produced using a FEI Nova Nanolab 200 Dualbeam FIB. A strip of
Pt was deposited onto the selected particles to fix them onto the Si matrix. A thin slice
about 100 nm thick was made with the FIB, and plucked with an ex-situ
micromanipulator, and then transferred onto a standard Cu grid with the carbon film.
The dark field micrographs were taken in JEOL 1400 TEM under 100kV.
3.3.2 Results and Discussion
Figure 3.18. (a) XRD pattern and (b) SEM image of the as-synthesized sample.
XRD pattern (Figure 3.18a) indicates the as-synthesized sample has the structure of
zeolite A. SEM image (Figure 3.18b) shows cube-like crystals with a particle size of
0.5–1.5 mm. This morphology with six {100} facets is typical for zeolite A, which has a
cubic structure with the unit cell parameter α = 2.461 nm, and space group Fm3c. The
Chapter 3
74
characteristic polyhedron normally indicates a single crystal property of zeolite A.
According to the classic crystal growth theory, crystals normally develop from nuclei
and the appearance of the facets is due to the differences in their growth rate [241-245].
Figure 3.19. TEM images of a typical zeolite A particle with a cube-like morphology
and the corresponding SAED patterns obtained from the entire particle. (a) Original
particle and (b) the same particle after beam irradiation for a few minutes.
TEM confirms the cube-like or rectangular morphology of the samples. Figure 3.19a
shows a TEM image of a typical rectangular particle of zeolite A with the
corresponding selected area electron diffraction (SAED) pattern, which is a standard
single crystal diffraction pattern viewed down the [100] zone axis. Many particles were
examined, and the single crystalline nature of zeolite A was observed in each case
without evidence of any polycrystallinity and twin defects. However, the image contrast
implies a core–shell structure, in which the core appears to be disordered. Under the
electron beam of the microscope, the disordered core of the zeolite A reduced in volume
and separated itself from the shell in a matter of few minutes. The shell remained intact,
which clearly appeared as a rectangle with a thickness of about 7nm (Figure 3.19b). The
SAED pattern from the particle in Figure 3.19b is almost identical to the pattern in
Figure 3.19a, which indicates that the shell structure was maintained after the irradiation
and the separation of the core. No other diffraction spots were observed, indicating that
the core is amorphous, rather than polycrystalline as in the case of zeolite analcime
[100]. As the material is very sensitive to the beam, HRTEM images of the crystalline
Chapter 3
75
shell were not acquired. The low-magnification TEM images and the SAED patterns
allow us to describe the zeolite A as a monocrystalline cube-like or rectangular box with
an amorphous core.
Figure 3.20. Dark field TEM images of the cross sections of cubes of zeolite A with an
amorphous core. These images indicate that the shell thickness varies in individual
crystals.
Figure 3.21. (a) SEM image and (b) XRD pattern of the sample treated with 0.35 M
acetic acid solution for 4 h. The insert in (a) is a TEM image of boxes.
Chapter 3
76
Figure 3.22. SEM images of sample prepared without addition of chitosan, before (a)
and after (b) treatment in 0.35 M acetic acid solution for 4 h.
The core-shell structure of the as-synthesized zeolite A was further supported by dark
field TEM images of the cross-sections of the cubes prepared by focused ion beam
milling (Figure 3.20) and by dissolution of the core component in an acidic aqueous
solution. During the latter process, 30 mL of 0.35 M acetic acid solution was added to 1
g of the zeolite A under stirring for 4 h. Most of the cubes lost their inner filling and
micrometer-sized hollow cube-like structures were produced (Figure 3.21a). According
to the XRD pattern these hollow structures were amorphous (Figure 3.21b). As the rate
of dissolution of the amorphous core is much faster than the crystalline shell, the cube-
like or rectangular outer shape was retained, although the crystallinity of the shell was
lost during the acidic treatment. For comparison, zeolite A crystals were prepared
without chitosan; however, the resulting particles were amorphous (Figure 3.22).
To investigate the crystallization of zeolite A during the hydrothermal reaction, the
products obtained during 1 to 6 h of reaction time were examined. The sample was
amorphous after 1 h, whereas those obtained after 2 and 3 h had zeolite A structure.
After 4 h and 6 h of hydrothermal reactions, a mixture of zeolite A and sodalite (Figure
3.23) was obtained. After the 0.35 M acetic acid treatment, all the crystalline samples
obtained from hydrothermal synthesis of 2–6 h had a rectangular or cube-like
morphology (Figure 3.24). These results indicate that extending hydrothermal reaction
time did not lead to the crystallization of the cores of zeolite A, but resulted in the
transformation of the crystal structure of the shell.
Chapter 3
77
Figure 3.23. XRD of the samples prepared with different hydrothermal times. (a) 1 h, (b)
2 h, (c) 3 h, (d) 4 h, and (e) 6 h (The peak labeled by asterisk is one of sodalite
characteristic peaks.).
Figure 3.24. SEM images of samples prepared with different hydrothermal times and
then treated with 0.35 M acetic acid solution for 4h. (a) 2 h, (b) 3 h, (c) 4 h and (d) 6 h.
Chapter 3
78
Scheme 3.2. Representation of the formation of cubes of zeolite A, with an amorphous
core (e). The rounded boxes in a) — d) are about 4 μm × 4 μm in dimension.
Scheme 3.2 shows the formation of cubes of zeolite A with an amorphous core. Initially,
the silica sol is dispersed in an acidified chitosan solution (Scheme 3.2a). After addition
of the alkaline solution, the chitosan molecules are deprotonated, resulting in a hydrogel
with microsized three-dimensional pores (Scheme 3.2b). The sodium aluminosilicate
gel is produced inside the chitosan hydrogel network by the reaction between silica and
the alkaline solution. During the hydrothermal treatment, zeolite nucleation takes place
mainly on the surface of the aluminosilicate aggregates (Scheme 3.2c). Similar to the
case of zeolite analcime [100], some crystalline islands might initially form on the
surface of the aluminosilicate. These islands then join together, leading to a
monocrystalline cube-like shape by self-alignment of their crystallographic orientations
(Scheme 3.2d). Thus, it is an interesting observation that a very thin-walled crystalline
cube-like or rectangular morphology can be developed on the surface of an amorphous
cluster without any specific relationship to the crystal growth rate of the crystal planes
from a nucleus in the amorphous center (Scheme 3.2d). The driving force behind the
formation of such polyhedral shells is the one that minimizes the surface energy.
Chapter 3
79
It has also been observed that organic polymers have significant effects on zeolite
crystallization [11, 12, 14, 99]. The addition of water-soluble polymers in the zeolite gel
could dramatically shorten the prenucleation and nucleation periods and thus accelerate
the crystal growth [14]. Mathematical modeling [235, 236] and experimental results
[237] indicate that zeolite nucleation takes place at the interface between the solution
and the gel by adsorption and rearrangement of the soluble precursors. In the synthesis
described herein, the uncrosslinked chitosan hydrogel networks are highly swollen by
the solution, and the interfaces between the chitosan polymer networks and zeolite
aluminosilicate gel can serve as ideal nucleation sites. Such unique interfaces facilitate
zeolite crystallization from the surface of the aluminosilicate gel aggregates. On the
other hand, the chitosan hydrogel may also play a role in confining the aluminosilicate
aggregates and thus controlling the sizes of the zeolite A cubes. During the
hydrothermal treatment, the crystalline shell limits the diffusion of the solution and thus
the crystallization of the cores is not able to proceed. It is worth mentioning that fully
crystallized zeolite A cubes were obtained when the gel was aged overnight at room
temperature before the hydrothermal treatment. The aging process most likely makes
the system more uniform, in which the chitosan-facilitated zeolite nucleation becomes
kinetically less pronounced. In addition, the desired network structure of chitosan
hydrogels is essential for the formation of cubes of zeolite A with an amorphous core.
Owing to the presence of crosslinked chitosan hydrogels, the small hydrogel pores
greatly confined zeolite growth leading to zeolite nanocrystals.
3.3.3 Summary
In summary, the cubes of zeolite A consisting of a thin crystalline shell and
anamorphous core have produced by using uncrosslinked chitosan hydrogels. Results
showed that the formation of cube-like or rectangular core-shell structures caused by
particle aggregation and surface-to-core crystallization induced by chitosan networks.
Chapter 3
80
3.4 Comparisons of zeolite formation mechanisms Chitosan was used as polymer in both Section 3.2 and 3.3, but resulting in the formation
of zeolite with different structures.
In Section 3.2, as shown in Scheme 3.3a, chitosan polymer was disscolved in acetic acid
solution and then crosslinked with glutaraldehyde (GA), forming a three-dimentional
network for the zeolite growth. After that, alkaline solution was introduced into GA-CS
hydrogels, which then was heated for zeolite crystallization inside the polymer pores.
Here, the GA-CS polymer plays a similar role as confined-space matrix, similar to
carbon black, starch, etc, in the conventional confined-space synthesis of zeolite. The
zeolite nanocrystals were formed and confined by the pore spaces of GA-CS. As
discussed in literature review, a confined-space matrix requires two basic requirements,
inertness and stability in hydrothermal synthesis conditions; and with pores of narrow
size distribution. Some previous studies have found that chitosan is not stable, possibly
swelling or degradable in the aqueous condition with heat or high pH, which normally
occurs in the zeolite hydrothermal synthesis [112, 113]. Moreover, Micro-pore sizes of
GA-CS may expand after exposure to water. Therefore, there large zeolite particles
were formed in some of my experimental results. However, by controlling the synthetic
conditions, zeolite nanoparticles can still be produced. Heated hydrogen peroxide can be
used to remove the polymer gel without using high-temperature combustion. The
products formed by using this method have high re-dispersibility in water, which can be
applied for the applications, such as the fabrication of zeolite-polymer mixed matrix
membranes.
In Section 3.3, CS polymer was similaly dissolved in an acetic solution as in Section 3.2,
with silica dispersing in. Upon raising the pH, it has been proven that amino groups in
CS are increasingly deportonated and become available for hydrogen bonding as shown
in Scheme 3.3b [246]. At high pH, the CS molecules in solution develop enough
hydrogen bonds to establish a gel netwok. As the pH is raised further, deportonation
continues and the molecules form additional miniature crystalline domains. This effect
results in an increase in gel stiffness and can be associated with minor gel contraction.
After the chitosan is deprotonated and uncharged in a high-pH alkaline solution, there
may be a hydrogel with micro-sized three dimnational pores formed, giving 3-D
Chapter 3
81
networks for the sodium aluminosilicate gel in the reaction between silica and the
alkaline solution. Chen et al. developed a mechanism demonstrating complex
geometrical structure built via the route of “Nanocrystallites” − “oriented Aggregation”
− “surface Recrystallization” − “Single crystals”, which is designated as the “NARS”
route [100]. In the chitosan gel, the zeolite nucleation may also take place mainly on the
surface of the aluminosilicate aggregates. Then the crystalline islands joined together
tightly, leading to a thin-walled crystalline cube-like or rectangular morphology by self-
alignment of their crystallographic orientations. Comparing with the formation of
zeolite nanoparticles in the Section 3.2, the formation of such polyhedral shells may be
driven by minimizing the surface energy. In other words, in a cluster of nanoparticles,
fewer and larger crystals with smaller surface-to-volume ratios may form rather than the
smaller particles, in order to reduce the energy of the entire system.
N
HH
O
nO
n
NH3
O
H
OO
H
(a)
n
NH3
O
H
O
NH2
n
+ n
(b)
Scheme 3.3. (a) Crosslinking reaction between glutaraldehyde (GA) and chitosan
molecules; (b) pH dependent protonation/deprotonation of the chitosan molecule [246].
Chapter 3
82
3.5 Conclusions In Chapter 3, chitosan hydrogels were found to have a significant effect on the zeolite
crystallization and growth. Crystallization of zeolite NaA and NaY in glutaraldehyde-
crosslinked chitosan (GA-CS) hydrogels was studied in the section 3.2. The zeolite
crystals were produced by penetration of Na2O-Al2O3-H2O alkaline solution into GA-
CS hydrogels filled with colloidal silica, followed by hydrothermal treatment and
removal of GA-CS hydrogels. The effects of the synthesis parameters —including the
amounts of silica, chitosan, and glutaraldehyde, and the aging and heating times — on
the size, size distribution and crystallinity of the particles were systematically
investigated. A hydrogen peroxide treatment method was shown to be an effective way
for removing GA-CS hydrogels, thereby avoiding the conventional calcination step. X-
ray diffraction (XRD), light scattering, scanning electron microscopy (SEM),
thermogravimetric analysis (TG), and N2 sorption were used to characterize the zeolite
samples. This work showed that GA-CS is a promising space-confinement medium for
the synthesis of zeolite nanocrystals with tunable crystal sizes and excellent
dispersibility. The zeolite NaA and NaY nanocrystals produced here are readily
dispersed in solvents such as deionized water, and therefore they may be useful for
applications such as in the fabrication of zeolite-polymer composite membranes and
hierarchical porous zeolitic structures. In section 3.3, the results showed cubes of zeolite
A consisting of a thin crystalline shell and an amorphous core can be grown within
uncrosslinked chitosan hydrogels. It is indicative that the formation of cube-like or
rectangular core-shell structures involves particle aggregation and surface-to-core
crystallization induced by chitosan networks. This work may provide a new model
system for studying complex zeolite nucleation and growth mechanisms.
As mentioned, Chapter 3 provides a novel way to produce NaA and NaY nanocrystals
with good redispersibility and high crystallinity, which are highly useful in the practical
applications, such as the fabrication of MMMs. There have been plenty of current works
on the use of LTA-type and FAU-type of zeolite nanocrystals in the zeolite-polymer
composite membranes for gas separation. In my research, the focus is on another type of
zeolite, with smaller pore sizes and channels, which is sodalite (SOD). Based on the
above Literature Review, pure inorganic sodalite membranes have been fabricated and
observed with an excellent hydrogen selectivity and permeability. However, no study
Chapter 3
83
has been reported on the MMMs with the incorporation of sodalite nanoparticles in the
literature. In order to improve the gas selectivity or permeability or both of them, one of
the suggestive ways is to organic functionalize zeolite nanoparticles, which is presented
in the following chapter, Chapter 4.
Chapter 4
84
Chapter 4 Organic-functionalized Sodalite Nanocrystals
4.1 Overview
Chapter 4 presents the synthesis of organic-functionalized sodalite nanocrystals and
their characterization. Hydroxy-sodalite nanocrystals with organic functional groups
(i.e.,=Si–(CH3)(CH2)3NH2, denoted Sod-N, or ≡Si-CH3, denoted Soc-C) were
synthesized by the direct transformation of organic-functionalized silicalite nanocrystals.
In the transformation process, silicalite with organic functional groups became
amorphous first and then re-crystallized, yielding a sodalite structure. The chemical
structure of organic-functionalized sodalite nanocrystals was confirmed by 29Si MAS
NMR spectroscopy. Gas sorption results showed that the sodalite nanocrystals
contained uniform pore channels that were accessible to hydrogen, but inaccessible to
nitrogen, as expected. The dispersion of Sod-N and Sod-C in organic solvents was
favored by the presence of organic functional groups.
4.2 Experimental
4.2.1 Synthesis of organic-functionalized silicalite nanocrystals
Clear synthesis solutions were prepared by dropwise addition of 20 g of 1 M
tetrapropylammonium hydroxide (TPAOH, Sigma-Aldrich) solution into the mixture of
17.8 g of tetraethyl orthosilicate (TEOS, 99%, Sigma-Aldrich) and 1.8 g of 3-
aminopropyl(diethoxy) methylsilane (ADMS, 97%, Sigma-Aldrich) or 1.3 g of
methyltrimethoxysilane (MTMS, 98%, Sigma-Aldrich) under vigorous stirring,
followed by continued stirring at room temperature for 3 h. The molar composition of
final solution was 1 TPAOH: 4.32 SiO2: 0.48 ADMS (or MTMS): 44 H2O.
Crystallization was carried out at 80 °C for 12-15 days. The milky silicalite suspensions
obtained were dried at 90-100°C leading to solid silicalites (denoted Sil-N and Sil-C for
silicalites prepared with ADMS and MTMS, respectively). To observe their
morphologies by scanning electron microscopy, the samples were prepared by repeated
Chapter 4
85
cycles of washing with deionized water and centrifuging, followed by drying at 90-100
°C overnight.
4.2.2 Synthesis of organic-functionalized sodalite nanocrystals
An alkaline solution with a molar composition of 6.07 Na2O:1 Al2O3:66 H2O was
prepared by mixing 20 g of sodium hydroxide (99%, Merck), 9.2 g of sodium
aluminates (anhydrous, Sigma-Aldrich), and 60 g of deionized water at room
temperature for 1-2 h. 1 g of the dried silicalite sample (i.e. Sil-N and Sil-C) was added
to 11 g of the alkaline solution during 2-3 min of stirring, and then aged at room
temperature for 4 h without further stirring. The transformation was carried out at 80 °C
for 0-4 h. The samples obtained were cooled to room temperature and collected by
repeated cycles of washing with deionized water and centrifuging, followed by drying at
90-100 °C overnight. The samples were denoted Sod-N and Sod-C, respectively, when
Sil-N and Sil-C were used as silica source, respectively. For comparison, hydroxy-
sodalite nanocrystals (denoted Sod) were also prepared from silicalite nanocrystals
according by applying similar method without adding organic silane.
4.2.3 Characterization
Scanning electron microscopy (SEM) images were taken with a JSM-6300F microscope
(JEOL). The particle size distributions for Sil-N, Sil-C, Sod-N and Sod-C were
determined by manual measurement of 300 nanocrystals each in SEM images with a
Photoshop software. X-ray diffraction (XRD) patterns were measured on a Philips
PW1140/90 diffractometer with Cu Kα radiation (25 mA and 40 kV) at a scan rate of 1
°/min with a step size of 0.02°. Fourier transform infrared spectra (FT-IR) were
recorded for the samples embedded in KBr pellets with a GX Spectrometer (Perkin
Elmer). Thermogravimetric analysis (TGA, Perkin Elmer, Pyris 1 analyzer) was
performed in air at a heating rate of 5 °C/min to 600 °C. 29Si solid-state nuclear
magnetic resonance (NMR) was conducted on a Bruker DSX300 spectrometer
(Germany) under conditions of cross polarization (CP) and magic angle spinning
(MAS). 29Si solid-state MAS NMR spectra were collected at room temperature with a
frequency of 59.6 MHz, a recycling delay of 600 s, a radiation frequency intensity of
Chapter 4
86
62.5 kHz, and a reference sample of Q8M8(sf[(CH3)3SiO]8Si8O12]). Nitrogen and
hydrogen adsorption-desorption experiments were performed at 77 K with a
Micrometritics ASAP 2020MC analyzer and a Micrometritics ASAP 2010MC analyzer,
respectively. The samples were degassed at 473 K before analysis. The surface areas
were determined by the Brunauer-Emmett-Teller (BET) method. Suspended particle
size distributions were quantified by light scattering with a Malvern Mastersizer 2000
analyzer. Different solvents ― deionized water, isopropanol (97%, Sigma-Aldrich),
dichloromethane (DCM, Sigma-Aldrich) and dimethylformamide (DMF, Sigma-
Aldrich) ― were used for sample dispersion. Approximately 10 mL of suspensions
were prepared in 30 mL vials by dispersing 20 mg of sample under ultrasonication, and
kept still for 30 min before the photos of samples were taken by a digital camera.
Approximately 12-15 mL of suspension was prepared by dispersing 50 mg of sample
into 50 mL of solvent under ultrasonication before injection into the Mastersizer for size
distribution analysis.
4.3 Results and Discussion
4.3.1 Transformation of silicalite
The XRD patterns (Figure 4.1) show the transformation of organic-functionalized
silicalites (Sil-N and Sil-C) under hydrothermal treatment at 80 °C. The organic-
functionalized silicalites (Sil-N and Sil-C) became amorphous after 1 h hydrothermal
treatment. However, in the previous study [84], plain silicalite (without organic groups)
was largely transformed into zeolite A after only 1 h hydrothermal treatment. This is
because the presence of =Si−(CH3)(CH2)3NH2 and ≡Si−CH3 in silicalite structures (Sil-
N and Sil-C, respectively) does not favor aluminosilicate structure rearrangement during
the incorporation of Al and Na. After 2 h treatment, both samples were a mixture of
zeolite A and sodalite. The pure organic-functionalized sodalite, Sod-N, was obtained
after 3 h. However, the transformation of Sil-C into Sod-C took a longer time (4 h) to
complete.
Chapter 4
87
10 20 30 40 50 60
3h
2h
1h
0h
In
tens
ity (a
.u.)
2θ (degrees)10 20 30 40 50 60
3h2h1h
0h
4h
Inte
nsity
(a.u
.)
2θ (degrees)
(a) (b)
Figure 4.1. XRD patterns of samples prepared with dried organic-functionalized
silicalites by hydrothermal treatment at 80 °C for different times. (a)Sil-N to Sod-N, (b)
Sil-C to Sod-C.
3500 3000 2500 2000 1500 1000 500
550
461
661714
990
10901220
Tran
smitt
ance
%
Wavelength (cm-1 )
Sod-N
Sil-N
428
3500 3000 2500 2000 1500 1000 500
550
428
461
661
714990
10901220
Sil-C
Wavelength (cm-1 )
Tran
smitt
ance
%
Sod-C
(a) (b)
Figure 4.2. FT-IR spectra of samples (a) Sil-N and Sod-N, obtained after 3 h
hydrothermal reaction, (b) Sil-C and Sod-C, obtained after 4 h hydrothermal reaction.
To investigate the transformation of silicalites to sodalites, Sil-N, Sil-C and Sod-N,
Sod-C samples were characterized by FT-IR spectroscopy (Figure 4.2a and b). The
characteristic bands of the silicalite (Sil-N and Sil-C) Si−O−Si framework are the
Chapter 4
88
double-ring vibration at approximately 550 cm-1 and the stretching vibrations at 1090
and 1220 cm-1. The characteristic adsorption band for the single four-membered ring of
the sodalite unit occurs at 428 cm-1. The adsorption between 714 cm-1 and 661 cm-1 is
due to the symmetric stretch of T-O-T (T=Si, Al), the band at 990 cm-1 is assigned to its
asymmetric stretch and the bands at 461 and 428 cm-1 arise from the bending vibration
of O-T-O [84, 247, 248].
Figure 4.3. SEM images (a, b, d, c and e) and particle size distributions (c, f) of organic-
functionalized silicalites and organic-functionalized sodalites. SEM images: (a) dried
silicalite Sil-N, (b) sodalite Sod-N obtained after 3 h hydrothermal reaction, (d) dried
silicalite Sil-C, and (e) sodalite Sod-C obtained after 4 h hydrothermal reaction. Particle
size distributions: (c) Sil-N and Sod-N obtained after 3 h hydrothermal reaction, and (f)
Sil-C and Sod-C obtained after 4 h hydrothermal reaction.
Figure 4.3 shows the SEM images and particle size distributions of organic-
functionalized silicalite nanocrystals (Sil-N and Sil-C) and organic-functionalized
sodalite nanocrystals (Sod-N and Sod-C). All samples exhibit similar morphologies. Sil-
N exhibits smaller particle sizes as compared with Sil-C, though the synthesis
Chapter 4
89
conditions were identical. This may be explained by the presence of the –NH2 groups
accelerating nucleation in the silicalite synthesis solution, leading to smaller particles on
average [249]. This is also consistent with the XRD results above showing that the
transformation of Sil-N into Sod-N took a shorter time. The particle sizes of the
organic-functionalized sodalite nanocrystals are larger than those of their precursor
silicalite nanocrystals. This is related to the recrystallization in the transformation as
indicated by XRD. The mean particle sizes are 95 nm, 105 nm, 105 nm and 140 nm for
Sil-N, Sod-N, Sil-C and Sod-C, respectively (Figure 4.3c and f).
4.3.2 Evidence of organic functionalization of sodalite
100 50 0 -50 -100 -150 -200 -250
Si-C (*)
Sod-NSod-C
ppm
Si (4Al)
(a)
SiO
OAl
OSi
O AlO
SiO
Si CH3
SiHO O
AlO
Si
(CH2)3NH2
SiO
OAl
OSi
O AlO
SiO
Si CH3
SiHO O
AlO
Si
Sod-N Sod-C
(b)
Figure 4.4. (a) 29Si-NMR of organic-functionalized sodalite nanocrystals and (b) the
bonding scheme for organic-functionalized sodalite nanocrystals.
Chapter 4
90
To prove that the organic functional groups have been attached onto the sodalite
nanoparticles, Sod-N and Sod-C samples were characterized by solid-state NMR
spectroscopy. The 29Si MAS NMR spectra shown in Figure 4.4a display a strong
resonance peak at around -85 ppm, which arises from Si (4Al) in Sod-N and Sod-C [250,
251]. The NMR spectra also exhibit a resonance peak at around -55 ppm, which is
ascribed to Si-C bonds [251]. The results confirm the existence of organic functional
groups in the Sod-N and Sod-C, and thus the organic groups have been attached with
the sodalites. The integrated area of the functionalized silicon peak represents 7.4 mole
% and 7.2 mole % of the total silicon in Sod-N and Sod-C respectively. The amounts of
organic functional groups attached with Sod-N and Sod-C are less than those added in
silicalite synthesis solutions (10 mole % was added for both Sil-N and Sil-C), but this is
reasonable given that a proportion of the hydrolyzed ADMS and MTMS would have
remained in the synthetic solutions. The Si-C bonds labeled with asterisk in organic-
functionalized sodalites are illustrated in Figure 4.4b.
0 100 200 300 400 500 600
85
90
95
100
c
b
Temperature (oC)
Mas
s (%
)
a
Figure 4.5. TGA curves of organic-functionalized sodalite nanocrystals and hydroxyl-
sodalite nanocrystals. (a) Sod-N, (b) Sod-C and (c) Sod.
The organic functionalization of the sodalite nanocrystals receives further support from
the TGA results, which are shown in Figure 4.5. The mass loss of the pure sodalite was
about 11 wt% owing to the loss of the structural water (Figure 4.5c) [84]. The mass
losses for Sod-N and Sod-C were 13.6 wt% and 11.5 wt% respectively. As compared
Chapter 4
91
with the pure sodalite nanocrystals, the additional mass loss of 2.6 wt% for Sod-N and
of 0.5 wt% for Sod-C was due to decomposition of organic functional groups (i.e.,
−(CH3)(CH2)3NH2 or −CH3) at high temperatures [22]. These figures are quite
consistent with the expected mass losses of 3.18 wt% for Sod-N and 0.65 wt% for Sod-
C that can be calculated from the proportion of Si–C bonds measured by 29Si MAS
NMR.
4.3.3 Gas adsorption and pore structures
To further compare the organic-functionalized sodalite nanocrystals (Sod-N and Sod-C)
and plain hydroxyl-sodalite nanocrystals (Sod), nitrogen and hydrogen adsorption-
desorption analyses were conducted. The isotherms of Sod-N, Sod-C and Sod are
shown in Figure 4.6. The amounts of nitrogen adsorbed in all three samples are very
low at low relative pressures, and substantially increase at high relative pressures (e.g.
P/Po>0.8). This is because well-grown sodalite pores are inaccessible to nitrogen (N2
kinetic diameter 3.6 Å is larger than sodalite pore size 2.8 Å), and the main nitrogen
adsorption arises from the external surfaces of nanocrystals. The BET surface areas are
calculated to be 22.8, 19.6 and 19.1m2/g for Sod, Sod-N, and Sod-C, respectively,
which is consistent with the particle size distributions observed by SEM. By contrast,
all samples exhibit much higher H2 adsorption at low relative pressures as compared
with N2 adsorption (Figure 4.6a, b), implying that the sodalite channels in these three
samples are readily accessed by H2 molecules. Furthermore, the organic-functionalized
sodalites (Sod-N and Sod-C) possess slightly lower H2 adsorption than pure sodalite
(Sod). At P/Po ≈1, the volume of hydrogen absorbed is around 33.0 cm3/g for Sod, 26.5
cm3/g for Sod-N, and 28.0 cm3/g Sod-C. Therefore, the organic groups do not
substantially change the hydrogen adsorption of the sodalite nanocrystals. Clearly, this
finding is essential if the functionalized nanoparticles are to be used successfully in H2
separation membranes.
Chapter 4
92
0.0 0.2 0.4 0.6 0.8 1.00
10
20
30
40
50
60
Vo
lum
e Ad
sorb
ed (c
m3 g-1
)
P/Po
Sod-N Sod-C Sod
0.0 0.2 0.4 0.6 0.8 1.0
0
5
10
15
20
25
30
35
Vol
umed
ads
orbe
d (c
m3 g-1
)
P/Po
Sod-N Sod-C Sod
(a) (b)
Figure 4.6. (a) Nitrogen and (b) hydrogen adsorption-desorption isothermals of plain
sodalites (Sod) and organic-functionalized sodalites (Sod-N and Sod-C).
4.3.4 Surface modification: dispersion in solvents
To study the effect of organic functionalization on the dispersibility of sodalite
nanocrystals, a series of solvents of different polarities was selected: deionized water,
isopropanol, dichlormethane (DCM), and dimethlformamide (DMF). The solvent
polarity of this series, in descending order, is water (100) > DMF (42.88) > isopropanol
(36.72) > DCM (23.04) [252]. The particle size distributions of Sod-N, Sod-C, and Sod
shown in Figure 4.7 are used as an indicator of their relative dispersibility. When
deionized water is used as a dispersion medium, both Sod-N and Sod have a similar
particle size distribution and their mean particle sizes are approximately 160 nm, which
is slightly greater than that observed by SEM due to the surface solvation effect (e.g.,
surface ionization and adsorption) [253, 254]; In contrast, Sod-C exhibits a wider
particle size distribution and its mean particle size is approximately 270 nm (Figure
4.7a). The different dispersibility between Sod-N/Sod and Sod-C arises from their
different surface energy components: Sod-N with –(CH3)(CH2)3NH2 groups and Sod
with –OH groups have similar hydrogen- bonding forces, whereas Sod-C with –CH3
groups is more hydrophobic. Sod-N, Sod-C, and Sod show similar dispersibility in
DMF (Figure 4.7b) because DMF combines a high polarity and high hydrogen-bonding
force with hydrophobic groups. In isopropanol, both Sod-N and Sod-C exhibit slightly
better dispersion than Sod (Figure 4.7c). Sod-N and Sod-C exhibit similar degrees of
Chapter 4
93
dispersion in DCM, but the Sod nanocrystals severely aggregate, leading to a mean
particle size of 880 nm (Figure 4.7d). These are because isopropanol and DCM, with
relatively low polarity and poor hydrogen-bonding force, preferentially interact with
organic-functionalized surfaces [255]. These results clearly show that the surface
properties of sodalite nanocrystals can be tailored by organic functionalization, which is
essential for preparing zeolite-polymer nanocomposites [256, 257].
100 1000
0
5
10
15
20
25
Num
ber (
%)
Particle size (nm)
Sod-N Sod-C Sod
100 1000
0
5
10
15
20
25
30
Num
ber (
%)
Particle size (nm)
Sod-N Sod-C Sod
(a) (b)
100 1000
0
5
10
15
20
25
Num
ber (
%)
Particle size (nm)
Sod-N Sod-C Sod
100 1000
0
5
10
15
20
25
30
Num
ber (
%)
Particle size (nm)
Sod-N Sod-C Sod
(c) (d)
Figure 4.7. Particle size distributions of organic-functionalized sodalite nanocrystals and
plain sodalite nanocrystals in different solvents: (a) deionized water, (b)
dimethylformamide (DMF), (c) isopropanol and (d) dichloromethane (DCM).
4.4 Conclusion
Organic functional groups have been successfully attached onto sodalite nanocrystals
through the direct transformation of organic-functionalized silicalite nanocrystals. The
organic-functionalized sodalite nanocrystals showed high crystallinity and well-grown
Chapter 4
94
pore structures based on XRD and nitrogen sorption measurements. The micropores of
the organic-functionalized sodalite nanocrystals were highly accessible to hydrogen
molecules, though there was a slight reduction of hydrogen adsorption compared with
sodalite nanocrystals without organic groups. Sodalite nanocrystals with –
(CH3)(CH2)3NH2 moieties showed good dispersibility in all four solvents (i.e., water,
isopropanol, dichloromethane, and dimethylformamide) tested whereas sodalite
nanocrystals with –CH3 groups were dispersible in isopropanol, dichloromethane and
dimethylformamide, but were agglomerated in water. Without organic functionalization,
sodalite nanocrystals showed very poor dispersibility in dichloromethane. Therefore, it
is expected that the organic-functionalized sodalite nanocrystals synthesized in this
work will be highly suited for fabricating sodalite-polymer nanocomposite membranes
and other zeolite nanostructures.
Chapter 5
95
Chapter 5 Preparation & Characterization of Mixed Matrix Membranes
5.1 Overview
Chapter 5 presents the preparation of mix matrix membranes (MMMs) fabricated from
organic-functionalized sodalite nanocrystals (Sod-N) dispersed in BTDA-MDA
polyimide matrices and their characterization for structure and gas-separation
performance. No voids are found upon investigation of the interfacial contact between
the inorganic and organic phases, even at a Sod-N loading of up to 35 wt%. This is due
to the functionalization of the zeolite nanocrystals with amino groups (=Si–
(CH3)(CH2)3NH2), which covalently link the particles to the polyimide chains in the
matrices. The addition of Sod-N increases the hydrogen-gas permeability of the
membranes, while nitrogen permeability decreases. Overall, these nanocomposite
membranes display substantial selectivity improvements. The sodalite-polyimide
membrane containing 35 wt% Sod-N has a hydrogen permeability of 8.04 Barrers and a
H2/N2 selectivity of 277 at 25 °C whereas the plain polyimide membrane exhibits a
hydrogen permeability of 6.94 Barrers and a H2/N2 selectivity of 193 at the same testing
temperature.
5.2 Experimental
5.2.1 Membrane fabrication
The amine-functionalized sodalite nanocrystals (denoted Sod-N) were synthesized by
transforming silicalite nanocrystals according to the method mentioned in Section 4.2.
Briefly, a clear synthesis solution was prepared by dropwise addition of 20 g of 1 M
tetrapropylammonium hydroxide (TPAOH, Sigma-Aldrich) solution into a mixture of
17.8 g of tetraethyl orthosilicate (TEOS, 99%, Sigma-Aldrich) and 1.8 g of 3-
aminopropyl(diethoxy) methylsilane (ADMS, 97%, Sigma-Aldrich) with vigorous
stirring, followed by continued stirring at room temperature for 3 h and then
Chapter 5
96
crystallization at 80 °C for 12-15 days. The milky silicalite suspensions so obtained
were dried at 90-100°C to obtain solid silicalites. An alkaline solution was prepared by
mixing 20 g of sodium hydroxide (99%, Merck), 9.2 g of sodium aluminates (anhydrous,
Sigma-Aldrich), and 60 g of deionized water at room temperature for 1-2 h. Around 1 g
of the dried silicalite sample (denoted Sil-N) was added to 10 g of the alkaline solution
during 2-3 min of stirring, and then allowed it to age at room temperature for 4 h
without further stirring. The transformation was carried out at 80 °C for 4 h. The
resulting amine-functionalized sodalite nanocrystals were cooled to room temperature
and collected by repeated cycles of washing with deionized water and centrifuging,
followed by drying overnight at 90-100 °C.
Monomers benzophenone-3,3’,4,4’-tetracarboxylic dianhydride (BTDA; Sigma-Aldrich)
and 4,4’-diaminodiphenylmethane (MDA; Sigma-Aldrich) were dried at ~150 °C and
~50 °C for at least 12 h under vacuum. Dimethylformamide (DMF) (GR, Merck) was
dried and stored with 4-Ǻ molecular sieves prior to use. To fabricate each composite
membrane, a given quantity of Sod-N nanocrystals was dispersed in 10 g of DMF under
ultrasonication at room temperature for 30 min. Then 1.5 g of BTDA and 0.92 g of
MDA were dissolved in the Sod-N suspension. The resulting mixture was stirred for 5 h
in an ice-water bath at approximately 0 °C under N2 gas to obtain a Sod-N/PAA
(poly(amic acid)) precursor, which was a cloudy yellow, viscous solution. The Sod-
N/PAA solution was cast directly onto a glass plate and placed into a vacuum oven and
heat treated for 2 h each at 50 °C, at 100 °C and at 150 °C, before it was held at 200 °C
overnight. The resulting sodalite-polyimide nanocomposite membrane (denoted Sod-
N/PI) was slowly cooling to room temperature. All of the yellow Sod-N/PI films were
immersed in hot water at 90 °C for 1 h to allow removal from the glass plates, after
which they were dried under vacuum at ~150 °C overnight before analysis. In this paper,
the sodalite-polyimide nanocomposite membranes were made with sodalite loadings of
15, 25 and 35 wt% (based on the weight of polyimide) and these are denoted PI-15, PI-
25, and PI-35, respectively. For comparison purposes, pure polyimide membranes were
prepared by applying the above procedures without any Sod-N additions and these are
referred to as PI-0.
Chapter 5
97
5.2.2 Characterization
Scanning electron microscopy (SEM) images were taken with a JSM-6300F microscope
(JEOL). X-ray diffraction (XRD) patterns were measured on a Philips PW1140/90
diffractometer with Cu Kα radiation (25 mA and 40 kV) at a scan rate of 1 °/min with a
step size of 0.01°. Fourier-transform infrared spectra (FT-IR) were recorded for the
samples embedded in KBr pellets with a GX Spectrometer (Perkin Elmer).
Thermogravimetric analysis (TGA, Perkin Elmer, Pyris 1 analyzer) was performed at a
heating rate of 5 °C/min to 700 °C in oxygen with a flow rate of 15 cm3⋅min-1.
Hydrogen adsorption–desorption experiments were performed at 77 K and room
temperature, and a pressure of up to 900 mmHg with a Micrometritics ASAP 2010MC
analyzer. The samples were degassed at 473 K before analysis.
Scheme 5.1. Apparatus for measuring gas permeance through the membrane.
To test gas separation properties, the pressure rise method was applied here as shown in
Scheme 5.1 [258]. The composite membrane or pure polyimide membrane samples
were firstly attached to a porous stainless-steel stand (pore size ~ 200 nm), which was
Chapter 5
98
then fixed in a stainless-steel holder (6) by using Torr Seal epoxy resin (Varian). Before
measurements, the samples were evacuated and dried in a vacuum oven at 200 °C
overnight to remove any residual solvent and adsorbed water. The gas permeation tests
were performed at 25 °C, 60 °C and 100 °C on pure H2 and pure N2. A pure gas flows
through the glass cell at atmospheric pressure by adjusting the valve (13). The
downstream is pumped by the vacuum pump (7), and its pressure rise of the permeate
stream was measured with a Series 901 Transducer (MKS) is monitored by the pressure
transducer (8). After equilibrium is reached, the value (11) is turned off; the
downstream pressure rises linearly with time, and is recorded by the computer (9).
Membrane permeability, Pi, was defined as [15, 259],
APNd
Pi
ii Δ=
……………………………..Equation 5-1
where d is the membrane thickness (cm), Ni the permeation rate of component i (cm3⋅s-1),
∆Pi the transmembrane pressure difference of i (cmHg), and A the membrane area (cm2).
1 Barrer = 10-10 cm3(STP)⋅cm⋅cm-2⋅s-1⋅cmHg-1. The selectivity, αij, between two gases, i
and j , was defined as [209, 260],
j
iij P
P=α
………………………………….Equation 5-2
The apparent activation energy Ep was analyzed according to the Arrhenius equation
[209-212],
⎟⎟⎠
⎞⎜⎜⎝
⎛ −=
RTE
PP pexp0
………………………….Equation 5-3
where P is the permeability, P0 the pre-exponential factor, R the ideal gas constant
(8.3143 J mol-1 K-1) and T is the temperature in Kelvin (K).
Chapter 5
99
5.3 Results and Discussion
5.3.1 Membrane characterization
Scheme 5.2. Preparation of sodalite-N/PI nanocomposite membranes.
A schematic diagram for the fabrication of Sod-N/PI nanocomposite membranes is
shown in Scheme 5.2. The sodalite nanocrystals with amino reactive functional groups
(=Si–(CH3)(CH2)3NH2) (Sod-N) is firstly well dispersed in DMF organic solvent,
following with the addition of MDA and BTDA monomers at 0 °C. In this paper, the
most widely applied method in polyimide synthesis is used, called two-step imidization
reaction [261]. The dianhydride (BTDA) and diamine (MDA) are mixed and stirred in a
Chapter 5
100
dipolar aprotic solvent (DMF), resulting in polyamic acid (PAA) precursor solution,
which is then cyclized and imidized into the final polyimide (Scheme 5.3).
Scheme 5.3. Fabrication of BTDA-MDA polyimide by two-step method.
Chapter 5
101
Figure 5.1. Digital photos of PI-0, PI-15, PI-25 and PI-35 showing the change in
transparency with increasing Sod-N content.
10 20 30 40 50 60
* *
*
*
* *
**
(330
)
(310
)(2
22)
(211
)
Sod-N
PI-35
PI-25
PI-15
Inte
nsity
(a.u
.)
2θ
PI-0
(110
)
Figure 5.2. XRD patterns for samples Sod-N, PI-0, PI-15, PI-25, and PI-35. The peaks
labeled with asterisks arise from Sod-N.
Chapter 5
102
Figure 5.1 shows photographs of the series of polyimide composite membranes with a
thickness of 50 μm, which were all intact and homogeneous, laid over the word
“Monash”. Pure polyimides are clear, flexible and have good tear strength. All of the
composite membranes have a yellow appearance, but their transparency decreases with
increasing content of Sod-N nanocrystals (Figure 5.1), as is evident from the gradual
obscuration of the word from PI-0 to PI-35. Figure 5.2 shows the XRD patterns for pure
Sod-N and for PI-0, PI-15, PI-25 and PI-35. The Sod-N nanocrystals exhibit good
crystallinity, giving sharp peaks in XRD pattern, which have been indexed in Figure
5.2. In contrast, the pure polyimide membrane (PI-0) appears to be amorphous, as
expected. With increasing contents of Sod-N nanocrystals in the polyimide membranes,
the peaks in Figure 5.2 increase in intensity from PI-15 to PI-35.
2000 1800 1600 1400 1200 1000 800 600
1780
Tran
smitt
ance
(a.u
.)
Wavenumbers (cm-1)
Sod-N
PI-0
661
661
72099013801720
PI-35
Figure 5.3. IR spectra of samples Sod-N, PI-0 and PI-35.
Figure 5.3 shows the IR spectra of Sod-N, PI-0 and PI-35. For the last two samples,
absorption bands, which correspond to the polyimide structure, are observed at 1780
cm-1 (C=O asymmetric stretching), 1720 cm-1 (C=O symmetric stretching), 1380 cm-1
Chapter 5
103
(C-N stretching), and 720 cm-1 (imide ring deformation); these indicate the successful
chemical imidization of the membranes [154, 262-264]. For the pure Sod-N sample, the
broad band at approximately 990 cm-1 is assigned to the asymmetric stretch (T-O-T, T =
Si, Al), and the adsorption at 661 cm-1 is ascribed to the symmetric stretch (T-O-T)
[84]. The presence of Sod-N in sample PI-35 causes the peaks at around 1000 cm-1 to
broaden in comparison with pure PI-0 film. Furthermore, there are new small peaks
appear for PI-35 at 661 cm-1, which is due to symmetric stretch T-O-T (T=Si, Al)
arising from added Sod-N.
Figure 5.4. SEM images for PI-0, PI-15, PI-25 and PI-35.
Figure 5.4 shows the SEM images for PI-0, PI-15, PI-25 and PI-35. These micrographs
confirm that Sod-N nanocrystals are well dispersed throughout the polyimide matrix at
all loadings of Sod-N. No voids are apparent between the nanocrystals and polyimide,
even at 35-wt% Sod-N where some large-scale surface roughness is evident, which
suggests good bonding between the zeolite and polymer. Other studies also have found
Chapter 5
104
that improving the interaction between zeolite and polymer tends to inhibit formation of
interfacial voids [176, 230-232].
0 100 200 300 400 500 600 700 800
0
20
40
60
80
100
PI-35
PI-0PI-15PI-25
Wei
ght l
oss
(%)
Temperature (°C)
Figure 5.5. TGA curves for samples PI-0, PI-15, PI-25, and PI-35.
The thermogravimetric (TG) curves of pure polyimide and the composite membranes
with different loadings of Sod-N are shown in Figure 5.5; Table 5.1 summarizes the
corresponding thermogravimetric (TG) and differential thermogravimetric (DTG)
results. Under flowing oxygen, the pure polyimide membrane, PI-0, lost 1.6% of its
mass in the temperature range from 30-400 °C. This is due to the loss of residual
organic solvent (DMF has a boiling point of 153 °C) and/or adsorbed water. In the
temperature range from 400-700 °C, the remaining 98.4% of mass was lost, leaving no
residue after the TGA run, which is ascribed to the complete decomposition and
combustion of the polyimide at high temperature [154, 264]. The DTG peak (Td) for the
corresponding mass-loss lies at 573 °C.
The mass losses varied for the composite membranes during heating between 30 °C and
400 °C—3.0%, 2.7% and 3.4% for PI-15, PI-25 and PI-35, respectively—but all the
composites lost more weight than PI-0. This might be due to increased adsorption of
Chapter 5
105
water and/or DMF caused by the hydrophilic Sod-N particles and/or by the presence of
inorganic-organic crosslinked networks after polymerization [264]. However, most of
the weight loss occurs in the temperature range from 400 °C to 700 °C and is 84.5% for
PI-15, 78.0% for PI-25 and 71.4% for PI-35. Interestingly, the Td values for the
composite materials are all higher than that of pure polyimide, and increase with Sod-N
content: 580 °C for PI-15, 595 °C for PI-25 and 600 °C for PI-35. Some previous
research attributed this kind of trend to the interaction between the amino moieties from
inorganic nanoparticles (Sod-N) and the polymer matrix, which can reduce the
movement (increase the rigidity) of the polymer chains and, thus, increase the
decomposition temperature of composite membranes [154, 262, 265].
The residual masses after TG analysis are 12.5%, 19.3% and 25.2% for PI-15, PI-25 and
PI-35, which would correspond to plain sodalite nanocrystals, given that organic
functional groups (i.e., −(CH3)(CH2)3NH2) would have been completely decomposed
and removed by the high temperatures [22]. The results of 29Si-NMR in Section 4.3.2
showed that 3.18 wt% of Sod-N comprised organic functional groups. This allows
recalculation of the actual Sod-N loading of PI-15, PI-25 and PI-35 as 14.8%, 24.7%
and 34.8%, respectively, which are close to the theoretical values.
5.3.2 H2 sorption of sodalite nanocrystals and gas permeation
of membranes
H2 sorption isotherms of amino-functionalized sodalite nanocrystals are shown in
Figure 5.6. It is clear that the temperature has substantial influence on H2 adsorption
capacity of sodalite nanocrystals. At 77 K, H2 adsorptive volume significantly increases
with increasing the adsorption pressure, and it reaches a maximum volume of 26.9
cm3/g. However, at room temperature (298 K), amino-functionalized sodalite
nanocrystals exhibit almost no H2 adsorption as P/Po is raised to 1 (Po = 900 mmHg).
This is due to sodalite cage contraction when the sorption temperature increases from 77
K to 298 K. XRD analysis confirms that the crystallinity in amino-functionalized
sodalite nanocrystals remains unchanged after H2 sorption analysis. According to Ref.
[266], sodalite cage expands and starts to uptake hydrogen at 573 K or above. These
Chapter 5
106
indicate that amino-functionalized sodalite nanocrystals may function as nonporous
nanoparticles in nanocomposite membranes in the gas permeation temperatures.
Figure 5.6. H2 adsorption-desorption isotherms of amino-functionalized sodalite
nanocrystals at 77 K and 298 K.
Table 2 summarizes the permeability values of two pure gases (H2 and N2) and the ideal
selectivity )( 22 NHα for pure polyimide films and composite membranes at three different
temperatures (25 °C, 60 °C and 100 °C). The permeability and ideal selectivity data for
pure polyimide membranes fabricated from BTDA and MDA is comparable to similar
polyimide membranes in the literature [149, 156]. The addition of Sod-N causes the
composite membranes to reduce the permeation of N2, leading to a substantial
improvement in H2/N2 selectivity. This should be attributed to the interfacial effects and
disrupted polyimide chain packing caused by the covalent bonding between Sod-N and
polyimide. The structure of sodalite-polyimide interface is illustrated in Scheme 5.4a.
Sodalite nanocrystals are composed of a crystalline sodalite core and a thin amorphous
aluminosilicate shell with amino-groups (=Si–(CH3)(CH2)3NH2). The thickness of the
amorphous aluminosilicate shell is roughly estimated to be around 2 nm assuming that
all amino-groups are contained in the shell. The high-quality bonding between the
sodalite nanocrystals and the polymer matrix is realized by forming covalent linkers via
the imidization reaction of the amino-groups with the polyimide monomers (Scheme
5.4b). The addition of Sod-N also affects the chain length of polyimide molecules
surrounding Sod-N nanocrystals because the polyimide chains reacting with amino-
groups are terminated. This would increase the rigidity of the polymer chains in the
Chapter 5
107
interfaces and polyimide matrix [194, 195]. These unique structures allow H2 to diffuse
through while reducing the passage of N2 molecules. This explains that the H2
permeability of all composite membranes at 25 °C is slightly higher than that of the
pure polyimide membrane. On the other hand, these data provide strong evidence that
there are no voids are present at the polyimide and sodalite interface in any of the
composite membranes, because such voids would have resulted in a large increase in
permeability of H2 or even N2 [231].
Scheme 5.4. Schematic representation of sodalite-polyimide interfacial structure (a) and
covalent linker between Sod-N and polyamide (b).
When the testing temperature is elevated, there is a subsequent increase in the
permeability of H2 or N2 for the pure-polyimide and the composite membranes. There
was a more significant increase in permeability for the pure polymer with temperature
than was found for the composite membranes, especially for N2 gas. PI-0 has a N2
permeability of 0.036 Barrer at 25 °C, compared with 0.127 Barrer at 100 °C, a 3.5-fold
increase. However, PI-35 showed an increase of only 1.1 times for 2NP between room
Chapter 5
108
temperature and 100 °C. In addition, at 60 and 100 °C, the permeabilities of H2 and N2
for the composite membranes are lower than those for the PI membrane (Table 5.2). As
the temperature increases, the permeabilities of both H2 and N2 increase because of the
increase of the diffusivity and the decrease of the solubility in polyimides [267]. This
result is attributed to the increase in polymer chain rigidity in the composite membranes
with increasing Sod-N loading, and the increase in the permeabilities of both H2 and N2
for the PI membrane is greater than those for the composite membranes.
Table 5.1. DTG and TG results of PI-0, PI-15, PI-25 and PI-35.
DTG TG
Mass loss (%) Sod-N content (%)
Sample Td (°C)
30-400 °C 400-700 °C
Mass residue
after TGA (%) Experimental Theoretical
PI-0 573 1.6 98.4 0 0 0
PI-15 580 3.0 84.5 12.5 14.8 15.0
PI-25 595 2.7 78.0 19.3 24.7 25.0
PI-35 600 3.4 71.4 25.2 34.8 35.0
Table 5.2. Gas permeation results of the PI-0, PI-15, PI-25, and PI-35 membranes.
Permeability (Barrer) Selectivity
)( 22 NHα
25 °C 60 °C 100 °C
Sample
H2 N2 H2 N2 H2 N2
25
°C
60
°C
100
°C
PI-0 6.96±0.11 0.036±0.001 12.04±0.12 0.096±0.001 13.87±0.15 0.127±0.002 193 125 109
PI-15 7.41±0.10 0.033±0.001 10.81±0.09 0.079±0.002 12.74±0.09 0.113±0.001 225 137 113
PI-25 8.05±0.05 0.034±0.001 9.92±0.06 0.056±0.000 11.28±0.16 0.073±0.001 237 177 155
PI-35 8.04±0.09 0.029±0.001 9.86±0.10 0.043±0.001 13.14±0.14 0.062±0.001 277 229 212
The H2/N2 selectivity for PI-0, PI-15, PI-25 as a function of temperature, which are
included in Table 5.2, are shown in Figure 5.7. At 25 °C, the nanocomposite
Chapter 5
109
membranes demonstrate perselectivities of 225, 237, and 277 for PI-15, PI-25, and PI-
35, respectively. These values represent 17%, 23%, and 44% greater ideal selectivity,
respectively, than PI-0.
It is also plain from Figure 5.7 that elevating the temperature lowers the ideal
selectivities of all the membranes, but that increasing the sodalite content considerably
retards the falling-off of gas selectivity from 25 to 100 °C. For instance, PI-0’s
selectivity drops from 193 at 25 °C to 109 at 100 °C, which is a fall of 44%. In contrast,
PI-35 sees a decrease in )( 22 NHα of only 23%. As Sod-N loading increases, the number
of Sod-N terminated increases substantially, affecting the chain configuration beyond
the interfaces; in other words, the interfacial area may be extended. Therefore, the
increased inorganic content in composite materials restricts the thermal motion of the
polymer segments, and thus reduces the decrease in the gas selectivity [209].
20 40 60 80 100100
150
200
250
300
S
elec
tivity
(H2/N
2)
Temeprature (°C)
PI-0 PI-15 PI-25 PI-35
Figure 5.7. Selectivity )( 22 NHα for PI-0, PI-15, PI-25, and PI-35 at different
temperatures (25 °C, 60 °C and 100 °C)
Figure 5.8 shows the apparent activation energy, Ep, of PI-0, PI-15, PI-25 and PI-35 for
the pure H2 and pure N2. It is apparent that all samples have higher values Ep for N2 than
H2, confirming that N2 molecules need more energy to penetrate the membranes than H2
molecules. Compared with composite membranes, pure polyimide polymer (PI-0) has
Chapter 5
110
the highest activation energies—8.5 kJ/mol for H2 and 15.8 kJ/mol for N2. In the
composite membranes, the presence of Sod-N lowers Ep below that of the pure polymer
membranes. For example, PI-15 and PI-25 have Ep values of 6.7 and 4.2 kJ/mol,
respectively, for H2 and 15.2 and 9.5 kJ/mol, respectively, for N2. Interestingly, PI-35
shows an increase in activation energy relative to PI-25 for H2, but not for N2.
Similarly, the H2 permeability for PI-35 increases largely from 8.04 Barrers to 13.14
Barrers as the temperature is increased from 25 °C to 100 °C. In the composite
membranes, gas diffusion requires relatively small segmental motions of polymer
matrix in the packing-disrupted polyimide chains and sodalite–polyimide interfaces,
because they possess relatively more unoccupied free space. When Sod-N loading is
increased to a certain point (e.g., 35%), the overlap of interfacial layers becomes
significant [268]. Such overlapped interfaces favour H2 diffusion, and are more
temperature-dependent in the permeation of small hydrogen molecules. It is clear that
the separation performance of polyimide membrane has been significantly enhanced.
The strategy of forming nanocomposite membranes demonstrated in this work could be
applied to fabricate practical H2 separation membranes by incorporating functional
sodalite nanocrystals into a more permeable polyimide skin layer. It would be
interesting to study water transport property of sodalite-polyimide nanocomposite thin
membranes for potential applications, such as in water/organic solvent separation and
water purification, given that sodalite membranes have been reported to exhibit good
water permeation property [269].
PI-0 PI-15 PI-25 PI-350
2
4
6
8
10
12
14
16
Ep (K
J/m
ol)
Hydrogen Nitrogen
Figure 5.8. Apparent activation energy (Ep) for PI-0, PI-15, PI-25 and PI-35.
Chapter 5
111
5.4 Conclusion
In Chapter 5, organic-functionalized sodalite nanocrystals (Sod-N) and polyimide have
been applied to fabricate nanocomposite membranes. Characterization by SEM showed
that Sod-N can be well distributed with polyimide phase, even at a loading of 35 wt%,
as is confirmed by the FTIR spectroscopy and XRD results. From TG and DTG analysis,
the DTG peaks for corresponding major mass loss increase with the increasing Sod-N
content of the composite, which is attributed to restricted movement of the main chains
arising from the interaction between the amino moieties from inorganic nanoparticles
(Sod-N) and polymer matrix. The gas permeation experiments were performed with two
pure gases, H2 and N2, and the results revealed that H2 permeability was improved,
while N2 permeability decreased. In particular, the PI-35 composite membranes had the
highest selectivity ( )( 22 NHα = 277) and a good permeability (8.04 Barrers) at room
temperature.
Chapter 6
112
Chapter 6 Conclusions & Recommendations for Future Work
6.1 Conclusions
In this thesis, uncrosslinked chitosan hydrogels or crosslinked chitosan hydrogels were
for the first time introduced into zeolite synthesis. The glutaraldehyde-crosslinked
chitosan (GA-CS) hydrogels with three-dimensional network structures were found to
be effective for controlling the growth of zeolite NaA and NaY. The zeolite crystal sizes
were significantly affected by the formulation of silica-containing GA-CS hydrogels
and alkaline solution, and by the aging and heating conditions. The zeolite NaA
nanocrystals with an average size of 148 nm and NaY with an average size of 192 nm
were synthesized in this study. A novel method of using hydrogen peroxide solution
was developed to remove GA-CS hydrogels after zeolite synthesis. TGA results
confirmed that polymer hydrogels were completely removed by this hydrogen peroxide
treatment method. The NaA samples obtained via this method exhibited much higher
crystallinity than those obtained via conventional calcination. This suggested that the
hydrogen peroxide treatment method be preferred for removal of GA-CS hydrogels. In
addition, the zeolite NaA and NaY nanocrystals produced here were readily dispersed in
solvents such as deionized water, and therefore they are useful for applications such as
in the fabrication of zeolite-polymer mixed matrix membranes (MMMs) and
hierarchical porous zeolitic structures. Within uncrosslinked chitosan hydrogels, the
cubes of zeolite A consisting of a thin crystalline shell and an amorphous core are found
grown. It is evident that the formation of cube-like or rectangular core-shell structures
involves particle aggregation and surface-to-core crystallization induced by chitosan
networks. This work provides a new model system for studying complex zeolite
nucleation and growth mechanisms.
To fabricate defect-free mixed matrix membranes (MMMs), organic-functional groups
have also successfully attached onto sodalite nanocrystals through the direct
transformation of organic-functionalized silicalite nanocrystals. XRD and nitrogen
sorption measurements showed that the organic-functionalized sodalite nanocrystals had
Chapter 6
113
high crystallinity and well-grown pore structures. The micropores of the organic-
functionalized sodalite nanocrystals were highly accessible to hydrogen molecules at
low temperature (77 K), though there was a slight reduction of hydrogen adsorption
compared with sodalite nanocrystals without organic groups. Sodalite nanocrystals with
–(CH3)(CH2)3NH2 moieties (denote as Sod-N) showed good dispersibility in all four
solvents (i.e., water, isopropanol, dichloromethane, and dimethylformamide) tested
whereas sodalite nanocrystals with –CH3 groups (denote as Sod-C) were dispersible in
isopropanol, dichloromethane and dimethylformamide, but were agglomerated in water.
The produced organic-functionalized sodalite nanocrystals (Sod-N) were added into the
DMF solvent with polyimide monomers to fabricate mixed matrix membranes (MMMs).
Characterization by SEM showed that Sod-N can be well distributed with polyimide
phase, even at loadings of 35 wt%. From TG and DTG analysis, the DTG peaks for
corresponding major mass loss increase with the increasing Sod-N content of the hybrid,
which was attributed to the interaction between the amino moieties from inorganic
nanoparticles (Sod-N) and polyimide matrix, which restricts the movement of the main
chains. The gas permeation experiments were performed and the results revealed that
the PI-35 mixed matrix membranes (MMMs) had the highest selectivity ( )( 22 NHα = 277)
and a good permeability (8.04 Barrer) at room temperature, which is an exciting finding
from my study.
6.2 Recommendations for future work
1. To date, SDA-free synthesis of zeolite nanocrystals with controllable sizes and
size distributions still remains a challenging task. My study shows that zeolite A
and Y nanocrystals can be prepared by using crosslinked chitosan hydrogel
networks. It is recommended to investigate the possibility of producing other
types of zeolite nanocrystals and synthesis of even smaller nanocrystals by using
similar gel system. Furthermore, it is necessary to further investigate the
mechanism of zeolite nanocrystal nucleation and growth in polymer hydrogels,
which can help design and optimize suitable hydrogel systems for practical
application.
Chapter 6
114
2. Cubic core-shell zeolite A has been produced by using uncrosslinked chitosan
hydrogels. It would be interesting to carry out more TEM study to investigate
crystal growth mechanism in details. Research should be also directed to study
the possibility of producing other types of zeolite in larger pore sizes (e.g.
zeolite Y and X) with a similar core-shell structure. It is possible to form hollow
crystals with large pores by dissolving amorphous core, which may be useful in
catalysis.
3. Sodalite nanocrystals and organic functionalization of sodalite nanocrystals were
synthesized for the first time by using direct transformation method. The
hydrogen adsorption results of the produced particles at 77 K were shown to be
accessible to hydrogen molecules. However, when the temperature was
increased to 25 °C or above that (e.g. 100 °C), there was only a limited
adsorption of hydrogen by Sod or organic-functionalized Sod-N at P/P0 ≈1. It is
recommended that future studies focus on the hydrogen adsorption of sodalite
nanocrystals at higher temperature or pressure to elucidate the accessibility of
SOD cage to hydrogen gas. Furthermore, micro-sized sodalite crystals have been
suggested to be a candidate for hydrogen storage, when at high temperature (e.g.
over 250 °C) in the literature. The nano-sized sodalite crystals are expected to
have higher hydrogen adsorption capacity, since nanocrystals have larger
surface area as compared with micro-sized ones.
4. The mixed matrix membranes (MMMs) by combining organic-functionalized
sodalite nanocrystals and polyimide matrix have been tested via permeation of
two different pure gases, i.e. hydrogen and nitrogen. In practice, it is desirable to
separate hydrogen from other gases, such as carbon dioxide and ammonia.
Hence, it is recommended that the study be extended to the separation of
hydrogen from other pure gases and gas mixtures. Operating conditions such as
pressure and temperature may play an important role in the gas separation
performance of the membranes and they should be investigated in the future.
5. In this study, MMMs with micro-sized thickness were fabricated for
characterization purpose. It is clear that MMMs as a thin layer (e.g. 100 nm)
exhibits higher gas flux and the processability of MMMs is very important for
industries. It is recommended that the strategy of forming MMMs demonstrated
Chapter 6
115
in this work could be applied to fabricate practical H2 separation membranes by
incorporating functional sodalite nanocrystals into a more permeable polyimide
skin layer.
6. It would be interesting to study water transport property of sodalite-polyimide
mixed matrix thin membranes for other potential applications such as in
water/organic solvent separation, and water purification, given that sodalite
membranes have been reported to exhibit a good activated water permeation.
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Appendix-Relevant Publications
137
Appendix-Relevant Publications
Microporous and Mesoporous Materials 116 (2008) 416–423
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Microporous and Mesoporous Materials
journal homepage: www.elsevier .com/locate /micromeso
Zeolite crystallization in crosslinked chitosan hydrogels: Crystal size control andchitosan removal
Dan Li a, Yi Huang a, Kyle R. Ratinac b, Simon P. Ringer b, Huanting Wang a,*
a Department of Chemical Engineering, Monash University, Clayton, Vic. 3800, Australiab Australian Key Center for Microscopy and Microanalysis, The University of Sydney, Sydney, NSW 2006, Australia
a r t i c l e i n f o
Article history:Received 15 January 2008Received in revised form 7 April 2008Accepted 30 April 2008Available online 24 May 2008
Keywords:Zeolite nanocrystalsChitosanHydrogelsHydrothermalHydrogen peroxide
1387-1811/$ - see front matter � 2008 Elsevier Inc. Adoi:10.1016/j.micromeso.2008.04.032
* Corresponding author. Tel.: +61 3 9905 3449; faxE-mail address: [email protected]
a b s t r a c t
For the purpose of controlling zeolite crystal size, crystallization of zeolite NaA and NaY in glutaraldehydecrosslinked chitosan (GA-CS) hydrogels was studied in this paper. The zeolite crystals were produced bypenetration of Na2O–Al2O3–H2O alkaline solution into GA-CS hydrogels filled with colloidal silica, fol-lowed by hydrothermal treatment and removal of GA-CS hydrogels. We systematically investigatedthe effects of the synthesis parameters – including the amounts of silica, chitosan, and glutaraldehyde,and the aging and heating times – on the size, size distribution and crystallinity of the particles. A hydro-gen peroxide treatment method was shown to be an effective way for removing GA-CS hydrogels, therebyavoiding the conventional calcination step. X-ray diffraction (XRD), light scattering, scanning electronmicroscopy (SEM), thermogravimetric analysis (TG), and N2 sorption were used to characterize the zeo-lite samples. This work showed that GA-CS is a promising space-confinement medium for the synthesis ofzeolite nanocrystals with tunable crystal sizes and excellent dispersibility.
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1. Introduction
There has been considerable interest in confined-space synthe-sis of zeolite nanocrystals since the principle was first reported byMadsen and Jacobsen in the late 1990s [1]. To date, a variety ofadditives such as carbon blacks and polymer hydrogels have beenused to confine zeolite crystallization. As a class of soft space-con-finement additives, polymer hydrogels comprise three-dimen-sional networks that are created via physical or chemicalcrosslinking [2, 3], which can be readily introduced into zeolitesynthesis due to good compatibility between zeolite precursorsand polymer gels [4]. We have recently demonstrated the con-trolled synthesis of zeolite crystals in chemically crosslinked poly-acrylamide hydrogel [4] and thermoreversible methyl cellulosehydrogels [5]. The crystal sizes of SAPO-34 molecular sieves weresubstantially reduced by forming crosslinked polyacrylamidehydrogel from the water-soluble organic monomers acrylamide(AM) and N,N0-methylenebisacrylamide (MBAM), followed by avapor phase transport process [4]. However, the synthesizedSAPO-34 nanocrystals exhibited a very poor dispersibility in sol-vents. Similarly, NaA (20–180 nm in size) and NaX (10–100 nmin size) nanocrystals were synthesized by employing thermore-versible methylcellulose hydrogels to confine crystal growth [5].The zeolite nanocrystals were easily collected by washing awaythe water-soluble methylcellulose at room temperature, and they
ll rights reserved.
: +61 3 9905 5686.u (H. Wang).
were highly dispersible in water and ethanol. Given the successof these techniques, it would be considerable interest to further ex-plore the feasibility for the synthesis of zeolite nanocrystals withcontrollable sizes and size distributions in other polymerhydrogels.
In this paper, therefore we report attempts to control the syn-thesis of zeolite nanocrystals in the system of Na2O–SiO2–Al2O3–H2O by using crosslinked chitosan hydrogels. Chitosan is a partiallydeacetylated polymer of chitin, which is found in a wide range ofnatural sources, such as crab, lobster and shrimp shells. Its ali-phatic primary amino groups are regularly distributed along thepolymer backbones, and can be crosslinked to form more rigidpolymer networks [6–9]. The crosslinked chitosan hydrogels havebeen studied for various applications such as in pervaporation sep-aration through chitosan [10] or chitosan-zeolite membranes [11],enzyme immobilization [12] and cationic specimen transportation[13], controlled ingredient-release [14, 15], environmental applica-tions [16] and fuel cells [17]. In the present work, zeolite crystalli-zation in glutaraldehyde crosslinked chitosan hydrogels wassystematically investigated to determine the effects of differentcompositions of the synthesis mixture (ratios of chitosan to silicato glutaraldehyde), and the duration of aging and heating. Unlikeother polymer hydrogels, chitosan is only soluble in an acidic solu-tion, and does not dissolve in an alkaline zeolite synthesis gel.Therefore, a two-step method involving the formation of silica-filled crosslinked chitosan hydrogel and the subsequent penetra-tion of Na2O–Al2O3–H2O alkaline solution was developed to forma sodium aluminosilicate gel inside the crosslinked chitosan hydro-
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gel. After zeolite synthesis, hydrogen peroxide was employed todegrade chitosan to retrieve zeolite crystals, and a comparison be-tween hydrogen peroxide degradation and high-temperature calci-nation was made.
Scheme 1. Synthesis of zeolite crystals within crosslinked chitosan hydrogels (GA-CS).
2. Experimental section
2.1. Synthesis of zeolite LTA (NaA)
Firstly, 0.6–1.5 g of chitosan (average molecular weight120,000 g/mol, �80% deacetylation, Sigma-Aldrich, denoted CS)was dissolved in 21 g of 1 M acetic acid (Sigma-Aldrich). Theresulting solution was stirred at room temperature for 1 h and thenleft overnight without stirring, after which 2.0–4.0 g of colloidalsilica (HS-30, 30%, Sigma-Aldrich) were added. A given amount(0.6–5.0 g) of glutaraldehyde (50%, Sigma-Aldrich, denoted GA)was added into the CS-silica solution, and left undisturbed at roomtemperature for 2 h, resulting in crosslinked chitosan hydrogel (de-noted GA-CS). Therefore, the silica-filled crosslinked chitosan (GA-CS) hydrogels were synthesized from molar compositions in therange 0.005–0.0125CS:10–20SiO2:1.88–25 glutaraldehyde(GA):21acetic acid (HAc):1243–2499H2O, corresponding to masscompositions of 0.6–1.5CS:0.6–1.2SiO2:0.4–5.0GA:1.3HAc:22.4–24.7H2O.
Secondly, an alkaline solution was prepared by dissolving 5.56 gof NaOH (99%, Merck) in 20.00 g of deionized water, with subse-quent addition of 2.45 g of NaAlO2 (anhydrous, Sigma-Aldrich) dur-ing stirring. The molar composition of the alkaline solution was7.7Na2O:1.0Al2O3:111.0H2O. This alkaline solution was introducedinto the silica-filled crosslinked chitosan hydrogel with a finalmolar composition of 0.005–0.0125CS:10–20SiO2:1.88–25GA:21-HAc:80Na2O:10Al2O3:2396–2455H2O, and aged for 12–72 h atroom temperature. After aging, the gel was removed from the alka-line solution, transferred to a sealed polypropylene bottle and thenheated at 90 �C for 1, 3 or 6 h to allow zeolite crystallization. Tomake a comparison, another sample prepared without aging washeated at 90 �C for 3 h in the presence of the alkaline solution.
2.2. Synthesis of zeolite FAU (NaY)
In this case, the synthesis of silica-filled hydrogels was per-formed from a system with a molar composition of0.01CS:17.5SiO2:12.5GA:21HAc:1302H2O. The same chemicals forthe synthesis of NaA crystals described earlier were used. Typically,1.2 g of CS was dissolved into 21 g of 1 M HAc. As with NaA, thesolution was stirred at room temperature for 1 h, and then leftovernight, after which 3.5 g of colloidal silica was added, and2.5 g of GA was added to form GA-CS. The alkaline solution wasprepared as follows: 4.14 g of NaOH was dissolved in 25.83 g ofdeionized water, with subsequent addition of 0.75 g of NaAlO2 dur-ing stirring. The molar composition of alkaline solution was17.3Na2O:1Al2O3:455.6H2O. The solution was stirred for 0.5–1 huntil it became clear and then it was introduced into the cross-linked chitosan gel system with a molar composition of0.01CS:17.5SiO2:12.5GA:21HAc:55Na2O:3.18Al2O3:2769H2O, andallowed to age at room temperature for 36 h. After aging, the gelwas removed from the alkaline solution, transferred to a sealedpolypropylene bottle, and then heated at 90 �C for 5 h.
2.3. Removal of crosslinked chitosan hydrogels
The heat-treated gels, which contained zeolites, were repeat-edly washed with deionized water until a pH of less than 8 was at-tained. Approximately 3 g of hydrogel was stirred into 150 mL of10% H2O2 solution and then heated at 80–90 �C for 1–2 h. The zeo-
lite crystals were retrieved by high-speed centrifugation and re-peated washing with deionized water; these were dried at 60 �C.For comparison, gels also were calcined to remove the crosslinkedCS. After washing, the zeolite-containing gels were dried at 80 �Covernight, ground by hand using a mortar and pestle, and calcinedat 550 �C under air for 2 h at an initial heating rate of 2 �C min�1.
2.4. Characterization
Scanning electron microscopy (SEM) images were taken with aJSM-6300 F microscope (JEOL). The particle size distributions forzeolite crystals were determined by manual measurement of 300crystals for each sample from the SEM images with Adobe Photo-shop software. Elemental Si/Al ratios of samples were determinedby energy dispersive X-ray spectroscopy (EDXS) on the JSM-6300Fmicroscope. X-ray diffraction (XRD) patterns were recorded on aPhilips PW1140/90 diffractometer with CuKa radiation (25 mAand 40 kV) at a scan rate of 2�/min and a step size of 0.02�. Thermo-gravimetric analysis (TGA, Perkin Elmer, Pyris 1 analyzer) was per-formed at a heating rate of 5 �C/min to 700 �C in oxygen with aflow rate of 15 cm3 min�1. Nitrogen adsorption–desorption exper-iments were performed at 77 K with a Micrometritics ASAP2020MC analyzer. The NaA sample and NaY sample were degassedat 673 K for 24 h, and 623 K for 4 h, respectively, prior to analysis,and the specific surface areas were calculated according to the Bru-nauer–Emmett–Teller (BET) method. To study the dispersibility ofzeolite nanocrystals, the particle size distributions of colloidal zeo-lite suspension were analyzed by light scattering with a MalvernMastersizer 2000 analyzer. Approximately 12–15 mL samples ofcolloidal zeolite suspension were prepared for this purpose by dis-persing 50 mg of each sample into 50 mL of deionized water duringultrasonication.
3. Results and discussion
A schematic diagram for the formation of zeolite nanocrystals inGA-CS hydrogels is shown in Scheme 1. A colloidal silica solution isdispersed in the solution of CS and acetic acid. When the cross-linker (GA) is added, the amino groups from the backbones ofchitosan are crosslinked [9], which causes chitosan solution tosolidify into yellow gels that contain colloidal silica. To form alumi-nosilicate zeolite gels, the alkaline solution is added into the yellow
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gel. During the aging, alkaline solution penetrates into the GA-CSgel and reacts with the entrapped silica. Aluminosilicate gels crys-tallize during the subsequent heating, producing zeolite crystalswithin the crosslinked chitosan hydrogel networks. To removethe chitosan hydrogels, hydrogen peroxide is added to solubilizecrosslinked chitosan by degradation of the network structure. Thenthe zeolite crystals are readily collected through high-speed centri-fugation and repeated washing (Scheme 1).
3.1. Effect of the amount of SiO2
Fig. 1 shows the XRD patterns of the samples prepared with mo-lar compositions of 0.01CS:ySiO2:12.5GA:21HAc:(1166 + 8y)H2O(y = 10-20). They are denoted A-10SiO2, A-12.5SiO2, A-15SiO2, A-17.5SiO2, and A-20SiO2, respectively. From Fig. 1, A-10SiO2 appearsto be amorphous, and A-12.5SiO2 exhibits very low crystallinity,whereas A-15SiO2, A-17.5SiO2, and A-20SiO2 are pure zeolite A.Interestingly, A-17.5SiO2 exhibits the smallest average particle size(148 nm), which is much smaller than the mean particle sizes of325 nm and 239 nm of A-20SiO2 and A-15SiO2, respectively. It iswell known that organic additives play an important role in zeolitenucleation and growth [18, 19]. Previous studies have found thatthe addition of water-soluble surfactants (e.g. sodium dodecyl sul-fate, sodium dioctylsulfosuccinate and cetyltrimethylammoniumbromide) and organic polymers (e.g. poly(ethylene glycol)) in thezeolite gel dramatically shortened prenucleation and nucleationperiods and accelerated crystal growth. [18] Some recent mathe-matical [20, 21] modeling and experimental results [22] haveshown that the zeolite nucleation takes place at the interface be-tween the solution and the gel solid by adsorption and rearrange-ment of soluble precursor. Our experimental results could possiblybe explained by the effect of the ratio of SiO2 to CS on the rates ofnucleation and growth. When the ratio of SiO2 to CS is high at2000:1, both initial nucleation rate and subsequent growth rateare presumably high due to the high concentration of aluminosili-cate gel, ultimately leading to the larger crystal sizes. As the ratio ofSiO2 to CS is lowered to 1750:1, the initial nucleation rate mightdrop slightly, accounting for far smaller crystals. When the ratioof SiO2 to CS decreases further to 1250:1, both nucleation rateand growth rate would significantly decrease, resulting in decreasein the number of nuclei formed in the system, but an overall in-crease in final particle size. As expected, if the ratio of SiO2 to CSbecomes too low (e.g. 1000), the amount of silica is insufficientfor zeolite crystallization.
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Fig. 1. XRD patterns of the samples prepared with molar compositions of0.01CS:ySiO2:12.5GA:21HAc:(1166 + 8y)H2O (y = 10 � 20) under the same aging(36 h) and heating (90 �C for 3 h) conditions: (a) A-10SiO2; (b) A-12.5SiO2; (c) A-15SiO2; (d) A-17.5SiO2 and (e) A-20SiO2. All of the samples were collected afterH2O2 treatment.
The SEM image, particle size distributions, and nitrogen sorp-tion isotherm of A-17.5SiO2 are shown in Fig. 2. The SEM image(Fig. 2a) reveals that the crystals exhibit irregular shapes. The par-ticle sizes determined by SEM range from 100 nm to 200 nm, aver-aging 148 nm (Fig. 2b), which is slightly smaller than their meanparticle size 174 nm measured by light scattering (Fig. 2c). Thesimilarity of these size distributions and mean sizes suggests thatthe produced particles were well dispersed in water. Furthermore,the suspension formed by dispersing the zeolite particles in wateris stable under lab condition for at least one week. In the nitrogenadsorption–desorption isotherm (Fig. 2d), the amount of nitrogenadsorbed in the sample is very low at low relative pressures, andsubstantially increase at high relative pressures (e.g. P/Po > 0.8).It is clear that the nitrogen adsorption arises from the external sur-faces of nanocrystals because the micropores of well-crystallizedzeolite LTA crystals are inaccessible to nitrogen molecules at the li-quid nitrogen temperature (77 K) [23]. The BET surface area is cal-culated to be 26.2 m2/g, and this further supports a highcrystallinity of the sample.
3.2. Effect of the amount of chitosan
Fig. 3 shows the XRD patterns of the samples produced withmolar compositions of xCS:17.5SiO2:1250xGA:21HAc:(1233 +6944x)H2O (x = 0.005-0.0125), and clearly the crystallinity of sam-ples declines as the amount of chitosan increases. For instance,when x = 0.0125 (Fig. 3a), the sample (A-0.0125CS) exhibits verylow crystallinity. This is probably due to an increase in the densityof the polymer networks with the CS concentration, resulting insignificant reduction of diffusion of zeolite precursors, and henceslow crystallization. The same argument can be applied to the dif-ference in particle size. The SEM results indicate that the crystalsize of zeolite A decreases with increasing amounts of CS; averagesizes are 399 nm, 341 nm and 148 nm for samples prepared withx = 0.0050, 0.0075, and 0.01, respectively. Again this can be ex-plained by the decrease in local concentration of aluminosilicateavailable within the chitosan hydrogels as the amount of chitosanincreases.
3.3. Effect of the amount of glutaraldehyde (GA)
The amount of GA was varied in the CS hydrogels to study theeffect of crosslinking. The samples were prepared from crosslinkedCS hydrogels with a molar composition of 0.01CS:17.5SiO2:zGA:21-HAc:(1233 + 6z)H2O (z = 1.88-25) under the same aging (36 h) andheating (90 �C for 3 h) conditions, and H2O2 treatment. The GA con-centration in the crosslinking gel was expressed as the molar ratio‘‘GA molecules:amino groups from chitosan” (GA:NH2), which wasvaried at 0.3GA:1.0NH2 (z = 1.88), 1.0GA:1.0NH2 (z = 6.25),2.0GA:1.0NH2 (z = 12.5), 4.0GA:1.0NH2. (z = 25). The samples ob-tained were denoted A-0.3GA, A-1.0GA, A-2.0GA and A-4.0GA,respectively. Fig. 4 displays the SEM images for A-0.3GA, A-1.0GA, A-2.0GA and A-4.0GA. It is clear that A-0.3GA has a widesize distribution ranging from 150 nm to over 500 nm (Fig. 4a),whereas A-1.0GA exhibits a narrower size distribution rangingfrom 100 nm to approximately 340 nm (Fig. 4b). A-2.0GA and A-4.0GA exhibit still smaller sizes and narrower size distributions(Fig. 4c and d).
Fig. 5 compares the particle size distributions measured by SEMand light scattering for A-2.0GA and A-4.0GA. In terms of SEM-de-rived particle size distributions, both A-2.0GA and A-4.0GA havesimilar particle size distributions, with average sizes of 148 nmand 147 nm, respectively (Fig. 5a and c). Light scattering measure-ments indicate that both A-2.0GA and A-4.0GA are well dispersedin deionized water, and that their mean particle size is 174 nmfor A-2.0GA (Fig. 5b) and 167 nm for A-4.0GA (Fig. 5d).
Fig. 2. (a) SEM image, (b) particle size distribution determined by SEM, (c) particle size distribution measured by light scattering, and (d) N2 sorption isotherm of the sampleA-17.5 SiO2.
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Fig. 3. XRD patterns of the samples prepared with molar compositions ofxCS:17.5SiO2:1250xGA:21HAc:(1233 + 6944x)H2O (x = 0.005-0.0125) under thesame aging (36 h) and heating (90 �C for 3 h) conditions: (a) A-0.0125CS; (b) A-0.01CS; (c) A-0.0075CS; (d) A-0.005CS. All of the samples were collected after H2O2
treatment.
Fig. 4. SEM images of the particles produced with different amounts of added GA:(a) A-0.3GA; (b) A-1.0GA; (c) A-2.0GA; (d) A-4.0GA.
D. Li et al. / Microporous and Mesoporous Materials 116 (2008) 416–423 419
These results suggest that the degree of crosslinking of chitosanhydrogel has a strong effect on the crystal size. As the amount ofGA increases, the three-dimensional 3 D networks of the hydrogelsbecome denser and more uniform [24]. The higher degree of cross-linking leads to lower degree of swelling, thereby decreasing thediffusion of zeolite precursors in solution [25]. Therefore,crosslinked chitosan gels with more GA provide more rigid, con-fined-spaces for zeolite crystallization and less material for crystal-lization, resulting in smaller and more uniform zeolitenanocrystals. A-2GA and A-4GA exhibit similar morphologies andparticle size distributions because, as has been pointed out in theliterature [25], there is little change in the degree of crosslinkingat high concentrations of GA and, therefore, little change in thecrystallization environment.
3.4. Effect of aging time
The crosslinked chitosan gels with a molar ratio of0.01CS:17.5SiO2:12.5GA:21HAc:1302H2O were aged for differentperiods of 12, 36 or 72 h, and then heated at 90 �C for 3 h. Theresulting samples were denoted A-12h, A-36h and A-72h, respec-tively. Fig. 6 shows the XRD patterns of A-12h, A-36h and A-72h.A-12h possesses a LTA crystal phase with a very low crystallinity
Fig. 5. Particle size distributions determined by SEM (a and c) and by light scattering (b and d) for A-2.0GA (a and b) and A-4.0GA (c and d).
Fig. 7. SEM images of the particles obtained from the chitosan gels after (a) 36 haging (A-36h), and (b) 72 h aging (A-72h).
420 D. Li et al. / Microporous and Mesoporous Materials 116 (2008) 416–423
(Fig. 6a). As the aging time increases, the crystallinity of sample in-creases (Fig. 6b and c). This is because the penetration of the alka-line solution through the crosslinked chitosan hydrogels isessential for producing aluminosilicate gels via reaction with silicaand longer aging times allow greater penetration. In addition, long-er aging allows more uniform nucleation to occur through the gelmatrix, which assists the formation of zeolite crystals with smallsizes and narrow size distributions. The SEM images of A-36hand A-72h are shown in Fig. 7.
Fig. 7a exhibits that the sample A-36h has particle sizes from100 to 200 nm, averaging 148 nm as mentioned above. When theaging period is extended to 72 h (sample A-72h), the resultantNaA particles have larger sizes ranging from approximately 200to 350 nm (Fig. 7b). This is probably due to the swelling behaviorof crosslinked chitosan, which is enhanced during the long expo-sure time to alkaline solution [26]. Moreover, it is possible thatthe crosslinked chitosan may partly degrade after days of exposurethese highly alkaline conditions [25]. As a result, the crosslinkedchitosan can adsorb more precursor Na and Al species and provideslarger spaces for the further growth of zeolite particles.
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Fig. 6. XRD patterns of the samples obtained from the crosslinked chitosan gelsafter (a) 12 h aging (A-12h), (b) 36 h aging (A-36h), and (c) 72 h aging (A-72h).
For further comparison, sample A-0h was made by directlyheating a gel after the addition of alkaline solution without anyaging period. Fig. 8 and Fig. 9 show the XRD patterns and SEMimages, respectively, for samples A-0h and A-36h. A-0h appearsto have a higher degree of crystallinity (as seen in greater peak
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Fig. 8. XRD patterns of the crystals in samples (a) A-0h (with alkaline solution) and(b) A-36h (without alkaline solution).
Fig. 9. SEM images of the crystals in samples (a) A-0h and (b) A-36h.
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Fig. 10. XRD patterns of the samples (a) A-1h, (b) A-3h, and (c) A-6h.
Fig. 11. SEM images of the particles produced after heating for (a) 3 h (A-3h) and(b) 6 h (A-6h).
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Fig. 12. XRD patterns of (a) A-H2O2 and (b) A-cal.
D. Li et al. / Microporous and Mesoporous Materials 116 (2008) 416–423 421
heights) than A-36h (Fig. 8b), as well as the presence of someamorphous material (Fig. 8a). The SEM images of A-0h (Fig. 9a) ex-hibit generally coarser particles with a broad size distribution –some particles even exceed 500 nm – and a mean particle size of380 nm. In contrast, A-36h possesses a far more uniform particlesize distribution (Fig. 9b) with a smaller average size of 148 nm.
The reasons for the difference in particle sizes can be explainedby the extent and uniformity of penetration of alkaline solutions.After the crosslinking reaction with GA, the CS hydrogels entrapcolloidal silica particles within their networks. To produce zeolitecrystals, an aluminosilicate gel of the desired composition needsto be formed by diffusion of the alkaline solution throughout theGA-crosslinked CS hydrogel network, and subsequent reactionwith colloidal silica. Without aging (i.e., sample A-0h), the largecompositional gradients of aluminosilicate gel exist during heating,such that the zeolite nucleation and growth can only start from theoutside of the polymer hydrogel surfaces, resulting in non-uniformgrowth. Thus, during the limited heating time of 3 h, the poorlydistributed precursor material deep within the polymer gel cannot fully crystallize into zeolite, and this unconverted is the amor-phous phase found in the XRD pattern (Fig. 8a). However, aging atroom temperature for 36 h (sample A-36h) allows the alkalinesolution to be evenly distributed throughout the crosslinkedhydrogels, leading to an aluminosilicate gels with a uniform com-position. After aging, the polymer hydrogels that incorporate thealuminosilicate gel are removed from the solutions for heating,which helps prevent the polymer hydrogels from over-swelling.On the other hand, during hydrothermal reaction, the wet gelsmay also slightly shrink due to water evaporation; presumably thismakes confined-space growth more effectively [27].
3.5. Effect of heating time
To study the effect of heating time, the molar composition ofthe gels was fixed at 0.01CS:17.5SiO2:12.5GA:21HAc:1302H2Oand all the synthesis gels were aged for 36 h and then heated at90 �C. The heating time was varied from 1, 3 or 6 h, and the corre-sponding as-synthesized samples were denoted A-1h, A-3h, and A-6h, respectively. The XRD patterns (Fig. 10) indicate that there is nocrystalline material formed after 1 h heating, whereas zeolite Acrystals are produced once the heating period is extended to 3 h(A-3h) or 6 h (A-6h). Fig. 11 shows the SEM images of A-3h andA-6h. It can be seen that the zeolite A which are produced under6 h hydrothermal treatment has larger particle sizes than thoseproduced during 3 h of heating. This difference might be attributedto a combination of the flexibility, interconnected pore channels,and large pore sizes (e.g. a few microns) of polymer hydrogel net-works. Therefore, optimized hydrothermal conditions are requiredfor controlled synthesis of zeolite nanocrystals with a narrow sizedistribution.
3.6. Comparison between the treatment of H2O2 and conventionalcalcination
In this study, we applied a novel method to remove the confin-ing polymer network by degradation of crosslinked chitosanhydrogels in a hydrogen peroxide solution. H2O2 easily decom-poses to form the highly reactive hydroxyl radical (HO�), especiallyunder heating. The hydroxyl radical attacks polymer hydrogels,degrading the crosslinked structure and chitosan molecules [28–32]. For comparison, high-temperature calcination, which is a con-ventional method for removing organic agents, was also used. Thesample for this comparison was crosslinked chitosan with zeoliteA, which was produced from gels with a molar ratio of0.01CS:17.5SiO2:12.5GA:21HAc:1302H2O, which were aged for
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Fig. 13. TG curves of the samples after the treatment of hydrogen peroxide: (a)plain crosslinked chitosan (GA-CS), (b) A-cal, and (c) A-H2O2.
422 D. Li et al. / Microporous and Mesoporous Materials 116 (2008) 416–423
36 h in alkaline solution and heated for 3 h at 90 �C. The sampletreated by hydrogen peroxide is denoted as A-H2O2 and that trea-
Fig. 14. (a) XRD pattern, (b) SEM image, and (c) particle size distribution by SEM and
Fig. 15. (a)TG curve and (b) nitrogen adsorption–desorption isotherm of
ted by calcination as A-cal. Fig. 12 compares the XRD patterns ofpure particles treated with H2O2 (Fig. 12a) and calcination at550 �C for 2 h (Fig. 12b). Clearly, the zeolite A crystals retain great-er crystallinity under the hydrogen peroxide treatment than thoseobtained from high-temperature calcination.
Fig. 13 shows the thermogravimetric (TG) curves of plain cross-linked chitosan, A-cal and A-H2O2. Under flowing pure oxygen,there is a continuous mass loss from 100 �C to 540 �C for the plaincrosslinked chitosan (Fig. 13a). Its total mass loss reaches 100%after 540 �C. A-H2O2 (Fig. 13c) has a total mass loss of approxi-mately 11% occurring, which is mainly attributed to loss of thestructural water from the zeolites as well as physically adsorbedwater. This confirms that the polymer hydrogels were completelyremoved by the hydrogen peroxide treatment method. For A-cal,however, there is another mass loss after 400 �C (Fig. 13b) in addi-tion to the loss of adsorbed and structural water at around 100 �C.This suggests that the crosslinked chitosan was not completelyburned off during calcination. Given the greater crystallinity re-tained and cleaner removal of the hydrogel, the hydrogen peroxidetreatment is clearly the preferred method for removal of GA-CSafter hydrothermal synthesis.
(d) particle size distribution by light scattering of zeolite FAU (NaY) nanocrystals.
zeolite Y samples obtained after treatment with hydrogen peroxide.
D. Li et al. / Microporous and Mesoporous Materials 116 (2008) 416–423 423
3.7. Synthesis of FAU nanocrystals
FAU nanocrystals were synthesized in GA-crosslinked CS hydro-gels by simply varying the compositions of the alkaline solutionand heating times. Fig. 14 displays the XRD pattern, SEM image,and particle size distributions (from SEM and light scattering) ofthe FAU nanocrystals. The Si/Al ratio of the synthesized crystalswas determined to be 2.07:1.00 by the EDXS analysis, suggestingthe nanocrystals are zeolite Y. The nanocrystals have a high crys-tallinity and a narrow size distribution. Their average crystal sizemeasured by SEM is 192 nm; the mean particle size from showsexcellent agreement at 193 nm. This confirms that the zeoliteNaY produced in this way can be well dispersed in deionized water.
Fig. 15 shows TG and nitrogen adsorption–desorption isothermof the zeolite Y nanocrystals. A mass loss of approximately 15% oc-curs, which is due mainly to loss of structural water from the zeo-lites. This confirms that no polymer molecules remain in the voidsof zeolites, a conclusion that is supported by the nitrogenadsorption–desorption isotherm in Fig. 15b. The sample exhibitsa much higher nitrogen adsorption capacity than the NaA samplesbecause NaY has larger, nitrogen accessible pores [33]. The BETspecific surface area of zeolite Y nanocrystals is calculated to be602.2 m2/g.
4. Conclusion
We have shown that glutaraldehyde crosslinked chitosan (GA-CS) hydrogels with three-dimensional network structures wereeffective for controlling the growth of zeolite NaA and NaY. Thezeolite crystal sizes were significantly affected by the formulationof silica-containing GA-CS hydrogels and alkaline solution, and bythe aging and heating conditions. The zeolite NaA nanocrystalswith an average size of 148 nm and NaY with an average size of192 nm were synthesized in this study. A novel method of usinghydrogen peroxide solution was developed to remove GA-CShydrogels after zeolite synthesis. TGA results confirmed that poly-mer hydrogels were completely removed by this hydrogen perox-ide treatment method. The NaA samples obtained via this methodexhibited much higher crystallinity than those obtained via con-ventional calcination. This suggested that the hydrogen peroxidetreatment method be preferred for removal of GA-CS hydrogels.In addition, the zeolite NaA and NaY nanocrystals produced hereare readily dispersed in common solvents, and therefore theymay be useful for applications such as in the fabrication of zeo-lite-polymer composite membranes and hierarchical porous zeo-litic structures.
Acknowledgments
This work was supported by the Australian Research Council(DP0452829) and by Monash University. The technical assistancefrom staff at the Monash Center for electron microscopy is grate-fully acknowledged. H.W. thanks the Australian Research Councilfor the QEII Fellowship.
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Zeolite CrystallizationDOI: 10.1002/anie.200802823
Cubes of Zeolite A with an Amorphous Core**Jianfeng Yao, Dan Li, Xinyi Zhang, Chun-Hua Kong, Wenbo Yue, Wuzong Zhou, andHuanting Wang*
The syntheses of zeolites involve very complex nucleation andgrowth processes. During the past decade, significant progresshas been made towards understanding zeolite crystallizationmechanisms. This progress has been made possible byadvanced analytical techniques, such as high-resolution trans-mission electron microscopy (HRTEM), small-angle X-rayscattering, and atomic force microscopy.[1–5] A number ofzeolite growth mechanisms were proposed based on therespective synthesis of the zeolites. For instance, by monitor-ing the crystallization of silicalite-1 from silica sols intetrapropylammonium ion (TPA) at room temperature, anoriented aggregation mechanism was proposed.[4] In thegrowth mechanism of zeolite A evolving from the nucleiinside the amorphous gel, the particles gradually grow intolarger crystals by consuming the surrounding amorphousgels.[2] The gel was formed by using aluminosilicate solutionsand tetramethylammonium hydroxide as the structure-direct-ing agent (SDA).[2] For zeolite A formation, evidences ofnucleation at the solid–liquid interface of the gel cavities werealso found in sodium aluminosilicate gels without organicSDA.[3] In addition, a reversed crystal growth process fromthe surface to the core of nanocrystallite aggregates wasobserved in the crystal growth of zeolite analcime icosite-trahedra.[6] These studies have undoubtedly provided newinsights into zeolite crystallization processes.
As non-structure-directing agents, organic polymers havesignificant effects on zeolite nucleation and growth. Theconfinement of sodium aluminosilicate zeolite gels in ther-moreversible methylcellulose hydrogels resulted in zeolite A
and X nanocrystals under hydrothermal treatment.[7] Cross-linked polyacrylamide hydrogels was used to reduce SAPO-34 crystal sizes in vapor-phase transport synthesis.[8] In allthese cases, the small crystal sizes is due to space confinementof the polymer hydrogel networks. Hollow sodalite spheresand zeolite A crystals were also synthesized hydrothermallyin the presence of crosslinked polyacrylamide hydrogels. Itwas suggested that the scaffolds of polyacrylamide hydrogelswere the preferential sites for zeolite nucleation, andpromoted the direction of nanoparticle aggregation subse-quent to the surface-to-core growth.[9] These results suggestthat the roles of polymer hydrogels in zeolite synthesis arecomplex, and syntheses of the zeolites depend on themicrostructure of the polymer hydrogels and the interactionbetween the polymer chains and the zeolite gels.
Herein we report the formation of cubes of zeolite A witha single crystalline shell and an amorphous core by in-situcrystallization of sodium aluminosilicate gel inside thepolymer networks of uncrosslinked chitosan hydrogel. Thiswork provides further direct evidence for the surface-to-corereversed-growth mechanism. Chitosan is a biopolymerderived from chitin that is found in a wide range of naturalsources, such as crab, lobster, and shrimp shells. Chitosan,containing abundant amino and hydroxy groups, was used asthe orientation-directing matrix for the synthesis of b-oriented TS-1 films.[10] Glutaraldehyde-crosslinked chitosan(GA-CS) hydrogels were recently used to control zeolitecrystallization, and thus zeolite A and Y nanocrystals weresynthesized.[11] It is noted that chitosan is only soluble in anacidic aqueous solution, and the resulting chitosan solutionturns into a polymer hydrogel when an alkaline solutionpenetrates through the gel. Therefore, for the synthesis ofcore-shell cubes of zeolite A, a two-step process, involving thedispersion of silica in a chitosan acidic solution and subse-quent penetration of Na2O/Al2O3/H2O alkaline solution wasemployed to form a sodium aluminosilicate gel inside theuncrosslinked chitosan hydrogel.
XRD pattern (Figure 1a) indicates the as-synthesizedsample has the structure of zeolite A. SEM image (Figure 1b)shows cube-like crystals with a particle size of 0.5–1.5 mm.This morphology with six {100} facets is typical for zeolite A,which has a cubic structure with the unit cell parameter a =
2.461 nm, and space group Fm3c. The characteristic poly-hedron normally indicates a single crystal property ofzeolite A. According to the classic crystal growth theory,crystals normally develop from nuclei and the appearance ofthe facets is due to the differences in their growth rate.[12–15]
TEM confirms the cube-like or rectangular morphologyof the samples. Figure 2a shows a TEM image of a typicalrectangular particle of zeolite A with the corresponding
[*] Dr. J. F. Yao,[+] D. Li, Dr. X. Y. Zhang, Dr. H. T. WangDepartment of Chemical EngineeringMonash University, Clayton, Victoria 3800 (Australia)Fax: (+ 61)3-9905-5686E-mail: [email protected]
W. B. Yue, Dr. W. Z. ZhouSchool of Chemistry, University of St. Andrews,St. Andrews, Fife KY16 9ST (United Kingdom)
Dr. C. H. (Charlie) KongElectron Microscope UnitUniversity of New South Wales, Sydney, NSW 2052 (Australia)
[+] Present address:State Key Laboratory of Materials-Oriented Chemical Engineeringand College of Chemistry and Chemical EngineeringNanjing University of Technology, Nanjing 210009 (P.R. China)
[**] This work was supported by the Australian Research Council (GrantNo.: DP0452829). H.W. thanks the Australian Research Council forthe QEII Fellowship. W.Z. thanks University of St Andrews for anEaStChem studentship to W.Y.
Supporting information for this article is available on the WWWunder http://dx.doi.org/10.1002/anie.200802823.
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selected area electron diffraction (SAED) pattern, which is astandard single crystal diffraction pattern viewed down the[100] zone axis. Many particles were examined, and the singlecrystalline nature of zeolite A was observed in each casewithout evidence of any polycrystallinity and twin defects.However, the image contrast implies a core–shell structure, inwhich the core appears to be disordered. Under the electronbeam of the microscope, the disordered core of the zeolite Areduced in volume and separated itself from the shell in amatter of few minutes. The shell remained intact, whichclearly appeared as a rectangle with a thickness of about 7 nm(Figure 2b). The SAED pattern from the particle in Figure 2bis almost identical to the pattern in Figure 2a, which indicatesthat the shell structure was maintained after the irradiationand the separation of the core. No other diffraction spots wereobserved, indicating that the core is amorphous, rather thanpolycrystalline as in the case of zeolite analcime.[6] As thematerial is very sensitive to the beam, HRTEM images of thecrystalline shell were not acquired. The low-magnificationTEM images and the SAED patterns allows us to describe thezeolite A as a monocrystalline cube-like or rectangular boxwith an amorphous core.
The core-shell structure of the as-synthesized zeolite Awas further supported by dark field TEM images of the cross-sections of the cubes prepared by focused ion beam milling(see Supporting Information, Figure SI1) and by dissolutionof the core component in an acidic aqueous solution. Duringthe latter process, 30 mL of 0.35m acetic acid solution wasadded to 1 g of the zeolite A under stirring for 4 h. Most of thecubes lost their inner filling and micrometer-sized hollow
cube-like structures were produced (Supporting Information,Figure SI2a). According to the XRD pattern these hollowstructures were amorphous (Supporting Information, Fig-ure SI2b). As the rate of dissolution of the amorphous core ismuch faster than the crystalline shell, the cube-like orrectangular outer shape was retained, although the crystal-linity of the shell was lost during the acidic treatment. Forcomparison, zeolite A crystals were prepared without chito-san; however, the resulting particles were amorphous (Fig-re SI3).
To investigate the crystallization of zeolite A during thehydrothermal reaction, the products obtained during 1 to 6 hof reaction time were examined. The sample was amorphousafter 1 h, whereas those obtained after 2 and 3 h had zeolite Astructure. After 4 h and 6 h of hydrothermal reactions, amixture of zeolite A and sodalite (Supporting Information,Figure SI4) was obtained. After the 0.35m acetic acid treat-ment, all the crystalline samples obtained from hydrothermalsynthesis of 2–6 h had a rectangular or cube-like morphology(Supporting Information, Figure SI5). These results indicatethat extending hydrothermal reaction time did not lead to thecrystallization of the cores of zeolite A, but resulted in thetransformation of the crystal structure of the shell.
Figure 3 shows the formation of cubes of zeolite A with anamorphous core. Initially, the silica sol is dispersed in anacidified chitosan solution (Figure 3a). After addition of thealkaline solution, the chitosan molecules are deprotonated,resulting in a hydrogel with microsized three-dimensionalpores (Figure 3b). The sodium aluminosilicate gel is producedinside the chitosan hydrogel network by the reaction betweensilica and the alkaline solution. During the hydrothermaltreatment, zeolite nucleation takes place mainly on thesurface of the aluminosilicate aggregates (Figure 3c). Similarto the case of zeolite analcime,[6] some crystalline islandsmight initially form on the surface of the aluminosilicate.These islands then join together, leading to a monocrystallinecube-like shapes by self-alignment of their crystallographicorientations (Figure 3d). Thus, it is an interesting observationthat a very thin-walled crystalline cube-like or rectangular
Figure 2. TEM images of a typical zeolite A particle with a cube-likemorphology and the corresponding SAED patterns obtained from theentire particle. a) Original particle and b) the same particle after beamirradiation for a few minutes.
Figure 3. Representation of the formation of cubes of zeolite A, withan amorphous core (e). The rounded boxes in a)–d) are about4 mm � 4 mm in dimension.
Figure 1. a) XRD pattern and b) SEM image of the as-synthesizedsample.
Communications
8398 www.angewandte.org � 2008 Wiley-VCH Verlag GmbH & Co. KGaA, Weinheim Angew. Chem. Int. Ed. 2008, 47, 8397 –8399
morphology can be developed on the surface of an amor-phous cluster without any specific relationship to the crystalgrowth rate of the crystal planes from a nucleus in theamorphous center (Figure 3d). The driving force behind theformation of such polyhedral shells is the one that minimizesthe surface energy.
It has also been observed that organic polymers havesignificant effects on zeolite crystallization.[7–9,11, 16] The addi-tion of water-soluble polymers in the zeolite gel coulddramatically shorten the prenucleation and nucleation peri-ods and thus accelerate the crystal growth.[16] Mathematicalmodeling[17, 18] and experimental results[3] indicate that zeolitenucleation takes place at the interface between the solutionand the gel by adsorption and rearrangement of the solubleprecursors. In the synthesis described herein, the uncros-slinked chitosan hydrogel networks are highly swollen by thesolution, and the interfaces between the chitosan polymernetworks and zeolite aluminosilicate gel can serve as idealnucleation sites. Such unique interfaces facilitate zeolitecrystallization from the surface of the aliminosilicate gelaggregates. On the other hand, the chitosan hydrogel may alsoplay a role in confining the aluminosilicate aggregates andthus controlling the sizes of the zeolite A cubes. During thehydrothermal treatment, the crystalline shell limits thediffusion of the solution and thus the crystallization of thecores is not able to proceed. It is worth mentioning that fullycrystallized zeolite A cubes were obtained when the gel wasaged overnight at room temperature before the hydrothermaltreatment. The aging process most likely makes the systemmore uniform, in which the chitosan-facilitated zeolitenucleation becomes kinetically less pronounced. In addition,the desired network structure of chitosan hydrogels isessential for the formation of cubes of zeolite A with anamorphous core. Owing to the presence of crosslinkedchitosan hydrogels, the small hydrogel pores greatly confinedzeolite growth leading to zeolite nanocrystals.[11]
In summary, we have shown that cubes of zeolite Aconsisting of a thin crystalline shell and an amorphous corecan be grown within uncrosslinked chitosan hydrogels. It isindicative that the formation of cube-like or rectangular core-shell structures involves particle aggregation and surface-to-core crystallization induced by chitosan networks. This workmay provide a new model system for studying complex zeolitenucleation and growth mechanisms.
Experimental SectionAcetic acid (99%, Sigma–Aldrich; 7 g of 1m) was dissolved indeionized water (14 g) in a polypropylene bottle. Chitosan (averagemolecular weight 120000 gmol�1, ca. 80% deacetylated, Sigma–Aldrich; 1.2 g) was dissolved in the prepared acetic acid solutionunder magnetic stirring for 1 h, followed by addition of the silica sol(HS-30 30 wt %, Sigma–Aldrich; 3.38 g) to the chitosan/acetic acidsolution. The alkaline solution was prepared by mixing NaOH (99%,Merck; (5 g), and NaAlO2 (anhydrous, Sigma–Aldrich; 2.45 g) withdeionized water (20 g). The solution was stirred for 0.5–1 h until itbecame clear. The Na2O/Al2O3/H2O alkaline solution was added tothe chitosan/acetic acid solution without stirring, resulting in a sodium
aluminosilicate gel entrapped inside the chitosan hydrogel. The finalmolar composition of chitosan/SiO2 was 1.18:1. After hydrothermaltreatment at 90 8C for 3 h, the samples were washed with sufficientwater and dried at 80–1008C overnight, followed by calcining thedried sample at 500 8C in oxygen, or treating them with 10%hydrogen peroxide to remove chitosan.[11] In addition, samples werealso synthesized at 90 8C with different hydrothermal reaction times(1, 2, 4, and 6 h).
Scanning electron microscopy (SEM) images were taken with aJSM-6300F microscope (JEOL). Transmission electron microscopy(TEM) images and selected-area electron diffraction (SAED) weretaken with a JEOL JEM-2011 electron microscope operated at200 kV. X-ray diffraction (XRD) patterns were recorded on a PhilipsPW1140/90 diffractometer with Cu Ka radiation at a scan rate of28min�1 and a step size of 0.028.
Received: June 14, 2008Published online: October 2, 2008
.Keywords: chitosan · crystal growth · hydrogels · polymers ·zeolites
[1] C. S. Cundy, P. A. Cox, Microporous Mesoporous Mater. 2005,82, 1 – 78.
[2] S. Mintova, N. H. Olson, V. Valtchev, T. Bein, Science 1999, 283,958 – 960.
[3] V. P. Valtchev, K. N. Bozhilov, J. Am. Chem. Soc. 2005, 127,16171 – 16177.
[4] a) T. M. Davis, T. O. Drews, H. Ramanan, C. He, J. S. Dong, H.Schnablegger, M. A. Katsoulakis, E. Kokkoli, A. V. McCormick,R. L. Penn, M. Tsapatsis, Nat. Mater. 2006, 5, 400 – 408; b) M. A.Snyder, M. Tsapatsis, Angew. Chem. 2007, 119, 7704 – 7717;Angew. Chem. Int. Ed. 2007, 46, 7560 – 7573.
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[6] X. Y. Chen, M. H. Qiao, S. H. Xie, K. N. Fan, W. Z. Zhou, H. Y.He, J. Am. Chem. Soc. 2007, 129, 13305 – 13312.
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[9] L. Han, J. F. Yao, D. Li, J. Ho, X. Y. Zhang, C. H. Kong, Z. M.Zong, X. Y. Wei, and H. T. Wang, J. Mater. Chem. 2008, 18,3337 – 3341.
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[12] A. Bravais, e�tudes Cristallographic, Gauthier-Villars, Paris,1866.
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Available online at www.sciencedirect.com
www.elsevier.com/locate/micromeso
Microporous and Mesoporous Materials 106 (2007) 262–267
Organic-functionalized sodalite nanocrystals and their dispersionin solvents
Dan Li a, Jianfeng Yao a, Huanting Wang a,*, Na Hao a, Dongyuan Zhao a,Kyle R. Ratinac b, Simon P. Ringer b
a Department of Chemical Engineering, Monash University, Clayton, VIC 3800, Australiab Australian Key Center for Microscopy and Microanalysis, The University of Sydney, Sydney, NSW 2006, Australia
Received 30 January 2007; received in revised form 2 March 2007; accepted 5 March 2007Available online 13 March 2007
Abstract
Hydroxy-sodalite nanocrystals with organic functional groups (i.e., @Si–(CH3)(CH2)3NH2, denoted Sod-N, or „Si–CH3, denotedSod-C) were synthesized by the direct transformation of organic-functionalized silicalite nanocrystals. The chemical structure oforganic-functionalized sodalite nanocrystals was confirmed by 29Si MAS NMR spectroscopy. Gas sorption results showed that the soda-lite nanocrystals contained uniform pore channels that were accessible to hydrogen, but inaccessible to nitrogen, as expected. The BETsurface areas are calculated to be 22.8, 19.6 and 19.1 m2/g for plain sodalite nanocrystals (Sod), Sod-N, and Sod-C, respectively; simi-larly, Sod-N and Sod-C exhibited slightly lower hydrogen adsorption than Sod. The dispersion of Sod-N and Sod-C in organic solventswas favored by the presence of organic functional groups. Therefore, the organic-functionalized sodalite nanocrystals prepared in thiswork may be very useful for fabricating zeolite nanostructures and sodalite-polymer nanocomposite membranes.� 2007 Elsevier Inc. All rights reserved.
Keywords: Sodalite; Silicalite; Organic functionalized; Nanocrystals; Dispersion
1. Introduction
Sodalite is a small-pore zeolite whose framework con-sists of a six-membered ring aperture with a pore size of2.8 A. Because of its small pore size and high ion exchangecapacity, sodalite has been considered as a good candidatematerial for a wide range of applications such as opticalmaterials, waste management, hydrogen storage, andhydrogen separation [1]. The active research into efficientstorage and separation of hydrogen has been driven byits potential as an essential component of future energyeconomies. Consequently, we are interested in developinghigh-selectivity, high-flux membranes for the separationand purification of hydrogen gas. Among the various
1387-1811/$ - see front matter � 2007 Elsevier Inc. All rights reserved.
doi:10.1016/j.micromeso.2007.03.006
* Corresponding author. Tel.: +61 3 9905 3449; fax: +61 3 9905 5686.E-mail address: [email protected] (H. Wang).
possible membranes, polymeric ones have been extensivelystudied for hydrogen separation because they are of low-cost and can be easily fabricated into compact hollow fibersand flat sheets with a high separation-area-to-volume ratio[2–4]. Although some polymer membranes exhibit goodhydrogen selectivity and permeability, there is still plentyof room for development of membranes with improvedperformance [2]. Previous studies by a number of groupshave suggested that the incorporation of zeolites into thepolymer matrix can significantly increase gas separationselectivity by enhancing selective gas adsorption and diffu-sion through the membranes [3–6]. Therefore, the additionof sodalite into polymers promises to yield sodalite-poly-mer composite membranes with superior selectivity forhydrogen separation. It has been suggested that template-free sodalite nanocrystals with good interfacial compatibil-ity with the chosen polymer are needed to effectivelyfabricate sodalite-polymer composite membranes [3]. As
D. Li et al. / Microporous and Mesoporous Materials 106 (2007) 262–267 263
part of our project aiming to design such membranes, thefocus of this paper is the synthesis of template-free sodalitenanocrystals with suitably tailored surface properties.
Our newly developed method for synthesizing colloidalhydroxy-sodalite nanocrystals by the transformation of sil-icalite nanocrystals is used in this study [1], since there is noother method for the synthesis of colloidal structure-direct-ing agent free hydroxy-sodalite nanocrystals available. Thehydroxy-sodalite nanocrystals obtained have a sodalitestructure whose framework charges are balanced byhydroxide ions, and they do not enclose template moleculeswithin their pore channels. The strategy of attachingorganic functional groups to zeolites is adopted to achievesuitable surface properties of the hydroxy-sodalite nano-crystals because it is one of the most effective ways formodifying surface properties or adding surface reactivityto zeolite crystals [7–12]. Two kinds of organic groupsincluding methyl and amino moieties are introduced intothe sodalite nanocrystals by adding controlled amountsof 3-aminopropyl(diethoxy) methylsilane and methyltri-methoxysilane during the growth of silicalite nanocrystals.The sodalite nanocrystals are thus expected to be mademore hydrophobic (–CH3) or reactive (–NH2). The prepa-ration and characterization of organic-functionalized soda-lite nanocrystals and their dispersion in solvents aredetailed in this paper.
2. Experimental section
2.1. Synthesis of organic-functionalized silicalite
nanocrystals
Clear synthesis solutions were prepared by dropwiseaddition of 20 g of 1 M tetrapropylammonium hydroxide(TPAOH, Sigma–Aldrich) solution into the mixture of17.8 g of tetraethyl orthosilicate (TEOS, 99%, Sigma–Aldrich) and 1.8 g of 3-aminopropyl(diethoxy) methylsilane(ADMS, 97%, Sigma–Aldrich) or 1.3 g of methyltrimethox-ysilane (MTMS, 98%, Sigma–Aldrich) under vigorous stir-ring, followed by continued stirring at room temperaturefor 3 h. The molar composition of final solution was 1TPAOH:4.32 SiO2:0.48 ADMS (or MTMS): 44 H2O. Crys-tallization was carried out at 80 �C for 12–15 days. Themilky silicalite suspensions obtained were dried at 90–100 �C leading to solid silicalites (denoted Sil-N and Sil-Cfor silicalites prepared with ADMS and MTMS, respec-tively). To observe their morphologies by scanning electronmicroscopy, the samples were prepared by repeated cyclesof washing with deionized water and centrifuging, followedby drying at 90–100 �C overnight.
2.2. Synthesis of organic-functionalized sodalite nanocrystals
An alkaline solution with a molar composition of 6.07Na2O:1 Al2O3:66 H2O was prepared by mixing 20 g ofsodium hydroxide (99%, Merck), 9.2 g of sodium alumi-nates (anhydrous, Sigma–Aldrich), and 60 g of deionized
water at room temperature for 1–2 h. 1 g of the dried silica-lite sample (i.e. Sil-N and Sil-C) was added to 11 g of thealkaline solution during 2–3 min of stirring, and then agedat room temperature for 4 h without further stirring. Thetransformation was carried out at 80 �C for 0–4 h. Thesamples obtained were cooled to room temperature andcollected by repeated cycles of washing with deionizedwater and centrifuging, followed by drying at 90–100 �Covernight. The samples were denoted Sod-N and Sod-C,respectively, when Sil-N and Sil-C were used as silicasource, respectively. For comparison, hydroxy-sodalitenanocrystals (denoted Sod) were also prepared from silica-lite nanocrystals according to our previous paper [1].
2.3. Characterization
Scanning electron microscopy (SEM) images were takenwith a JSM-6300F microscope (JEOL). The particle sizedistributions for Sil-N, Sil-C, Sod-N and Sod-C were deter-mined by manual measurement of 300 nanocrystals each inSEM images with a Photoshop software. X-ray diffraction(XRD) patterns were measured on a Philips PW1140/90 dif-fractometer with Cu Ka radiation (25 mA and 40 kV) at ascan rate of 1�/min with a step size of 0.02�. Thermogravi-metric analysis (TGA, Perkin Elmer, Pyris 1 analyzer) wasperformed in air at a heating rate of 5 �C/min to 600 �C.29Si solid-state nuclear magnetic resonance (NMR) wasconducted on a Bruker DSX300 spectrometer (Germany)under conditions of cross polarization (CP) and magic anglespinning (MAS). 29Si solid-state MAS NMR spectra werecollected at room temperature with a frequency of59.6 MHz, a recycling delay of 600 s, a radiation frequencyintensity of 62.5 kHz, and a reference sampleof Q8M8([(CH3)3SiO]8Si8O12]). Nitrogen and hydrogenadsorption–desorption experiments were performed at77 K with a Micrometritics ASAP 2020MC analyzer anda Micrometritics ASAP 2010MC analyzer, respectively.The samples were degassed at 473 K before analysis. Thesurface areas were determined by the Brunauer–Emmett–Teller (BET) method. Suspended particle size distributionswere quantified by light scattering with a Malvern Master-sizer 2000 analyzer. Different solvents-deionized water, iso-propanol (97%, Sigma–Aldrich), dichloromethane (DCM,Sigma–Aldrich) and dimethylformamide (DMF, Sigma–Aldrich) – were used for sample dispersion. Approximately12–15 ml of suspension was prepared by dispersing 50 mgof sample into 50 ml of solvent under ultrasonication beforeinjection into the Mastersizer for size distribution analysis.
3. Results and discussion
3.1. Transformation of silicalite
The XRD patterns (Fig. 1) show the transformationof organic-functionalized silicalites (Sil-N and Sil-C)under hydrothermal treatment at 80 �C. The organic-functionalized silicalites (Sil-N and Sil-C) became amorphous
10 20 30 40 50 60
3h
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Inte
nsity
(a.u
.)
2θ (degrees)10 20 30 40 50 60
3h2h1h
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nsity
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.)
2θ (degrees)
Fig. 1. XRD patterns of samples prepared with dried organic-functionalized silicalites by hydrothermal treatment at 80 �C for different times. (a)Sil-N toSod-N and (b) Sil-C to Sod-C.
264 D. Li et al. / Microporous and Mesoporous Materials 106 (2007) 262–267
after 1 h hydrothermal treatment. However, in our previ-ous study [1], plain silicalite (without organic groups) waslargely transformed into zeolite A after only 1 h hydrother-mal treatment. This is because the presence of @Si–(CH3)(CH2)3NH2 and „Si–CH3 in silicalite structures(Sil-N and Sil-C, respectively) does not favor aluminosili-cate structure rearrangement during the incorporation ofAl and Na. After 2 h treatment, both samples were a mix-ture of zeolite A and sodalite. The pure organic-functional-ized sodalite, Sod-N, was obtained after 3 h. However, thetransformation of Sil-C into Sod-C took a longer time (4 h)to complete.
Fig. 2 shows the SEM images and particle size distribu-tions of organic-functionalized silicalite nanocrystals (Sil-Nand Sil-C) and organic-functionalized sodalite nanocrystals(Sod-N and Sod-C). All samples exhibit similar morpholo-gies. Sil-N exhibits smaller particle sizes as compared withSil-C, though the synthesis conditions were identical. Thismay be explained by the presence of the –NH2 groups
Fig. 2. SEM images (c and f) and particle size distributions a, b, d, and e of oimages: (a) dried silicalite Sil-N, (b) sodalite Sod-N, (d) dried silicalite Sil-C, andSil-C and Sod-C.
accelerating nucleation in the silicalite synthesis solution,leading to smaller particles on average [13]. This is alsoconsistent with the XRD results above showing that thetransformation of Sil-N into Sod-N took a shorter time.The particle sizes of the organic-functionalized sodalitenanocrystals are larger than those of their precursor silica-lite nanocrystals. This is related to the recrystallization inthe transformation as indicated by XRD. The mean parti-cle sizes are 95 nm, 105 nm, 105 nm and 140 nm for Sil-N,Sod-N, Sil-C and Sod-C, respectively (Fig. 2c and f).
3.2. Evidence of organic functionalization of sodalite
To prove that the organic functional groups have beenincorporated into the sodalite nanoparticles, Sod-N andSod-C samples were characterized by solid-state NMRspectroscopy. The 29Si MAS NMR spectra shown inFig. 3a display a strong resonance peak at around�85 ppm, which arises from Si (4Al) in Sod-N and Sod-C
rganic-functionalized silicalites and organic-functionalized sodalites. SEM(e) sodalite Sod-C. Particle size distributions: (c) Sil-N and Sod-N, and (f)
100 50 0 -50 -100 -150 -200 -250
Si-C (*)
Sod-NSod-C
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SiO
OAl
OSi
O AlO
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SiHO O
AlO
Si
(CH2)3NH2
SiO
OAl
OSi
O AlO
SiO
Si CH3
SiHO O
AlO
Si
Sod-N Sod-C
Fig. 3. (a) 29Si-NMR of organic-functionalized sodalite nanocrystals and (b) the bonding scheme for organic-functionalized sodalite nanocrystals.
0 100 200 300 400 500 600
85
90
95
100
cb
Temperature (°C)
Mas
s (%
)
a
Fig. 4. TGA curves of organic-functionalized sodalite nanocrystals andhydroxy-sodalite nanocrystals. (a) Sod-N, (b) Sod-C and (c) Sod.
D. Li et al. / Microporous and Mesoporous Materials 106 (2007) 262–267 265
[14,15]. The NMR spectra also exhibit a resonance peak ataround �55 ppm, which is ascribed to Si–C bonds [14]. Theresults confirm the existence of organic functional groupsin the Sod-N and Sod-C, and thus the organic groups havebeen incorporated into the sodalites. The integrated area ofthe functionalized silicon peak represents 7.4 mole% and7.2 mole% of the total silicon in Sod-N and Sod-C, respec-tively. The amounts of organic functional groups incorpo-rated into Sod-N and Sod-C are less than those added insilicalite synthesis solutions (10 mole% was added for bothSil-N and Sil-C), but this is reasonable given that a propor-tion of the hydrolyzed ADMS and MTMS would haveremained in the synthetic solutions. The Si–C bondslabeled with asterisk in organic-functionalized sodalitesare illustrated in Fig. 3b.
The organic functionalization of the sodalite nanocrys-tals receives further support from the TGA results, whichare shown in Fig. 4. The mass loss of the pure hydroxy-sodalite was about 11 wt% owing to the loss of the struc-tural water (Fig. 4c) [1]. The mass losses for Sod-N andSod-C were 13.6 wt% and 11.5 wt%, respectively. As com-pared with the pure sodalite nanocrystals, the additionalmass loss of 2.6 wt% for Sod-N and of 0.5 wt% for Sod-C was due to decomposition of organic functional groups(i.e., –(CH3)(CH2)3NH2 or –CH3) at high temperatures[7]. These figures are quite consistent with the expectedmass losses of 3.18 wt% for Sod-N and 0.65 wt% for Sod-C that can be calculated from the proportion of Si–Cbonds measured by 29Si MAS NMR.
3.3. Gas adsorption and pore structures
To further compare the organic-functionalized sodalitenanocrystals (Sod-N and Sod-C) and plain hydroxy-soda-
lite nanocrystals (Sod), nitrogen and hydrogen adsorp-tion–desorption analyses were conducted. The isothermsof Sod-N, Sod-C and Sod are shown in Fig. 5. Theamounts of nitrogen adsorbed in all three samples are verylow at low relative pressures, and substantially increase athigh relative pressures (e.g., P/P0 > 0.8). This is becausewell-grown sodalite pores are inaccessible to nitrogen (N2
kinetic diameter 3.6 A is larger than sodalite pore size2.8 A), and the main nitrogen adsorption arises from theexternal surfaces of nanocrystals. The BET surface areasare calculated to be 22.8, 19.6 and 19.1 m2/g for Sod,Sod-N, and Sod-C, respectively, which is consistent withthe particle size distributions observed by SEM. By con-trast, all samples exhibit much higher H2 adsorption atlow relative pressures as compared with N2 adsorption(Fig. 5a and b), implying that the sodalite channels in thesethree samples are readily accessed by H2 molecules.
0.0 0.2 0.4 0.6 0.8 1.00
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60
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orbe
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-1)
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orbe
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-1)
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Fig. 5. (a) Nitrogen and (b) hydrogen adsorption–desorption isothermals of plain sodalites (Sod) and organic-functionalized sodalites (Sod-N and Sod-C).
266 D. Li et al. / Microporous and Mesoporous Materials 106 (2007) 262–267
Furthermore, the organic-functionalized sodalites (Sod-Nand Sod-C) possess slightly lower H2 adsorption than puresodalite (Sod). At P/P0 = 0.99, the volume of hydrogenabsorbed is around 33.0 cm3/g for Sod, 26.5 cm3/g forSod-N, and 28.0 cm3/g Sod-C. Therefore, the organicgroups do not substantially change the hydrogen adsorp-tion of the sodalite nanocrystals. Clearly, this finding isessential if the functionalized nanoparticles are to be usedsuccessfully in H2 separation membranes.
3.4. Surface modification: dispersion in solvents
To study the effect of organic functionalization on thedispersibility of sodalite nanocrystals, a series of solventsof different polarities was selected: deionized water, isopro-panol, dichlormethane (DCM), and dimethlformamide
100 1000
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100 1000
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25
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ber
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Sod-N Sod-C Sod
a b
c d
Fig. 6. Particle size distributions of organic-functionalized sodalite nanocrysta(b) dimethylformamide (DMF), (c) isopropanol and (d) dichloromethane (DC
(DMF). The solvent polarity of this series, in descendingorder, is water (100) > DMF (42.88) > isopropanol(36.72) > DCM (23.04) [16]. The particle size distributionsof Sod-N, Sod-C, and Sod shown in Fig. 6 are used asan indicator of their relative dispersibility. When deionizedwater is used as a dispersion medium, both Sod-N and Sodhave a similar particle size distribution and their mean par-ticle sizes are approximately 160 nm, which is slightlygreater than that observed by SEM due to the surface sol-vation effect (e.g., surface ionization and adsorption)[17,18]. In contrast, Sod-C exhibits a wider particle size dis-tribution and its mean particle size is approximately270 nm (Fig. 6a). The different dispersibility betweenSod-N/Sod and Sod-C arises from their different surfaceenergy components: Sod-N with –(CH3)(CH2)3NH2 groupsand Sod with –OH groups have similar hydrogen-bonding
100 1000
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ls and plain sodalite nanocrystals in different solvents: (a) deionized water,M).
D. Li et al. / Microporous and Mesoporous Materials 106 (2007) 262–267 267
forces, whereas Sod-C with –CH3 groups is more hydro-phobic. Sod-N, Sod-C, and Sod show similar dispersibilityin DMF (Fig. 6b) because DMF combines a high polarityand high hydrogen-bonding force with hydrophobicgroups. In isopropanol, both Sod-N and Sod-C exhibitslightly better dispersion than Sod (Fig. 6c). Sod-N andSod-C exhibit similar degrees of dispersion in DCM, butthe Sod nanocrystals severely aggregate, leading to a meanparticle size of 880 nm (Fig. 6d). These are because isopro-panol and DCM, with relatively low polarity and poorhydrogen-bonding force, preferentially interact withorganic-functionalized surfaces [19]. These results clearlyshow that the surface properties of sodalite nanocrystalscan be tailored by organic functionalization, which isessential for preparing zeolite-polymer nanocomposites[10,20].
4. Conclusion
We have successfully incorporated organic functionalgroups into hydroxy-sodalite nanocrystals through thedirect transformation of organic-functionalized silicalitenanocrystals. The organic-functionalized sodalite nanocrys-tals showed high crystallinity and well-grown pore struc-tures based on XRD and nitrogen sorption measurements.The micropores of the organic-functionalized sodalite nano-crystals were highly accessible to hydrogen molecules,though there was a slight reduction of hydrogen adsorptioncompared with sodalite nanocrystals without organicgroups. Sodalite nanocrystals with –(CH3)(CH2)3NH2 moi-eties showed good dispersibility in all four solvents (i.e.,water, isopropanol, dichloromethane, and dimethylform-amide) tested whereas sodalite nanocrystals with –CH3
groups were dispersible in isopropanol, dichloromethaneand dimethylformamide, but were agglomerated in water.Without organic functionalization, sodalite nanocrystalsshowed very poor dispersibility in dichloromethane. There-fore, we expect that the organic-functionalized sodalitenanocrystals synthesized in this work will be highly suitedfor fabricating sodalite-polymer nanocomposite mem-branes and other zeolite nanostructures.
Acknowledgments
This work was supported by the Australian ResearchCouncil (Discovery Project No. DP0559724) and MonashUniversity. The facilities and technical assistance from staffat the Electron Microscopy and Microanalysis Facility,Monash University, are gratefully appreciated. H.W.thanks the Australian Research Council for the QEIIFellowship.
References
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Microporous and Mesoporous Materials xxx (2009) xxx–xxx
ARTICLE IN PRESS
Contents lists available at ScienceDirect
Microporous and Mesoporous Materials
journal homepage: www.elsevier .com/locate /micromeso
Synthesis and characterization of sodalite–polyimide nanocomposite membranes
Dan Li a, Huai Yong Zhu b, Kyle R. Ratinac c, Simon P. Ringer c, Huanting Wang a,*
a Department of Chemical Engineering, Monash University, Clayton, VIC 3800, Australiab School of Physical and Chemical Sciences, Queensland University of Technology, Brisbane, QLD 4001, Australiac Australian Key Center for Microscopy and Microanalysis, The University of Sydney, Sydney, NSW 2006, Australia
a r t i c l e i n f o a b s t r a c t
Article history:Received 28 March 2009Received in revised form 10 May 2009Accepted 12 May 2009Available online xxxx
Keywords:SodalitePolyimideNanocomposite membraneHydrogen separation
1387-1811/$ - see front matter � 2009 Elsevier Inc. Adoi:10.1016/j.micromeso.2009.05.014
* Corresponding author. Tel.: +61 3 9905 3449.E-mail address: [email protected]
Please cite this article in press as: D. Li et al., M
Nanocomposite membranes are fabricated from sodalite nanocrystals (Sod-N) dispersed in BTDA-MDApolyimide matrices and then characterized structurally and for gas separation. No voids are found uponinvestigation of the interfacial contact between the inorganic and organic phases, even at a Sod-N loadingof up to 35 wt.%. This is due to the functionalization of the zeolite nanocrystals with amino groups(@SiA(CH3)(CH2)3NH2), which covalently link the particles to the polyimide chains in the matrices. Theaddition of Sod-N increases the hydrogen-gas permeability of the membranes, while nitrogen permeabil-ity decreases. Overall, these nanocomposite membranes display substantial selectivity improvements.The sodalite–polyimide membrane containing 35 wt.% Sod-N has a hydrogen permeability of 8.0 Barrersand a H2/N2 ideal selectivity of 281 at 25 �C whereas the plain polyimide membrane exhibits a hydrogenpermeability of 7.0 Barrers and a H2/N2 ideal selectivity of 198 at the same testing temperature.
� 2009 Elsevier Inc. All rights reserved.
1. Introduction
During the past two decades, mixed matrix membranes(MMMs) have attracted much attention due to their potential forsuperior gas separation performance [1–4]. A variety of inorganicfillers such as zeolite, porous carbon, and nonporous silica havebeen used to fabricate inorganic–polymer composite membranes.Of the many possible separations, hydrogen purification is ofindustrial importance because of its applications in the chemicalindustry, and its use as a fuel in fuel cells. Until now, the many at-tempts to develop zeolite–polymer composite membranes withimproved hydrogen permeability and selectivity have met withlimited success. For instance, S�en et al. developed polycarbonate-matrix membranes filled with highly crystalline zeolite-4A withparticles sizes of 3 lm [5]. At a zeolite loading of 30 wt.%, the com-posite membranes had an improved H2/N2 selectivity of 73.2 com-pared with 56.7 for plain polycarbonate. However, they also founda decrease in hydrogen permeability, which they attributed to in-creased rigidity of the polymer chains in the presence of the zeoliteparticles [6,7], the partial blockage of the zeolite pore by the poly-mer chains [7] and/or the extended diffusion pathways of thehydrogen molecules through the membrane [8,9]. A similar trendwas reported by Li et al. [10] and Huang et al. [7]. Li et al.demonstrated that membranes of polyethersulfone and zeolite5A (1–5 lm) exhibited about 25% higher H2/N2 selectivity than aplain polyethersulfone membrane, but had a decrease in gas
ll rights reserved.
u (H. Wang).
icropor. Mesopor. Mater. (2009
permeability of at least 25% [7,9]. Huang et al. prepared their com-posite membranes by incorporating 20 wt.% of micrometer-sized(1–5 lm) or nanometer-sized (50–140 nm) zeolite A in polyether-sulfone (PES) [7]; the hydrogen permeability of the PES membranedropped from 8.96 Barrers to 8.3 Barrers when filled with thenano-zeolite and down to 4.94 Barrers for the micro-zeolite. Inter-estingly, the gas permselectivity enhancement was much morepronounced when zeolite-4A nanocrystals were incorporated in aPES membrane. Indeed, nano-sized zeolites are required for fabri-cating composite membranes because the polymeric membranesare usually shaped into asymmetric hollow fibers or flat sheetswith a thin selective layer (e.g., <1 lm) for practical applications[10]. Up to 40.2 wt.% silicalite-1 (MFI) nanocrystals (80 nm) werecombined with Telfon AF 1600 polymers by Golemme et al. [11].The composite membranes had a hydrogen permeability of 3580Barrers, a 15-fold increase relative to the pure polymer mem-branes. However, the H2/N2 selectivity of the composite mem-branes, at just 4.6, was 50% less than the pure Telfon film. Thisresult was probably due to interfacial voids between the zeoliteand polymer, which were formed because of low adhesion be-tween the polymer matrix and the zeolite crystals [5,6,12–14].
Several approaches have been proposed to fabricate the mixed-matrix membranes that are free of voids and have enhanced selec-tivity [15–17]. One of the most effective ways is the surface mod-ification of the zeolite particles with silane-coupling agents[3,16,18]. For instance, Duval et al. promoted the adhesion betweenzeolite particles and polymer matrices by modifying the zeolitesurfaces with silane-coupling agents (e.g., y-aminopropyltriethoxysilane, N-p-(aminoethy1)-y-aminopropyltrimethoxy silane and
), doi:10.1016/j.micromeso.2009.05.014
2 D. Li et al. / Microporous and Mesoporous Materials xxx (2009) xxx–xxx
ARTICLE IN PRESS
styryl amino functional silane). Unfortunately, the measured per-meabilities were slightly lower compared with polymer matrixwhile the ideal selectivities were largely unchanged [16]. Pecharet al . developed mixed-matrix membranes from polyimide andzeolite L or ZSM-2 zeolites, which were functionalized by APTES(aminopropyl-triethoxysilane) coupling agents. The gas selectivityof the composite membranes was enhanced, but the gas perme-ability was unexpectedly lowered relative to the pure polyimidemembranes [18,19]. Similarly, Li et al. found that the increase inselectivity of membranes made with zeolite A, which had beenmodified with APDEMS (3-aminopropyl)-diethoxymethyl silane),was offset by a decrease in permeability [3,9].
Nonporous fillers such as silica nanoparticles were also incorpo-rated into polymer to yield inorganic–organic polymer compositemembranes. Merkel et al. found that silica-poly(4-methyl-2-pen-tyne) nanocomposite membranes exhibited significantly enhancedmembrane permeability and selectivity for large organic moleculesover small permanent gases. This was because physical dispersionof nanoporous nanoparticles yielded polymer-particle interfaces,disrupted polymer chain packing and thus affected moleculartransport [20].
The objective of our work is to study the feasibility of fabricatingnanocomposite membranes with improved separation propertiesby incorporating organic-functionalized sodalite nanocrystals intopolymer. Sodalite is a type of small-pore zeolite, which has asix-membered-ring aperture with a 2.8 Å pore size. Sodalite wasreported to exhibit good hydrogen adsorption property at high tem-peratures (e.g., >573 K) [21]. Our recent study shows that sodalitenanocrystals exhibit attractive hydrogen adsorption–desorptionbehavior at very low temperatures (e.g., 77 K) [22]. This interestingtemperature-dependent hydrogen sorption property is related tothe change in the size of sodalite cage at different temperatures.It would be of fundamental interest to investigate how sodalitenanocrystals affect the microstructure and separation propertiesof polymer membranes. In this study, we used modified sodalitenanocrystals, functionalized with = SiA(CH3)(CH2)3NH2 groups aswe previously reported [22,23], as the inorganic phase in compositemembranes. We chose polyimide as the continuous polymer matrixfor this study; polyimides have attracted considerable interest forhydrogen separation, because of their good gas transport proper-ties, their thermal and chemical stability, and their mechanicalproperties [24–28]. Previous research has reported excellent selec-tivity, varying from 64.8 to 365, for separating hydrogen from nitro-gen by using polyimide prepared from different kinds of monomers[26,27]. Fluorinated polyimides usually possess higher H2 perme-ability and lower selectivity over other gases as compared withnon-fluorinated polyimides. For instance, 6FDA-DDBT polyimideexhibits a H2 permeability of 156 Barrers and a H2/CH4 selectivityof 78.8 [29] whereas BPDA-ODA polyimide has a H2 permeabilityof 1.33 Barrers and a H2/N2 selectivity of 365 [27]. Here, we havechosen a polyimide with a moderate H2/N2 selectivity for thefabrication of sodalite-polyimide membranes. Two monomers(benzophenone-3,30,4,40-tetracarboxylic dianhydride and 4,40-diaminodiphenylmethane) are thus used to synthesize polyimidethat is bonded directly to nanoparticles of organic-functionalizedsodalite, resulting in membranes free from interfacial defects. Thefabrication, characterization and separation performance of thesecomposite membranes are detailed in this paper.
2. Experimental
2.1. Sodalite synthesis and membrane fabrication
The amino-functionalized sodalite nanocrystals (denoted Sod-N) with a mean size of 105 nm were synthesized by transforming
Please cite this article in press as: D. Li et al., Micropor. Mesopor. Mater. (2009
silicalite nanocrystals according to our reported method [22].Briefly, a clear synthesis solution was prepared by dropwise addi-tion of 20 g of 1 M tetrapropylammonium hydroxide (TPAOH, Sig-ma–Aldrich) solution into a mixture of 17.8 g of tetraethylorthosilicate (TEOS, 99%, Sigma–Aldrich) and 1.8 g of 3-aminopro-pyl(diethoxy) methylsilane (ADMS, 97%, Sigma–Aldrich) withvigorous stirring, followed by continued stirring at room tempera-ture for 3 h and then crystallization at 80 �C for 12–15 days. Themilky silicalite suspensions so obtained were dried at 90–100 �Cto obtain solid silicalites. An alkaline solution was prepared bymixing 20 g of sodium hydroxide (99%, Merck), 9.2 g of sodium alu-minates (anhydrous, Sigma–Aldrich), and 60 g of deionized waterat room temperature for 1–2 h. We added 1 g of the dried silicalitesample (denoted Sil-N) to 11 g of the alkaline solution during 2–3 min of stirring, and then allowed it to age at room temperaturefor 4 h without further stirring. The transformation was carriedout at 80 �C for 4 h. The resulting amino-functionalized sodalitenanocrystals were cooled to room temperature and collected by re-peated cycles of washing with deionized water and centrifuging,followed by drying overnight at 90–100 �C.
Monomers benzophenone-3,30,4,40-tetracarboxylic dianhydride(BTDA; 96%, Sigma–Aldrich) and 4,40-diaminodiphenylmethane(MDA; 97%, Sigma–Aldrich) were dried at �150 �C for at least12 h under vacuum. Dimethylformamide (DMF) (GR, Merck) wasdried and stored with 4-ÅA
0
molecular sieves prior to use. To fabri-cate each composite membrane, a given quantity of Sod-N nano-crystals was dispersed in 10 g of DMF under ultrasonication atroom temperature for 30 min. Then 1.5 g of BTDA and 1.92 g ofMDA were dissolved in the Sod-N suspension. The resulting mix-ture was stirred for 5 h in an ice-water bath at approximately0 �C under N2 gas to obtain a Sod-N/PAA (polyamic acid) precursor,which was a cloudy yellow, viscous solution. The Sod-N/PAA solu-tion was cast directly onto a glass plate and placed into a vacuumoven and heat treated for 2 h each at 50 �C, at 100 �C and at 150 �C,before it was held at 200 �C overnight. The resulting sodalite–poly-imide nanocomposite membrane (denoted Sod-N/PI) was slowlycooled to room temperature. All of the yellow Sod-N/PI films wereimmersed in hot water at 90 �C for 1 h to allow removal from theglass plates, after which they were dried under vacuum at 150 �Covernight before analysis. In this paper, the sodalite–polyimidenanocomposite membranes were made with sodalite loadings of15, 25 and 35 wt.% (based on the mass of polyimide) and theseare denoted PI-15, PI-25, and PI-35, respectively. For comparison,pure polyimide membranes were prepared by applying the aboveprocedures without any Sod-N additions and these are referredto as PI-0.
2.2. Characterization
Scanning electron microscopy (SEM) images of cross sections ofmembranes were taken with a JSM-6300F microscope (JEOL). X-raydiffraction (XRD) patterns were measured on a Philips PW1140/90diffractometer with Cu Ka radiation (25 mA and 40 kV) at a scanrate of 1�/min with a step size of 0.01�. Fourier-transform infraredspectra (FT-IR) were recorded for the samples embedded in KBrpellets with a GX Spectrometer (Perkin–Elmer). Thermogravimet-ric analysis (TGA, Perkin–Elmer, Pyris 1 analyzer) was performedat a heating rate of 5 �C/min to 700 �C in oxygen with a flow rateof 15 cm3 min�1. Hydrogen adsorption–desorption experimentswere performed at 77 K and room temperature, and a pressure ofup to 900 mm Hg with a Micrometritics ASAP 2010MC analyzer.The samples were degassed at 473 K before analysis. To test gasseparation properties, the composite membrane or pure polyimidemembrane samples were firstly attached to a porous stainless-steel stand (pore size � 200 nm), which was then fixed in asample holder by using Torr Seal epoxy resin (Varian). Before
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Fig. 2. XRD patterns of samples Sod-N, PI-0, PI-15, PI-25, and PI-35. The peakslabeled with asterisks arise from Sod-N.
Fig. 3. IR spectra of samples Sod-N, PI-0 and PI-35.
D. Li et al. / Microporous and Mesoporous Materials xxx (2009) xxx–xxx 3
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measurements, the samples were evacuated and dried in a vacuumoven at 200 �C overnight to remove any residual solvent and ad-sorbed water. The gas permeation tests were performed at 25, 60and 100 �C on pure H2 and pure N2. The pressure rise of the perme-ate stream was measured with a Series 901 Transducer (MKS),which was connected to computer. Membrane permeability, Pi,was defined as [10,30],
Pi ¼dNi
DPiA
where d is the membrane thickness (cm), Ni the permeation rate ofcomponent i (cm3 s�1), DPi the transmembrane pressure differenceof i (cm Hg), and A the membrane area (cm2). 1 Barrer = 10�10
cm3(STP) cm cm�2 s�1 cm Hg�1. The ideal selectivity, aij, betweentwo gases, i and j, was defined as, [31,32]
aij ¼Pi
Pj
The apparent activation energy Ep was analyzed according tothe Arrhenius equation [31,33–35],
P ¼ P0 exp�Ep
RT
� �
where P is the permeability, P0 the pre-exponential factor, R theideal gas constant (8.3143 J mol�1 K�1) and T is the temperaturein Kelvin (K).
3. Results and discussion
3.1. Membrane characterization
Fig. 1 shows photographs of the series of polyimide compositemembranes with a thickness of 50 lm, which were all intact andhomogeneous, laid over the word ‘‘Monash”. Pure polyimides areclear, flexible and have good tear strength. All of the compositemembranes have a yellow appearance, but their transparency de-creases with increasing content of Sod-N nanocrystals (Fig. 1), asis evident from the gradual obscuration of the word from PI-0 toPI-35. Fig. 2 shows the XRD patterns of pure Sod-N and for PI-0,PI-15, PI-25 and PI-35. The Sod-N nanocrystals exhibit good crys-tallinity, giving sharp peaks in XRD pattern, which have been in-dexed in Fig. 2. In contrast, the pure polyimide membrane (PI-0)appears to be amorphous, as expected. With increasing contentsof Sod-N nanocrystals in the polyimide membranes, the peaks inFig. 2 increase in intensity from PI-15 to PI-35.
Fig. 1. Photos of PI-0, PI-15, PI-25 and PI-35 showing the change in transparencywith increasing Sod-N content.
Please cite this article in press as: D. Li et al., Micropor. Mesopor. Mater. (2009
Fig. 3 shows the IR spectra of Sod-N, PI-0 and PI-35. For the lasttwo samples, absorption bands, which correspond to the polyimidestructure, are observed at 1780 cm�1 (C@O asymmetric stretch-ing), 1720 cm�1 (C@O symmetric stretching), 1380 cm�1 (CANstretching), and 720 cm�1 (imide ring deformation); these indicatethe successful chemical imidization of the membranes [25,36–38].For the pure Sod-N sample, the broad band at approximately990 cm�1 is assigned to the asymmetric stretch (T-O-T, T = Si, Al),and the adsorption at 661 cm�1 is ascribed to the symmetricstretch (T-O-T) [22,39]. The presence of Sod-N in sample PI-35causes the peaks at around 1000 cm�1 to broaden in comparisonwith pure PI-0 film. Furthermore, there is a new small peak ap-pears for PI-35 at 661 cm�1, which is due to asymmetric stretchT-O-T (T = Si, Al) arising from added Sod-N.
Fig. 4 shows the SEM images of cross sections of PI-0, PI-15, PI-25 and PI-35. These micrographs confirm that Sod-N nanocrystalsare well dispersed throughout the polyimide matrix at all loadingsof Sod-N. No voids are apparent between the nanocrystals andpolyimide, even at 35-wt.% Sod-N where some large-scale surfaceroughness is evident, which suggests good bonding and compati-bility between the zeolite and polymer. Other studies also havefound that improving the interaction between zeolites and poly-mer tends to inhibit formation of interfacial voids [3,18,19,40].
The thermogravimetric (TG) curves of pure polyimide and thecomposite membranes with different loadings of Sod-N are shownin Fig. 5; Table 1 summarizes the corresponding thermogravimet-ric (TG) and differential thermogravimetric (DTG) results. Underflowing oxygen, the pure polyimide membrane, PI-0, lost 1.6% ofits mass in the temperature range from 30–400 �C. This is due tothe loss of residual organic solvent (DMF has a boiling point of153 �C) and/or adsorbed water. In the temperature range from400 to 700 �C, the remaining 98.4% of mass was lost, leaving no
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Fig. 5. TGA curves of samples PI-0, PI-15, PI-25, and PI-35.
Table 1DTG and TGA results of PI-0, PI-15, PI-25 and PI-35.
Sample DTG TG
Td (�C) Mass loss (%) Mass residueafter TGA (%)
Sod-N content (%)
30–400 �C 400–700 �C Experimental Theoretical
PI-0 573 1.6 98.4 0 0 0PI-15 580 3.0 84.5 12.5 14.8 15.0PI-25 595 2.7 78.0 19.3 24.7 25.0PI-35 600 3.4 71.4 25.2 34.8 35.0
Fig. 6. H2 adsorption–desorption isotherms of amino-functionalized sodalitenanocrystals at 77 and 298 K.
Fig. 4. SEM images of cross sections of PI-0, PI-15, PI-25 and PI-35.
4 D. Li et al. / Microporous and Mesoporous Materials xxx (2009) xxx–xxx
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residue after the TGA run, which we ascribed to the completedecomposition and combustion of the polyimide at high tempera-ture [25,36]. The DTG peak (Td) for the corresponding mass loss liesat 573 �C.
The mass losses varied for the composite membranes duringheating between 30 and 400 �C – 3.0%, 2.7% and 3.4% for PI-15,PI-25 and PI-35, respectively – but all the composites lost moremass than PI-0. This might be due to increased adsorption of waterand/or DMF caused by the hydrophilic Sod-N particles and/or bythe presence of inorganic–organic cross-linked networks after
Please cite this article in press as: D. Li et al., Micropor. Mesopor. Mater. (2009
polymerization [36]. However, most of the mass loss occurs inthe temperature range from 400 to 700 �C and is 84.5% for PI-15,78.0% for PI-25 and 71.4% for PI-35. Interestingly, the Td valuesfor the composite materials are all higher than that of pure polyim-ide, and increase with Sod-N content: 580 �C for PI-15, 595 �C forPI-25 and 600 �C for PI-35. Some previous research attributed thiskind of trend to the interaction between the amino moieties frominorganic nanoparticles (Sod-N) and the polymer matrix, whichcan reduce the movement (increase the rigidity) of the polymerchains and, thus, increase the decomposition temperature of com-posite membranes [25,37,41].
The residual masses after TG analysis are 12.5%, 19.3% and25.2% for PI-15, PI-25 and PI-35, which would correspond to plainsodalite nanocrystals, given that organic functional groups (i.e.,A(CH3)(CH2)3NH2) would have been completely decomposed andremoved by the high temperatures [42]. The use of 29Si-NMR inour previous work showed that 3.18 wt.% of Sod-N comprises or-ganic functional groups [22]. This allows recalculation of the actualSod-N loading of PI-15, PI-25 and PI-35 as 14.8%, 24.7% and 34.8%,respectively, based on the mass of polyimide, which are close tothe theoretical values.
3.2. H2 sorption of sodalite nanocrystals and gas permeation ofmembranes
H2 sorption isotherms of amino-functionalized sodalite nano-crystals are shown in Fig. 6. It is clear that the temperature hassubstantial influence on H2 adsorption capacity of sodalite nano-crystals. At 77 K, H2 adsorptive volume significantly increases withincreasing the adsorption pressure, and it reaches a maximum vol-ume of 26.9 cm3/g. However, at room temperature (298 K), amino-functionalized sodalite nanocrystals exhibit almost no H2 adsorp-tion as P/Po is raised to 1 (Po = 900 mm Hg). This is due to sodalitecage contraction when the sorption temperature increases from77 K to 298 K. XRD analysis confirms that the crystallinity in ami-no-functionalized sodalite nanocrystals remains unchanged afterH2 sorption analysis. According to Ref. [21], sodalite cage expandsand starts to uptake hydrogen at 573 K or above. These indicatethat amino-functionalized sodalite nanocrystals may function asnonporous nanoparticles in nanocomposite membranes in ourgas permeation temperatures.
Table 2 summarizes the permeability values of two pure gases(H2 and N2) and the ideal selectivity aðH2=N2Þ for pure polyimidefilms and composite membranes at three different temperatures(25, 60 and 100 �C). Our permeability and ideal selectivity datafor pure polyimide membranes fabricated from BTDA and MDA iscomparable to similar polyimide membranes in the literature
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Table 2Gas permeation results of the PI-0, PI-15, PI-25, and PI-35 membranes.
Sample Permeability (Barrers) Ideal Selectivity aðH2=N2Þ
25 �C 60 �C 100 �C 25 �C 60 �C 100 �C
H2 N2 H2 N2 H2 N2
PI-0 7.0 0.036 12.0 0.096 13.9 0.13 198 124 110PI-15 7.4 0.033 10.8 0.079 12.7 0.11 223 137 113PI-25 8.1 0.034 9.9 0.056 11.3 0.073 238 176 154PI-35 8.0 0.029 9.9 0.043 13.1 0.062 281 230 210
Fig. 8. Ideal selectivity aðH2=N2 Þ of PI-0, PI-15, PI-25, and PI-35 at differenttemperatures (25, 60 and 100 �C).
D. Li et al. / Microporous and Mesoporous Materials xxx (2009) xxx–xxx 5
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[29]. The addition of Sod-N causes the composite membranes to re-duce the permeation of N2, leading to a substantial improvement inH2/N2 selectivity. This should be attributed to the interfacial effectsand disrupted polyimide chain packing caused by the covalentbonding between Sod-N and polyimide. The structure of sodalite-polyimide interface is illustrated in Fig. 7a. Sodalite nanocrystalsare composed of a crystalline sodalite core and a thin amorphousaluminosilicate shell with amino-groups (@SiA(CH3)(CH2)3NH2).The thickness of the amorphous aluminosilicate shell is roughlyestimated to be around 2 nm assuming that all amino-groups arecontained in the shell [22]. The high-quality bonding betweenthe sodalite nanocrystals and the polymer matrix is realized byforming covalent linkers via the imidization reaction of the ami-no-groups with the polyimide monomers (Fig. 7b). The additionof Sod-N also affects the chain length of polyimide molecules sur-rounding Sod-N nanocrystals because the polyimide chains react-ing with amino-groups are terminated. This would increase therigidity of the polymer chains in the interfaces and polyimide ma-trix [6,7]. These unique structures allow H2 to diffuse throughwhile reducing the passage of N2 molecules. This explains thatthe H2 permeability of all composite membranes at 25 �C is slightlyhigher than that of the pure polyimide membrane. On the otherhand, these data provide strong evidence that there are no voidspresent at the polyimide and sodalite interface in any of the com-posite membranes, because such voids would have resulted in alarge increase in permeability of H2 or even N2 [18].
When the testing temperature is elevated, there is a subsequentincrease in the permeability of H2 or N2 for the pure-polyimide andthe composite membranes. There was a more significant increasein permeability for the pure polymer with temperature than wasfound for the composite membranes, especially for N2 gas. PI-0has a N2 permeability of 0.036 Barrer at 25 �C, compared with0.13 Barrer at 100 �C, a 3.6-fold increase. However, PI-35 showed
Fig. 7. Schematic representation of sodalite–polyimide interfacial structure (a) andcovalent linker between Sod-N and polyamide (b).
Please cite this article in press as: D. Li et al., Micropor. Mesopor. Mater. (2009
an increase of only 1.1 times for PN2 between room temperatureand 100 �C. In addition, at 60 and 100 �C the permeabilities of H2
and N2 for the composite membranes are lower than those forthe PI membrane (Table 2). As the temperature increases, the per-meabilities of both H2 and N2 increase because of the increase ofthe diffusivity and the decrease of the solubility in polyimides[43]. We attribute this result to the increase in polymer chain rigid-ity in the composite membranes with increasing Sod-N loading,and the increase in the permeabilities of both H2 and N2 for thePI membrane is greater than those for the composite membranes.
The H2/N2 selectivity for PI-0, PI-15, PI-25 as a function of tem-perature, which are included in Table 2, are shown in Fig. 8. At25 �C, the nanocomposite membranes demonstrate perselectivitiesof 223, 238, and 281 for PI-15, PI-25, and PI-35, respectively. Thesevalues represent 13%, 20%, and 46% greater ideal selectivity,respectively, than PI-0.
It is also plain from Fig. 8 that elevating the temperature lowersthe ideal selectivities of all the membranes, but that increasing thesodalite content considerably retards the falling-off of gas selectiv-ity from 25 to 100 �C. For instance, PI-0’s selectivity drops from 198at 25 �C to 110 at 100 �C, which is a fall of 44%. In contrast, PI-35sees a decrease in aðH2=N2Þ of only 25%. As Sod-N loading increases,the number of Sod-N terminated increases substantially, affectingthe chain configuration beyond the interfaces; in other words,the interfacial area may be extended. Therefore, the increased inor-ganic content in composite materials restricts the thermal motionof the polymer segments, and thus reduces the decrease in the gasselectivity [31].
Fig. 9 shows the apparent activation energy, Ep, of PI-0, PI-15, PI-25 and PI-35 for the pure H2 and pure N2. It is apparent that allsamples have higher values Ep for N2 than H2, confirming that N2
molecules need more energy to penetrate the membranes thanH2 molecules. Compared with composite membranes, pure poly-imide polymer (PI-0) has the highest activation energies – 8.5 kJ/mol for H2 and 15.9 kJ/mol for N2. In the composite membranes,the presence of Sod-N lowers Ep below that of the pure polymermembranes. For example, PI-15 and PI-25 have Ep values of 6.7and 4.1 kJ/mol, respectively, for H2 and 14.9 and 9.5 kJ/mol, respec-tively, for N2. Interestingly, PI-35 shows an increase in activationenergy relative to PI-25 for H2, but not for N2. Similarly, the H2
permeability for PI-35 increases largely from 9.9 Barrers to 13.1Barrers as the temperature is increased from 25 to 100 �C. In thecomposite membranes, gas diffusion requires relatively smallsegmental motions of polymer matrix in the packing-disruptedpolyimide chains and sodalite–polyimide interfaces, because theypossess relatively more unoccupied free space. When Sod-Nloading is increased to a certain point (e.g., 35%), the overlap of
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Fig. 9. Apparent activation energy (Ep) for PI-0, PI-15, PI-25 and PI-35.
6 D. Li et al. / Microporous and Mesoporous Materials xxx (2009) xxx–xxx
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interfacial layers becomes significant [44]. Such overlapped inter-faces favor H2 diffusion, and are more temperature-dependent inthe permeation of small hydrogen molecules. It is clear that theseparation performance of polyimide membrane has been signifi-cantly enhanced. The strategy of forming nanocomposite mem-branes demonstrated in this work could be applied to fabricatepractical H2 separation membranes by incorporating functionalsodalite nanocrystals into a more permeable polyimide skin layer.It would be interesting to study water transport property of soda-lite-polyimide nanocomposite thin membranes for potential appli-cations, such as in water/organic solvent separation and waterpurification, given that sodalite membranes have been reportedto exhibit good water permeation property [45].
4. Conclusions
We have used organic-functionalized sodalite nanocrystals(Sod-N) and polyimide to fabricate nanocomposite membranes.Characterization by SEM showed that Sod-N can be well distrib-uted with polyimide phase, even at a loading of 35 wt.%, as is con-firmed by the FTIR spectroscopy and XRD results. From TG and DTGanalysis, the DTG peaks for corresponding major mass loss increasewith the increasing Sod-N content of the composite, which isattributed to restricted movement of the main chains arising fromthe interaction between the amino moieties from inorganic nano-particles (Sod-N) and polymer matrix. The gas permeation experi-ments were performed with two pure gases, H2 and N2, and theresults revealed that H2 permeability was improved, while N2 per-meability decreased. In particular, the PI-35 composite membraneshad the highest ideal selectivity (aðH2=N2Þ = 281) and a good perme-ability (8.0 Barrers) at room temperature.
Acknowledgments
This work was supported by the Australian Research Council(ARC) and the CSIRO Flagships – Advanced Membrane Technology
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for Water Treatment Cluster. H.W. thanks the ARC for the QEII Fel-lowship. D.L. gratefully acknowledges Monash University for thepostgraduate scholarships.
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