Recent Advances in Gas Separationby Microporous Ceramic Membranes N.K. Kanellopoulos(Editor) e 2000 Elsevier Science B.V. All rights reserved.
335
Microporous Silica Membranes
Nieck Benes, Arian Nijmeijer and Henk Verweij
Laboratory of Inorganic Materials Science, Department of Chemical Technology, University of Twente, PO Box 217, 7500AE Enschede, the Netherlands
Introduction
Microporous silica membranes have a high potential for gas separation and pervaporation at high temperatures in chemically aggressive environments. Well-prepared silica membranes show high fluxes for small gas molecules such as H2, CO2 and O2 and considerable selectivities for these gases with respect to larger gas molecules such as SF6 and hydrocarbons [ 1,2]. This offers perspectives on applications such as natural gas purification, molecular air filtration, selective CO2 removal and industrial 1-12 purification. A specific application for these membranes is the use in high temperature membrane reactors in which silica membranes can be of particular use to remove Ha selectively with high fluxes to acieve conversion enhancement in thermodynamically limited reactions. Examples of such reactions can be found in steam reforming, the water-gas shift process, dehydrogenation of hydrocarbons and coal gasification [3,4]. Two different types of molecular sieving silica membranes can be distinguished: 9
Chemical Vapour Infiltrated (CVI) membranes, which are commercially available*.
9
Sol-gel silica membranes, which are not commercially available yet.
CVI membranes are produced by reacting a gaseous silica precursor such as Tetra-Ethyl-Ortho-Silicate (TEOS) with an oxidising agent inside the pores of a macro- or mesoporous support [5,6]. These membranes normally have very high permselectivities towards hydrogen, values as high as 3000 have been measured for H2/N2. A large drawback of such membranes is, however, their relatively low permeance (2-4x 10s mol/m2sPa at 200~
due to the presence of the nearly dense silica plugs inside the
pores of the supporting system. By changing reaction conditions it is possible to obtain a higher hydrogen permeance, but at the expense of selectivity. CVI membranes recently developed in our group have a H2/N2 permselectivity of 43, but with a H2 permeance of 1.7x 10-7 mol/m2sPa at 200~ [6]. An possible advantage of CVI membranes, however, is that the vulnerable separative silica layer is located inside the pores of the supporting system, where it is to some extent protected to aggressive environments. Sol-gel coated silica membranes have a separative layer that is coated on top of a supporting Media and Process TechnologyInc. (MPT), Pittsburgh, PA, USA.
336
system and show fluxes that are again a factor of 10 higher than the above-mentioned CVI membranes and hence in the range of 1-2x 10-6 mol/m2sPa. With such high permeances the supporting system may very well become the limiting factor instead of the active membrane layer. Compared to polymeric membranes, inorganic microporous membranes with molecular sieve-like properties have a good chemical, mechanical and thermal stability [7]. Nevertheless, the stability of silica membranes towards water and water vapour at elevated temperatures and how they affect the membrane performance is not yet elucidated. The issue of thermal and chemical resistance is not only relevant during applications but also in membrane cleaning procedures which often specify strong acids and bases. A general rule is that more acidic metal oxides or ceramics show greater resistance towards acids but are more prone to attack by bases and vice versa [8]. For example, alumina or zirconia membranes generally are more stable than silica when exposed to alkaline solutions. On the other hand, silica membranes have better acidic resistance than most other metal oxide membranes. Recently some interesting self-organising mesoporous silica structures which can be used for membrane purposes have been realised by the group of Brinker [9,10]. By templating methods, using small molecules, they were also able to prepare microporous silica membranes with a controlled pore-size [11,12]. For those ceramic membranes, which contain two oxides, the chemical resistance towards acids and bases often lies between those of the constituents. Even within a given metal oxide system, the chemical resistance may vary with the particular phase. For example, ct-alumina is very stable towards strong acids and bases, the "/-alumina phase however has been known to be subject to some attack at high and low pH. The objective of this chapter is to describe recent developents in synthesis, thermochemical stability and transport properties of supported silica membranes. We did not attempt to provide a complete literature survey on the subject but instead, focused on recent results obtained in the 'Inorganic Materials Science' group of the University of Twente.
Synthesis
The microporous silica membranes prepared in our group consist of three layers. First a support is prepared from or-alumina powder with a fiat or tubular shape. Flat supports are prepared by either die pressing [2] or colloidal filtration [13] and the tubular supports are prepared by the centrifugal casting technique [14]. The use of colloidal processing techniques such as filtration and centrifugal casting come more and more into scope because they result in an extremely homogeneous and hence strong porous structure and a high surface quality. The latter is of importance to be able to apply very thin
337 defect-free membrane layers. On top of the supports a 7-alumina intermediate layer is coated under clean-room conditions. Finally on top of this intermediate layer the final molecular sieving silica top layer is coated. In this section all synthesis steps to obtain a supported microporous silica membrane will be treated, starting with the silica top layer. After that some highlights on 7-alumina intermediate layers, such as pore-sizes and stability, will be given and finally the preparation of flat and tubular supports will be described in detail.
Silica top layer Two types of silica top layers are of interest, the conventional hydrophilic layers and the newly developed hydrophobic silica top layers [ 15]. These layers are prepared by dipping supported y-A1203 membranes in polymeric silica dip solution, followed by drying and calcining. A standard silica sol is prepared by acid-catalysed hydrolysis and condensation of tetra-ethyl-ortho-silicate (TEOS)* in ethanol. A mixture of acid and water is carefully added to a mixture of TEOS and ethanol under vigorous stirring. During the addition of the acid/water mixture the TEOS/ethanol mixture is placed in an ice-bath to avoid premature (partial) hydrolysis. After the addition is complete the reaction mixture is refluxed for 3 hours at 60~ in a water bath under continuous stirring. The reaction mixture had a final molar TEOS/ethanol/water/acid ratio of 1/3.8/6.4/0.085 in agreement with the "standard" recipe of silica sol preparation, as defined in [ 16]. The reacted mixture was cooled and diluted 19 times with ethanol to obtain the final dip solution. After dipping the membranes were calcined at 400~ for 3 hours in air with a heating and cooling rate of 0.5~
The whole process of dipping and calcining can be repeated once again to repair any de-
fects in the first silica membrane layer. Recent results showed, however, that this second coating step is not absolutely necessary anymore if one works under class 100 cleanroom conditions. The membranes are henceforth referred to as "Si(400) membranes". Another type of membranes was prepared by the same procedure as described above but with the only difference that the calcination temperature was 600~
These membranes will be referred to as "Si(600) membranes". More recently the first de-
fect-free silica membranes "Si(800)" were prepared with a firing temperature of 800~ Hydrophobic silica layers [ 1,15] To make the silica membrane material more hydrophobic, methyl-tri-ethoxy-silane (MTES) "r is incorporated at a certain stage of sol preparation. The hydrolysis/condensation rate at room temperature of MTES is --7 times higher than that of TEOS [ 17]. This implies that the reaction time of MTES should
P.a. grade, Aldrich ChemicalCompanyInc., Milwaukee(WI), USA. *P.a. grade, Aldrich ChemicalCompanyInc., Milwaukee(WI), USA.
338 be ---7 times shorter to obtain silica polymers with dimensions similar to those obtained with hydrolysis and condensation of TEOS. This simple consideration led us to the idea to start with a "standard" silica sol solution preparation and add MTES after 6/7 of the normal total reaction time at least. If MTES was added earlier, for instance after 2/3 of the total reaction time, more "bulky" polymers were formed, visible through light scattering in the sol solution. This implied that in that case the polymer particles formed had dimensions of >10 nm, hampering the formation of a microporous membrane structure in a later stage of the process. The complete sol preparation procedure for hydrophobic membranes was as follows: TEOS was mixed with ethanol and placed in an ice-bath to avoid premature (partial) hydrolysis. A mixture of acid and water was added under vigorous stirring. After addition the reaction mixture was heated for 2 90hr at 60~ in a water bath under continuous stirring. The reaction mixture had a molar ratio (based on unreacted components) TEOS/ethanol/water/acid of 1/3.8/6.4/0.085. MTES was mixed with ethanol in the ratio of 1:3.8 and placed in an ice-bath. This mixture was added to the TEOS reaction mixture that has refluxed 2 90hr. The resulting MTES/TEOS reaction mixture obtained was heated for another 15 min. at 60~
The mixture then had a molar ratio MTES/TEOS/ethanol/water/acid (based on unreacted
components) of 1/1/7.6/6.4/0.085. Subsequently, the resulting sol was cooled and diluted 19-fold with ethanol to obtain the final dip-coat solution. After coating the membranes were calcined at 400~ for 3 hrs in pure N2 using a heating and cooling rate of 0.5~
Some active coal pellets' were placed in
the vicinity of the membranes to capture traces of oxygen in the N2 stream. Calcination was performed under a constant nitrogen flow (instead of air for the standard membranes) to avoid premature oxidation of the CH3 groups. The whole process of dipping and calcining was repeated once to repair any defects in the first silica membrane layer. The membranes obtained in this way are henceforth referred to as "MeSi(400) membranes". Unsupported silica material Unsupported microporous silica material was made for characterisation by means of physical sorption measurements as follows [2]: 60 cm 3 of 19x ethanol-diluted, hydrolysed silica sol was allowed to evaporate in a 10 cm O petri-dish at room temperature, so that 0.1-0.3 mm thick silica gel flakes were obtained overnight. These flakes were calcined at 400~
or 600~
for 3 hours with a heating and
cooling rate of 0.5~ The unsupported (microporous silica) membrane material was characterised with Ar physical sorption at 87K to determine its micropore volume, porosity and pore size distribution t. Nitrogen sorption measuraments were performed to investigate the amount of hydroxyl groups present in microporous
Norit RGM 0.8, QualityA3687, Norit N.V., Amersfoort, The Netherlands. t Sorptomatic 1900, Carlo Erba Instruments, Milan, Italy.
339 materials. The physical gas sorption set-up was provided with a turbo molecular pump system* and an extra pressure transducer for the low-pressure range (10 -3 Torr to 10 Torr) to be able to determine microporosity. This was checked with measuring zeolites [ 18]. All samples were degassed at 300~ for 24 hours prior to the sorption experiments. The pore size distribution is calculated according to the Horvfith-Kawazoe method [ 19], combined with the 10:4 Lennard-Jones potential functions for sorption of Ar on SiO> Hydrophobicity The hydrophobicity of the unsupported membrane material is determined by measuring the hydrophobicity index H I = Xoc,a,,Jx,,a,e,. as described in [20,21 ]. For that purpose the sample was dried first for 12 hrs at 250~ in a pure Ar stream. After that an Ar stream containing defined and equal concentrations of water and octane was used to load the sample until saturation at a temperature of 30~
The
Ar, water and octane flow rates were controlled by mass flow controllers. The breakthrough curves of the individual components were obtained by on-line gas chromatography. Numerical integration of the normalised breakthrough curves provided the sample loading of water, Xwater, and octane, Xoctane. These values were corrected for background signals, originating from the reactor. Measurement of breakthrough curves of water and octane resulted in/-//--0.3 for the unsupported Si(400) material and HI=3.0 for the MeSi(400) material [1,15]. The methylated unsupported membrane material is thus very hydrophobic whiie the standard silica material is strongly hydrophilic. The value of HI=3.0 for amorphous microporous silica is similar to a value of HI=2.9 found by Klein et al for a methylated silica-titania hybrid material [20]. For zeolites, however, higher values are measured such as H I = 1 0 . 3 for Silicalite I [21 ] which offers perspectives for further improvement of our silica material.
*Turbotronik, NT50 Leybold,Germany
340
Figure 1:
Drop of water on (A) MeSi(400) and (B) Si(400) membrane [1,15]
An impression of the difference in hydrophobicity of the membranes could also be obtained directly by putting a drop of water on both membrane types and observing the difference in curvature of the drops. As can be seen in Figure la water drop becomes more spread out on top of a Si(400) membrane than on a MeSi(400) membrane. This confirms that the MeSi(400) membranes are more hydrophobic than the Si(400) membranes. Thermogravimetric analysis Thermogravimetric analysis (TGA)* was performed on unsupported silica material to obtain a qualitative impression of the amount of hydroxyl groups at the surface. The TGA samples were stored in normal air at room temperature and hence at normal relative humidity before measurement. Unsupported material, made from the standard dip solution [Si(400)], but not calcined was examined by TGA to determine the burnout of the organic groups. The TGA experiments were performed with a heating rate of 1~
to 800~ in a pure N2 stream with a water and oxygen content less than 5 ppm.
The thermogravimetric experiments demonstrated a clear difference between the thermochemical properties of Si(400) and MeSi(400) materials as can be seen in Figure 4.
*Type 1136, Setaram, Lyon, France.
341
100 MeSi(400)~
98
A
U)
o
.,(=
.
.
.
.
.
.
.
.
.
.
.
.
.
.
.
.
.
.
.
.
.
.
.
.
96
o~
94
.~
92
~
9O
88
Figure 2:
Si(nr
,,,
0
I
,,
200
I
I
I
400 600 Temperature (~
800
,
1000
TGA, relative weight loss vs T for MeSi(400), Si(400) and dried silica sol [1,15]
MeSi(400) material does not show any weight loss up to 500~ starts to loose weight below 100~
while the Si(400) material already
This low-T weight loss of Si(400) is ascribed, as usual, to evapo-
ration of physisorbed water [22,23]. The lack of such a low-temperature weight loss for the MeSi(400) material indicates that in this material water sorption from ambient air hardly occurs. The total weight loss of both materials is quite different as well. Si(400) shows a weight loss of about 2%, caused by the loss of adsorbed water and surface hydroxyls at elevated temperatures. MeSi(400) shows a weight loss starting at much higher temperatures that saturates at 4% around 800~ cluded to be thermochemically stable until-500~
Hence MeSi(400) is con-
in pure N2. The weight loss of MeSi(400) in N2 at
higher temperatures might be due to the loss of incorporated methyl groups with the formation of CH4 or H2. Theoretically a weight loss of 12% is expected if all methyl groups, that are initially introduced in the synthesis, are removed in this way. It was found that after the TGA experiments the MeSi(400) material had turned black while the Si(400) material remained white. This led us to the conclusion that after the TGA run not all carbon atoms of the CH3 groups were removed from the MeSi(400) material and that it is likely that thermally induced condensation of methyl groups had occurred, for instance by:
~Si~CH
3 +
H3C~si"
/
~
~
~Si~CH2~Si~
~
4"
OH4
(1)
If, for instance, every two methyl groups combine, as happens in reaction (1), a maximum weight loss of 6.4% is expected. The fact that the actual weight loss observed is much less than 6.4% can be ex-
342 plained by the rather low concentration of methyl groups. This makes it unlikely that all methyl groups can participate in a reaction such as (1). Infra-Red Spectroscopy The presence of methyl groups in the microporous silica structure of the hydrophobic membrane material was demonstrated with infrared (FTIR) spectroscopy'. For that purpose, a KBr pellet was made of 20 mg unsupported material and 200 mg KBr. The pellet was heated in a pure Ar stream in an IR-cell t with KBr windows at 400~
for 20 hrs to remove water and weakly bound surface hydroxyls. The
spectra were recorded at 30~ (200 Scans) in the diffuse reflectance mode and represented by application of the Kubelka-Munk-function [24]. In Figure 3 the IR-spectra are given of unsupported Si(400) and MeSi(400) material. In the methylated material a sharp extra absorption peak is found around 1280 cm ~. This peak is ascribed to a symmetric deformation vibration of the CH3 groups [24].
*IFS 46, Bruker, Ettlingen, Germany. t HVC/diffuse reflectance unit DRA-XX, Harrick Scientific Corporation, Ossining (NY), USA.
343
O0
MeSi(400) m
mm
C C
1,0
m
m
|
'
9
9
m
CH3
tt~
t~ m
03
r
0
m
!
1500
1400
1300
1200
1100
Wavenumber Figure 3
1000
900
I
800
c m "1
IR spectraof MeSi(400) and Si(400) [1,15]
?,-Alumina intermediate layer ),-Alumina membranes are prepared by dip coating sintered a-alumina supports (either flat or tubular) in a home-made boehmite (),-A1OOH) sol. The boehmite sol is prepared by the colloidal sol-gel route as follows: 70 moles of double-distilled water are heated to about 90~
and 0.5 mole of alumin-
ium-tri-sec-butoxide (ATSB)* is added drop-wise under a nitrogen flow to avoid premature hydrolysis. The temperature of the reaction mixture should at least be 80~ to prevent the formation of bayerite (AI(OH)3) [25]. After the addition of ATSB is complete, the mixture is kept at 90~ to evaporate the formed butanol. Subsequently, the solution is cooled down to about 60~
at which temperature the
boehmite mixture is peptised with HNO3t at pH of 2.5. During the full synthesis the mixture is stirred
Acros, 97% purity, Belgium. +E. Merck, Darmstadt,Germany.
344 vigorously. The peptised boehmite mixture is refluxed for 20 hours at 90~ resulting in a homogeneous and stable 0.5 molar boehmite sol. During refluxing the pH increases to 3.5. If the sol is peptised at pH of 3.5 before the 20 hours of refluxing, the final pH will be about 4.4. This will result in fast aggregation of boehmite particles and subsequent precipitation of part of the boehmite from the sol. This lowers the concentration of the remaining sol, thus making it unusable for membrane preparation. Before dipping the sol is mixed with a PVA" solution of 3 grams PVA per 100 ml 0.05 HNO3., prepared at 80~
This "dip solution" has a PVA : boehmite ratio of 2:3.
Dip coating is performed under class 100 cleanroom conditions in order to minimise particle contamination of the membrane layer. After dipping, the membranes are dried in a climate chamber* at 40~ and 60% R.H. that is situated inside the cleanroom. The drying rate at such conditions is sufficiently low to avoid any crack formation in the boehmite layer [26]. Standard )'-alumina membranes are formed by firing at 600~
for 3 hours in air with a heating and cooling rate of l~
The total
,/-alumina layer thickness is in the order of 3 ~m, with an average Kelvin radius of 2.0 nm, as determined by permporometry [27,28]. Hydrothermally stable )'-alumina membranes An important drawback of the "standard" )'-alumina membranes described above is that they are not stable towards steam atmospheres that are, for example, used in steam reforming. For a membrane steam reformer, normal operation conditions are: 600~
30 bar gas pressure with a ratio H20:CH4 =
1:3, these conditions are further indicated as SASRA (Simulated Ambient Steam Reforming Atmospheres). SASRA conditions were found to be highly detrimental to our standard ~/-alumina membranes: the complete ),-alumina layer peeled-off within several hours. It was discovered that the peeling-off of the ),-alumina layer is due to insufficient adherence of the layer to the a-alumina support. The use of supports with a less smooth surface, to enhance mechanical anchoring, can not solve this problem. Hence a chemical "anchoring" material is coated onto the a-alumina support before coating the ,/-alumina layer. This treatment leads to )'-alumina membranes that do not show any degradation under steam reforming conditions anymore [29]. The treatment is performed as follows: A Mono-Aluminium Phosphate (MAP) layer is coated on the supports according to the following procedure: A commercial 50 wt-% MAP solution s is diluted either 10 or 20 times, further indicated as MAP 10 and MAP20, respectively. The shiny surface of a flat support is brought in contact with this solution for 3 seconds, after which it is dried. Next to this pre-treatment, the supports are coated under class 100 clean room conditions with either pure or doped 0.5M boehmite sols as described above. *MW 72.000 (g/mol)P.a., Merck, Darmstadt, Germany. *Heraeus V0tsch, Balingen, Germany. ~tAlfa, Johnson Matthey GmbH, Karslruhe, Germany.
345 The extent of degradation of the membranes is related to the tensile strength at the or/y-interface at which peeling-off or blistering is observed. Membrane degradation (peeling-off or blistering) was here tested by the so-called Scotch Tape Test [31 ]. In this test a piece of Scotch Tape is stuck onto the membrane surface and torn off rapidly. If the layer is of good quality it will not be torn off together with the tape. For the membranes with sufficient adherence, the change in pore-size during steamreforming treatment is measured with permporometry Standard y-alumina membrane-layers on ~-A1203 supports always came off in the Scotch Tape Test after SASRA treatment. When the support was treated with MAP, however, after SASRA treatment no delamination was observed. We suggest that the beneficial effect of MAP treatment results from chemical bonding between the membrane-layer and the support. The concentration of the MAP solution is found to be critical. Treatment with 5 wt-% MAP-solution gave good adherence, while a 2.5 wt-% solution resulted in some delamination, possibly due to insufficient phosphate on the surface of the supports. To reduce possible pore growth during the steam reforming treatment, the y-alumina membranes are sintered at temperatures much higher than usual. Such high temperatures, even up to 1000~
are pos-
sible provided an appropriate amount of lanthanum doping is present. The stabilising effect of lanthanum doping is well known [32,33]. Doping of the boehmite sol is performed by thorough mixing with the appropriate amount of a 0.3M lanthanum nitrate solution. The mixing is done directly before coating to avoid possible ageing effects that have been reported in the literature, for example by Lin and Burggraaf [32]. No such ageing studies are, however, performed in the present work. After the pore-size is established, the membranes are SASRA treated in a steel reactor. Heating and cooling is performed in an argon atmosphere at the same total pressure at a rate of 1~
In a few
experiments a pure steam treatment is carried out at 0.2 MPa total pressure at 150~ or 300~ in the same manner as for SASRA treatment. A pure CO2 treatment is done likewise, but at 500~
at
1.2 MPa pressure. Table 1 summarises the most important results from the investigation of metal doping. In this table the results of MAP treatment are combined with effects of firing temperature and doping. As can be seen in Table 1, y-alumina membranes with pore radii as low as 2.0 nm (Kelvin radius) may be obtained after firing at 600~
Note that, since an instrumental standard error of 0.5 nm (90% reliability) is
common in permporometry this technique should only be used for comparison purposes and to obtain a qualitative impression of the pore-size and pore-size distribution of the material under investigation.
346
Support Treatment
Tcalc(oc)
Test conditions
rKe,v,, (nm)
None
600
None
2.0
None
825
None
3.6
None
100(
None
8.7
MAP 10
825
None
4.2
MAP 10
825
SASRA
6.2
MAP 10
825
2xSASRA
7.5
y +3La
None
825
None
3.3
V+ 3La
MAP 10
100r
None
8.4
y+ 3La
MAP 10
100s
SASRA
9.3
y + 6La
MAP 10
100(~
None
6.0
y + 6La
MAP 10
1000
SASRA
6.1
~,+ 6La
MAP 10
1000
150~ steam
6.0
y + 6La
MAP 10
1000
300~ steam
6.0
? + 6La
MAP 10
1000
C02
6.3
V+ 9La
MAP 10
1000
None
8.6
Membrane
Table 1
Influence of support treatment, y-alumina doping, membrane firing temperature and SASRA-treatment on the pore-size of y-alumina. MAP 10 indicates a 10 times diluted standard MAP solution, which results in an effective MAP concentration of 5 mol-%., 3La indicates a 3 mol-% La-doped membrane, 6La indicates a 6 mol-% La doped membrane.
The pore-growth o f undoped "/-alumina strongly depends on temperature with a large increase in poresize between 825 and 1000~ 825~
The MAP-treated membranes have somewhat larger pores after firing at
The cause o f this effect is not clear yet. For undoped '/-alumina membranes, the pores g r o w
during S A S R A from 4.2 to 6.2 nm, and after a second S A S R A treatment to 7.5 nm. Thus, it appears that the pore-growth continues within the time scale o f our S A S R A treatment experiments. C o m p a r e d to undoped materials, 3 mol-% lanthanum doping gives hardly any beneficial effects on stability (Table 1). A significant improvement is found, however, for 6 mol-% lanthanum doping. For this case a pore-size o f only 6.0 nm is found after firing at 1000~
and no pore growth during S A S R A
347 treatment is observed at all. Additionally, after SASRA treatment, the pore-size distribution of a 6 mol-% doped y-alumina membrane is still very narrow, as can be seen in Figure 4.
6.E+17 In (D
5.E+17_
8v
4.E+17 _
"6 3.E+17
_
t_
,I2
2.E+17 1.E+17
t
0.E+00_
A
.4k A
0
10
A
4k
IA
20
30
Kelvin radius (nm)
Figure 4
Pore size distribution of a SASRA-treated,/-aluminamembrane. The support was treated with 5 mol-% MAP (MAP 10). The ,{-aluminawas doped with 6 mol-% La and sintered at 1000~ for three hours.
As one can see from Table 1, a spin-off result of this work is a list of recipes for the preparation of membranes with different amounts of doping, covering a complete range of pore-sizes with a resolution of 1-2 nm. This shows that we are now able to produce membranes with a tailor-made pore-size, which may be important for retaining certain large molecules by high-flux nanofiltration.
Flat supports
Flat supports are relatively easy to prepare. In our group two different methods are used. The first method is die pressing of a commercially available spray dried m-alumina powder*. The resulting disk is then pressed isostatically at 4000 bar. Final sintering is performed at 1260~ for 3 hours. The second method is the so-called colloidal filtration method. In this method a colloidal suspension is made of pure alumina powder [AKP30 or AKP-15t]. A 50wt-% suspension is obtained by dispersing the a-alumina powder in a 0.02M nitric acid solution [for AKP-30 powder] or a 0.02M nitric acid solution, mixed with Poly Vinyl Alcohol PVA ~ (5 g/l) [for AKP-15 powder] and using of ultrasonic
*PAI, Philips, Uden, The Netherlands t Sumitomo Chemical Company,Ltd, Tokyo, Japan. E. Merck, Darmstadt, Germany.
348 treatment* for 15 minutes. The resulting suspension is filtered over polyester filters t, consisting of a biological mixture of cellulose nitrate and cellulose acetate, with a pore size of 0.8 ~tm using a waterjet evacuation. The resulting filter cake (cast) is dried overnight at ambient temperature and fired at 1100~
[AKP-30] or 1150~ [AKP-15] for 1 hour. After firing the supports are machined to the re-
quired dimensions and polished until a shiny surface is obtained. Support pore-diameters obtained are 80 nm for the AKP-30 supports, 120 nm for the die-pressed supports and 160 nm for the AKP- 15 supports.
Tubular supports The single-bore tube geometry is currently most-well known in inorganic membrane technology. However to enhance area/volume ratios, multi-bore tubes and hollow fibres have emerged. All large membrane producing companies, such as US Filter, Noritake and Mitsui are able to deliver multi-bore tubes with various geometries. More recently, hollow fibre [34] supports consisting of porous or-alumina ceramics have been developed by TNO/CTK in Eindhoven. In a number of cases multibore may be less suitable due to limitations on reactor lay-out and the possible complications with high temperature sealing. Sealing problems can be expected for the hollow fibre geometry as wellbut the largest difficulty that must still be overcome is finding suitable techniques for the application of separative layers inside the hollow fiber. Coating a layer on the outside of the fibre is much easier but has the drawback that such a layer is much more subject to damage. Hence for hollow fibres supports, application of permselective layers by CVI seems to be the most suitable technique. Porous ct-A1203 tubes are frequently used as support for inorganic membranes. The normal way of producing such tubes is by extrusion or isostatic pressing followed by sintering. These techniques are fully accepted for the production of dense ceramic tubes, but may be less suitable for the production of porous membrane supports. Especially the occurrence of unroundness, inhomogeneities and a considerable surface roughness may impose problems. For the application of defect-poor meso- and microporous membrane layers for gas separation [ 16,35] a very smooth inner surface together with a narrow pore-size distribution of the membrane support tube is needed as well [36]. To meet increasing demands on roundness, homogeneity and surface quality ceramic tubes can be made by centrifugal casting (CC) of colloidal particles [37,38,39]. In this process a ceramic powder is dispersed in a liquid with a stabilising agent, followed by rotating for some time in a cylindrical mould around its axis. The resulting cast is dried, released from the mould and slightly sintered. If particles are used with a narrow size distribution and a low degree of agglomeration one may expect the forma*Model 250 Sonifier, BransonUltrasonicsCorporation,Danbury, USA. t ME 27, Schleicher& Schuell, Dassel, Germany.
349 tion of a nearly random-close-packed (RCP) green compact [40]. This requires the use of a proper colloidal stabiliser at a concentration such that the particles stay well dispersed in the liquid but form a coherent rigid structure in the compact. Examples of possible stabilisers are nitric acid [41,42,43] or polyacrylate-based products [39,44,45]. If the concentration of stabiliser is too low the particles will already flock in the liquid and form a low-density compact that will exhibit a rough surface. At higher stabiliser concentrations the dispersion may become too stable so that the compact remains fluid-like [46] and redispersion might occur as soon as the rotation stops. At optimum conditions the compact shape will closely follow the cylindrical mould shape which can be made with roundness near to perfection. In addition the surface roughness of the inside surface of the compact can be expected to be of the order of the particle size. Sintering mainly serves to obtain sufficient strength by the formation of necks without significant grain growth and shrinkage. The starting ot-A1203 powders were the above-mentioned AKP-30 and AKP-15 with a mean particle size of 0.40 and 0.62 ~tm and a BET surface of 6.2
m2/gand 3.5 m2/g respectively. Both powders have
narrow particle size distributions of(1.5%<0.25 pm + 95%<1 ~tm) and (1.5%<0.27 ~tm + 89%<1 pm), respectively and a chemical purity of >99.99% as stated by the producer.
Figure 5:
AKP-30 tubes made by CC deposition: 1,3 with sinter warping and cracking defects ([APMA] = 417 kg/m3]) and 2,4 without visible processing defects ([APMA] = 167 kg/m3]).
350
To obtain tubes with 2 mm wall thickness and -~20 mm diameter, 120 gram of powder was mixed with different amounts of APMA (Ammonium PolyMethAcrylate aqueous solution, Darvan | C*) and distilled water. The mixture of water and APMA, 120 ml in total, was brought on pH = 9.5 by adding (-1.5 ml) concentrated ammonia t. The resulting suspension was ultrasonically treated for 15 minutes using a frequency of 20 kHz and a transducer output power of 100 W. With this suspension tubes were prepared with three different lengths: short, 6&10 cm, tubes in a home-built apparatus, using steel moulds and long tubes (16 cm) in a commercial centrifuge t using Delrin | moulds. The inner diameter of the tubes was -~20 mm. Before pouring the suspension into the moulds, the moulds were coated at the inside with a solution of Vaseline |
in petroleum ethertt (boiling range 40-60~
to obtain easy
mould release. The tubes were centrifuged for 20 minutes at 20.000 rpm and the remaining liquid was poured out of the moulds afterwards. The green tubes were horizontally dried inside the moulds in a climate chambertt for two days at 30~ and 60% RH. After drying the green tubes were removed from the moulds and sintered horizontally on a fiat support at 1150~ for 1 hour with a heating/cooling rate of 1~ To study the influence of the amount of APMA on the drying and sintering behaviour of the AKP-30 tubes a series with different APMA concentrations in the suspension was made.
R.T. VanderbiltCompany, Inc., Norwalk, USA. t E. Merck, Darmstadt, Germany. CEPA, GLE, Carl Padberg GmbH, Lahr, Germany. Du Pont de Nemours, Dordrecht, The Netherlands. **Elida Faberg6, Bodegraven,The Netherlands. tt E. Merck, Darmstadt, Germany. ~ Heraus V6tch, Ballingen,Germany
351
Results on the preparation of tubular supports
The suspension mixtures of 120 gram powder and 120 ml stabilising liquid were sufficient for two
Observations
[APMA] (kg/m 3) 0
No suspension possible
8
Low green strength; green tube difficult to re-
could be varied between 1 and 2 mm at least, de-
42
Better green strength; some surface roughness
pending on solid concentration. It was found that
83
Good green strength; some surface roughness
porous tubes, visually free of processing defects
loo
Ibid
tubes with a length of 16 cm (and a wall thickness of 2 mm). The wall thickness of the supports
lease
could be obtained only with a certain, optimum,
167
No visible processing defects
APMA concentration, [APMA], in the suspen-
250
Ibid
sion. With [APMA] below optimum drying
292
Reasonable quality; some sinter-cracking
cracks were observed after drying. At [APMA]
333
Some sinter-cracks; some surface roughness
417
Considerable cracks after sintering, warping
higher than optimum typical defects were ob-
and surface roughness
tained such as surface corrugation and excessive warping and cracking during sintering. Examples
Table 2:
Influence of liquid phase [APMA] on the quality of sintered porous AKP-30 tubes.
are given in Figure 5. Influence of binder concentration
The results of the study on the influence of the APMA concentration on the quality of AKP-30 tubes after sintering are summarised in Table 2. It was found that [APMA] = 167 kg/m 3 (addition of 20 ml APMA) gave optimal results. The AKP-30 results could be used to obtain rapidly the optimum [APMA] = 83 kg/m 3 (10 ml APMA) for tubes made from AKP-15 powders. Pore size dlstnbutlon tubular supports 140 -
i
120 100 o
E
8o
I
co
o-AKP-30
-~ 60 R
4o
g
a~ |
9AKP-151
LA 41'
v
20
40
60
80
100
20
140
radius (nrn)
Figure 6:
Pore-size distribution of AKP-30 and AKP-15 tubes made by CC at optimum conditions.
352 Properties ---...
The porosity of AKP-30 (AKP-15) tubes, made with optimum [APMA] was 42.5% (43.2%)* after firing at 500~
vL x
measured with the Archimedes
method by immersion in mercury. The sintered compacts had a porosity of 34.8% (34.5%). Their pore-size distributions, measured by mercury porosimetry t are given in Figure 6. The mean pore radius was found to be 60(92)nm. The surface roughness ~ of the tubes was found to b e - 0 . 2 5 gm
0.097mm Figure 7:
Roundnessdiagram of an AKP-30 centrifuged tube.
for the inside and -0.9 ~tm for the outside. The mean unroundness "~was --0.025 mm, based on a 100 point measurement; the unroundness diagram is shown in Figure 7. In this diagram the drawn line around the concentrical circles gives the deviation from a perfectly round object. Figure 7 shows a slight elliptic deformation, possibly caused by "gravitational stress" during sintering. For comparison: the unroundness of a typical" extruded ~-A1203 tube was measured to be 0.16 mm as shown Figure 8. The surface roughness of this tube was -~6 ~tm. Discussion and conclusions on the preparation of tubular supports With the CC technique excellent tubular membrane
vL
supports can be prepared with a very low surface
x
roughness (0.25 ~tm). The roundness of tubes was
/ /
//7I If;I~ ;tt',,, \
found to b e - 1 0 x better than that of a typical ex-
.d..,-~ ..
,
"\'~ \
,
truded tube. The roundness can possibly be im-
,ss; i
ifl
proved further if special attention is paid to mould roundness and drying and/or sintering is done ver0.1)97 mm
tically or in a rotating set-up. This roundness is very important for application in reactors. If the
Figure 8:
Roundnessdiagram for a typical extruded ot-Al203 tube.
tubes are glass-soldered in ring-shaped machined flanges, a good roundness may result in a minimal and evenly filled solder space between the tube and
"The numbers between parentheses refer to AKP-15 tubes. t Series 200, Carlo Erba, Milan, Italy ~tMeasured with Mitutoyo Surftest III, Mitutoyo Mfg CO.,Ltd., Tokyo, Japan. wMeasured with MC 850, Carl Zeiss, Oberkochen, Germany. **
ECN (1990), Petten, The Netherlands.
353
the flange. This, in turn, will generally result in a better sealing process, quality and stability. If deformable (graphite) gaskets are used in a removable flange, a large unroundness will result in a radially inhomogeneous stress distribution in the tubes near the sealing, increasing the risk of brittle fracture. The wall-thickness of 2 mm provides the CC tubes with sufficient mechanical strength to withstand gas and liquid pressures that are common in membrane technology. The measured pore diameters of 120 and 184 nm are well in the range used for mesoporous membrane preparation and it is expected that 3t-alumina layers can be applied on the supports by conventional dip-coating and without the need of further intermediate layers [ 16]. The minimum membrane thickness that can be obtained defect-free can be expected to be of the order of the support roughness. This leads to the conclusion that membrane thickness' on CC tubes can be 50x less than those on extruded tubes. This, in turn, may result in a large flow increase for gasses and liquids. The best tube quality was obtained with an optimum APMA concentration that is proportional to the specific surface area of the powders. In the present experiments (with ot-A1203 powders) the optimum ratio between [APMA] and specific surface area was found to be -0.03 kg2/m5. An [APMA] of 8 kg/m 3 only, showed to be sufficient for electrosteric stabilisation and a rather stable suspension but also resulted in some drying cracks and roughness on the inside tube surface. This is likely to be caused by the fact that the suspension is partly flocked, leading to a poor particle packing that densities significantly during drying. In addition it was found that the green strength was insufficient at low [APMA]. This can be ascribed to poor particle packing too but it is more likely that APMA acts as a polymeric binder. At optimum [APMA] tubes can be prepared with sufficient handling strength and no surface roughness or cracking during drying or sintering. With higher [APMA], the condition of the green state looks all right but significant warping and cracking is obtained during sintering. This observation can be explained best by the presence of internal stresses in the green state caused by green state handling, or thermal processing. These stresses neither relax nor lead to cracks in the green state because of the combined effect of a particle packing, close to RCP and a significant amount of interparticle bonding. Perspectives for the use of the CC-technique It may be questioned whether the high degree of perfection of the CC tubes justifies possibly higher cost prices in mass-scale production and the limitation to circular shapes. Extrusion processes can be relatively cheap and easily continuous and enable more complex shapes such as multi-bore tubes. On the other hand the CC technique allows a radial variation of composition and particle morphology. The optimum support structure for a given application can be obtained using design rules that take into account the desired support shape, the material's strength and permeability [13]. This optimum is achieved at a certain thickness and spatial distribution of porosity and pore size. The use of a suspen-
354 sion that consists of largely different sizes of particles may naturally result in specific desired radial variations that can be predicted quantitatively on basis of the method described in [47]. In addition it is possible to inject small amounts of suspension layer-wise by using a tangential injection technique [37] so that all thinkable radial distributions can be realised.
Characterisation of membrane morphology
FE-SEM / TEM
Morphological membrane characterisation was done by Field-Emission Scanning Electron Microscopy (FE-SEM)" and Transmission Electron Microscopy (TEM)*. The FE-SEM recordings were made of a perpendicular fracture surface. TEM recordings were performed on a thin cross-section of the membrane, made as follows: One silica membrane was cut in halves and both parts were glued together with the silica top layers facing each other. From this "sandwich structure" a small slab was cut, which was abraded to a thickness of 200 lam. The specimen was further thinned by making dimples on both sides, in the middle down to 15 lam using a Dimple grinder. Finally the thickness of the centre of the specimen was reduced further by ion milling until a centre hole had just appeared. During the thinning procedure the specimen was carefully positioned so that the opposing silica layers formed the centre of the thinned area. The TEM recording was made near the centre hole, at the thinnest part of the sample. The FE-SEM an TEM recordings, given in Figure 9 and Figure 10 reveal a very thin silica layer of ~30 nm, obtained after 2 times dipping. This result is a little in contrast with earlier suggestions in the literature [ 16] in which a thickness of 100 nm is proposed.
*Hitachi, Type $800. tCM30 TWIN/(S)TEM,PhilipsAnalytical,Eindhoven,The Netherlands.
355 The TEM micrograph indicates that the silica layer is deposited on top of the 7-A1203 layer, as a distinct separation between the two layers is present. The y-AI203 layer is about 3 lain thick after 2 times dipping. The boundary between the first and second y-AI203 layer at approximately 250 nm from the surface is clearly visible in TEM. There is a clear "colour" difference between the two y-AI203 layers visible in the TEM micrograph. The second 7-alumina layer has a lighter appearance than the first layer that can be caused by two effects: 9
The first layer is calcined two times at 600~
9
The first layer is applied on the ~-A1203 support that has a much courser structure than the y-AI203 layer on which the second layer is applied. Since capillary forces play an important role in the layer formation this can result in a denser second ]t-A1203 layer.
The colour difference is not caused by a difference in pore size. This was checked by permporometry in which the one- and two-layer dipped membranes
Figure 10:
TEM Micrograph of Si(400) membrane cross section showing a part of the yalumina intermediate layers and the silica top-layer [48].
were found to have both a Kelvin radius of --2.5 nm. The permoporometry method is based on the measurement of permeation of a noncondensable gas while at the same time pores are blocked by a condensable gas. This makes that for multi-layer stacks of continuous defect-free membranes, pore radii are obtained that correspond to those of the layer with the smallest pores. Pore size characterisation of unsupported silica The unsupported Si(600) material shows a type II sorption isotherm, which is typical for nitrogendense materials. This result is in agreement with the pore size of the supported Si(600) membranes, estimated from the relation between permeance of various gases and their kinetic diameter.
356 180 160 140
~
120 100
o)
~
8o
~
6o 40 20 0
Figure 11"
N2 sorption isotherm at 77K of unsupported MeSi(400) and Si(400) material [1,15]
In Figure 11 the N2 physical sorption isotherms of Si(400) and MeSi(400) at 77K are presented. For both materials the sorption isotherms are type I, characteristic for microporous materials. From the figure it is clear that more N2 can be sorbed by the Si(400) material. This could be a result of a difference in pore structure of the silica materials, but also of the difference in amount of hydroxyl groups at the internal surface of these materials: N2 is expected to show particular interaction with the hydroxyl groups [50], so that more sorption of N2 is expected for Si(400). 180
MeSi(400)
160 140 1. 120
.
.
.
.
S~400)
100
~
8o
~
6o 40 20 0 0
Figure 12
1
Ar sorption isotherm at 77K of unsupported MeSi(400) and Si(400) material [1,15]
Figure 12 shows the Ar sorption behaviour of Si(400) and MeSi(400). Since the Ar molecule is spherical and more "inert" than the N2 molecule', it is assumed that the Ar sorption is determined by the structure of the pores and is less dependent on the concentration of hydroxyl groups [49]. A striking
Besides the hydroxyl interactions already mentioned, the interpretation of N2 adsorption data is further complicated by quadrupolar interactions.
357
observation is that the sorption behaviour of Si(400) is similar for Ar and N2, i.e. the amount of Ar that can be sorbed is only slightly higher, while for MeSi(400) material much more Ar can be adsorbed than N2. This may be explained from the molecular dimensions of the gas molecules: Ar has a kinetic diameter, dk, of 0.340 nm and covers a sorption area of 0.133 nm 2, while N2 has a dk of 0.365 nm and covers a sorption area of 0.166 nm 2 [22]. This suggests that MeSi(400) contains many small pores that are not available for the larger N2 molecules. The pore size distribution, calculated with the Horvfith-Kawazoe method [19] from the Ar sorption isotherms at 87K is given in Figure 13. This plot shows that the pore size distribution of the unsupported Si(400) material is narrow with a maximum at Deff~7A. The distribution in pore size of the MeSi(400) material is much broader, but has a maximum at approximately the same Deff. It should be kept in mind that the results for unsupported material can not be transferred quantita-
0.2 ~9
tively to the supported membrane situation. They can only be of use to show trends in changes in
o
pore structure with processing. It must be clear that
~;
the supported silica layer cannot be expected to
E u
have the same structure as similar processed unsupported silica material, since the forces present during the drying process of both materials are different.
0.4
MeSi(400)
o.15
0.3
0.1
0.2
0.05
0.1
;i(400 "'~-.
0 0 Figure 13
1 2 3 pore width (nm)
"a
.
~ E
0 4
Pore size distribution for unsupported MeSi(400) and Si(400) material, calculated according to the Horv~ith-Kawazoemethod [19] from Ar adsorption isotherms at 87K
Transport
[1,15]
The movement or rate of mass transport, of a species in a microporous material is generally determined by two parameters: the amount of a species present in the material and the mobility of the species. A proper description of mass transport in a microporous medium such as the amorphous silica layers studied here should consequently include a good understanding of the sorption properties of the system, as well as the diffusion mechanism by which transport occurs.
Diffusion The silica layer may be regarded as an interconnected network of voids with a concentration
qSat. It can
be assumed that during sorption each translational degree of freedom of a gas phase species is convetted to a vibrational one. Sorbed species vibrate around a minimum in the potential field inside a
358 void. Transport occurs when molecules jump from a minimum in one pore to a minimum in a neighbouring one. In such a system the species obey the statistics of a lattice gas. Irreversible thermodynamic description of single species permeation fluxes According to the theory of irreversible thermodynamics [51 ], the flux J of a single species i in the silica material is proportional to the gradient in its chemical potential
J, = -L.Vl.t,
(2)
The phenomenological (macroscopic) Equation (2) can be worked out further in terms of experimental quantities if the adsorption equilibrium and the mechanical mobility are studied in more detail including microscopic information. In the present paper we assume that the relevant adsorption data can be described with sufficient accuracy by the Langrnuir isotherm. This implies that the silica micropores can be either occupied by one molecule or empty and that the bonding energy of the molecules is independent of @. In that case a simple expression is obtained for the chemical potential of species i
(3)
where ~t,.0 is the concentration-independent part of ~t, and O, = q,/qSat is the fraction of occupied voids. In fact, this is the chemical potential of a so-called "building-unit" [52], which reflects the fact that species i the vacancies are conjugated. Combination of (3) and (2) results in
(4)
The direct coefficient Lii is the product of the mobility bi of the species and the concentration
qsat'0i.
When the mobility is expressed in terms of a component-diffusion coefficient, 1), = b~RT, the expression for the flux can be written as
ji=_q~at
D, VOi
1-0,
(5)
The component-diffusion coefficient is not necessarily constant. In fact, it seems plausible that it will decrease with rising occupancy of the voids, since a jump of a species to an adjacent occupied site will not take place. The mechanical mobility will be proportional to the probability that an adjacent void is empty, b, = b~(1- 0, ). In that case the flux obeys a, =
-13,vo,
(6)
359 where Di is the constant single component chemical diffusion coefficient and is related to the uncon-
ditional mechanical mobility, b~ , through /9, = b~ through/9i
=
(and to the component-diffusion coefficient,
Di )
Microscopic description of diffusion mobility In the derivation of equation (6) we already included some information about the system on a microscopic level. To obtain more information about the factors that influence diffusion mobility the flux expression must be derived with a less general but entirely microscopic molecular jump method [53]. When transport takes place by species jumping from void to void, the net flux of a species is determined by the overall number of successful jumps in the direction considered and the average jump distance a. The number of successful jumps is determined by: 9
a geometric factor g that is representative for the spatial arrangements of voids and their (percolative) connectivity,
9
the concentration of the species 0i,
9
the frequency at which a species attempts to jump v,(T)
9
and a Boltzmann factor that determines the fraction of species with sufficient energy Em., to cross the potential bridge between two adjacent voids.
In the case of single site occupancy a jump will only be successful if the adjacent site is unoccupied. The flux expression obtained from a microscopic approach is identical to (6), but the chemical diffusion coefficient contains information about the system on a microscopic level
Em ~)= ga2vA(T)e Rr
(7)
It is clear that also in this case it is found the chemical diffusion coefficient is not dependent on concentration, for the simple lattice model.
Sorption and permeance In the derivation of the diffusion flux equation it is assumed that the sorption behaviour can be described with the well-known Langrnuir-isotherm
0,-
K,p 1 + Kip
For the T- and p-range used in this study the Langmuir isotherm reduces to the Henry isotherm
(8)
360
q, = K,p
(9)
If the rate of transport through the silica layer is limited by diffusion through the silica layer, thermodynamic equilibrium can be assumed at the interfaces, and the permeance is F, =
J----J-'-
Ap,
sat ~-~ r.,-
q u,t~,
L(l + K,p,,h)(l + K,p,,,)
(10)
where P,.h and p ,.1 are the high (h) and low (1) partial pressures of i at both sides respectively of the silica layer and L is the thickness. (In the derivation of (10) from (6) and (8) we made use of the fact that D, is independent of 0, so that for the case of stationary diffusion (6) can simply be integrated to J, =-/9, 01'' -Oh" ). In the Henry regime the denominator of (10) reduces to L and the permeance 8 becomes independent of pressure. The validity of Henry's law implies that the concentration of molecules absorbed in the microporous solid is small compared to the number of available voids.
Temperature dependence of permeance Diffusion in the silica layer is an activated process, since species have to pass a potential bridge between two adjacent voids. Furthermore, the attempt frequency v, depends on the vibrations of both the species in the voids, as well as their surrounding molecules in the solid and is thus also a function of temperature. Era.,
D, (T) = ~)~o(T)e RT
(I I)
For relevant temperatures it is safe to assume that the temperature dependence of the pre-exponential factor is negligible compared to the influence of the exponential factor. The temperature dependence of the sorption process is expressed in the Langmuir coefficient
Q~._.._~,
K, (T) = Kf ~( T)e Rr
(12)
where Qst., is the isosteric heat of sorption. The pre-exponential factor contains information such as the partition functions for the vibrational or rotational modes of a species and is thus slightly temperature dependent. It is again safe to assume that its influence on the temperature dependent behaviour is small compared to that of the exponential factor. In experiments an "effective" temperature dependence of the mass transport process will be observed. In the Henry regime this temperature dependence will be described by
361
sat ,-~
Qst,,-Em,,
F, = q--c-Df~ L
Rr
Ea,,
= jo e Rr
(13)
in which we define a new (nearly) temperature-independent proportionality constant J0 =/~r,,Kr,, and an effective activation energy for permeance:
Ea -- Em - P s t
(14)
Ea can have any sign; here a sign convention is used such that J increases with temperature if Ea is positive. Since Em and Qst can be expected to be about the same the value of Ea may well be close to zero so that slightly temperature-dependant factors in J0 must be taken into account to obtain a proper quantitative expression for the temperature dependence of J.
S u p p o r t resistance
The overall permeance of the membrane will be determined by the resistances for mass transport of all subsequent layers it consists of. These resistances are considered to be connected in series. 1
Foverall
1 - ~ 1+ ~ +1 ~ Fo~-layer F~-layer Fsilica
(15)
Due to the different structures of the layers, each layer will influence the motion of a species in a distinct manner, resulting in disparate mass transport behaviour in each layer. The mass transport all three layers is generally a function of pressure. In both the o~- and ~/-layer, a gas phase species has three degrees of translational freedom and transport will take place by a combination of viscous flow, Knudsen diffusion and, in the case of a multicomponent system, bulk diffusion [54]. Furthermore, for a strongly adsorbable gas diffusion over the pore surface may take place. A quantitative mathematical description of such mass transport involves solving a set of coupled non-linear partial differential equations, for which one needs to resort to numerical techniques [55]. Here a only a single gas is present, the contribution of viscous flow is only small and the temperature is sufficiently high for surface diffusion to be negligible. As a result the permeance F of the ix- and y-layer will only depend on the molecular weight of the permeating gas and the temperature, i.e. F-x/T-.xf-M can be expected constant. For reasons of simplicity the support will be treated as one effective layer. Since for most gasses the largest resistance for mass transport is situated in the silica layer this approach is justifiable.
362
Transport measurements Membrane gas permeance was measured in the pressure-controlled dead-end mode [56] in the temperature range of 50 to 300~
Prior to the permeance measurements the membranes were dried for
several hours at 200~ to remove adsorbed water from the micropores. The disk-shaped membranes were placed in stainless-steel permeance cells with the microporous top-layer at the feed side. The pressure difference over the membrane was adjusted by an electronic pressure controller.' The gas flow through the membrane was measured by mass flow meters with maximum flow ranges of 25 or 100 (cm3/min SPT). The pressure over the membrane was measured with an electronic pressure transducer.* A schematic representation of the permeance set-up is given in Figure 14.
...... j
Gas feed
i i = i '.
9
i
PC FI PI TC
~
= Pressure controller = Flow Indicator = Differential Pressure indicator = Temperature controller furnace ;
L.. . . . . . .
.J
Furnace and test cell
my
To atmosphere
"
;
,
=
Figure 14
~
.
"
Schematic diagram of the experimentalset-up for permeancemeasurements.
Supports The dead-end single gas permeance of the small gasses H2, CO2, 02, N2 and CI-I4 through the die pressed support was measured as a function of pressure and temperature. Data obtained from these measurements show that the permeance of the support is pressure-independent. The product F-x/-f.x/-M, (mol.K-kg)~
the so-called modified permeance, is for all gasses approximately 9.8-10 -5 and shows only little variation. From this it may be concluded that in the sup-
port diffusion takes predominately place by Knudsen diffusion. Only for CO2 the modified permeance is observed that is higher than that for other gasses while it increases slightly with pressure. This can be explained by the occurrence of diffusion over the pore surface of the support.
*Type5866, Brooks Instruments,USA. tValidine Inc., Northridge(CA), USA.
363
Si(400) Single gas permeance results for Si(400), corrected for support resistance, are independent of pressure. This was expected since the sorption occurs in the Henry regime. The highest permeance is observed for the smallest gas used in this study, i.e. F m = 4.4-10 -6 mol/m2-s-l.pa-1 at 200~ F,~
Si(400)
Si(600)
H2/CO 2
13
22
H2/O 2
1.6.102
43
H2/N 2
41
H2/CH 4
>>500
In Table 3 the permeance data of the other gasses compared to H2, at T=200~
are presented. Al-
though the kinetic diameters of 02 and N2 are quite similar, 3.46A and 3.65A respectively, our microporous membranes have an 02 permeance that is approximately 4 times higher than that of N2. The permselectivity F~ >>500 for H2/CH4 can not be determined more accurately, since the
Table 3:
Permselectivitiesat T=200 ~ and p=l bar
permeance of CH4 is very low and consequently difficult to measure. The fact that CH4, with a dk=3.8A has some detectable permeance and that SF6, with a kinetic diameter of 5.5A, does not permeate at all leads us to an estimate of the membrane pore size between 3.8
3.8 CH4
\
289H 2 \
\ Si(400)
3 46 02
82 \ "5 E
5000
-5000 2.5
\
\
\
\
\
//
/
O
3.64 N2
d //
\
/
\ ~ 3 3 CO2 3.0
35
40
ak [A]
Figure 15:
Apparent energyof activation(Ea)versus kinetic diameter(dk).
For all gasses, except for C02, the permeance increases with temperature, i.e. the apparent energy of activation (E,) is positive and hence the activation energy of mobility heat of sorption
(Era) is larger than the isosteric
(Qst). The CO2 permeance appears to decrease with T at higher temperatures. This
demonstrates that E, is not necessarily always positive. The apparent energy of activation for the different gasses is presented in Figure 15. The permselectivities for H2/CO2, H2]N2 and HJO2 increase with increasing temperature due to the fact that the apparent energy of activation Ea is largest for H2.
364 Due to the inaccuracies in the measurement of the low CH4 flow, no systematic temperature dependence of the permselectivity for H2/CH4 could be determined. Using the experimental data for E,, together with typical Q~, data, shown in Table 4~ values for Em are obtained by making use of equation (14). The results are given in Figure 16, from which it is clear that E,, increases with the kinetic diameter dk of the gas. As dk increases the molecule has more difficulty to jump from void to void. Due to the low accuracy of the CH4 permeance data, the absolute value of Em for CH4 will also be imprecise. It is clear however, that CH4 has the greatest difficulty moving through the pores and its mobility energy is much higher than that of the other gasses. Si(600) The Si(600) membranes show a permeance behaviour that is clearly different from that of Si(400). The permeance at 200~
CH4
CO2
H2
De Lange [571
I0
22
6
Rees 1581
20.0
24.6
Golden [59]
18.6
24.0
Choudary 136]
28
20
Dunne 1601
20.9
27.2
Used in this study
20
24
N2
02
of H2 is lower,
2.4.10 -7 mol/mLs-~-Pa"I and at the same temperature the CO2 permeance is even much lower, i.e. 1.1.10 "8 mol/m2.sl-Pa -l, resulting in F~=22 for H2/CO2. The 02 permeance has reduced to 6.1.10 -9 mol/m2.s-l.Pa -~, at 200~
Source
and the per-
meances of both N2 and CH4 were not sufficiently high to be determined with a reasonable accuracy. The observation that the supported Si(600) mate-
Table 4:
17.3 6.0
6
15.0
17.6
16.3
17
16
Microporous silica isosteric heats of adsorption, Qst (kJ/mol) at low coverage.
rial is not 'open' for N2 is consistent with results of the sorption behaviour of the unsupported Si(600) material described in the section "Pore size characterisation of unsupported silica" of this chapter. Using a very sensitive qualitative soap-solution test, some N2 flow was detected. Even with this very sensitive test, no permeance of CH4 was observed. Determination of the pore size by size exclusion by means of the permeance experiments thus results in a pore size 3.6<0<3.8 A which is significantly smaller than that of the Si(400) membranes.
365 35000
38 0
30000
"--'
%
25000
3.46 02 / 3.3 C02.I ;;~ ~ --0
E
Si(400)t 101 /
LLIE 20000
2.89 H2
Jl
15000
I0000
--
2
5
/
364 N2
J
o-/
/
CH 4
/
/
/
//Si(600)
3.0'
3'5-
40
d,, [A]
Figure 16:
Calculated energyof activationfor mobility(E,,) versuskinetic diameter(dk).
The decrease of the permeance is presumably a result of densification of the structure and a smaller pore size. The large increase of F~ for HJCO2 from 13 to about 22 can also be, partially, attributed to a decrease in the amount of terminal hydroxyl groups at the internal surface of the silica, since at higher calcination temperatures more hydroxyl groups are irreversibly removed [22]. It is recognised that the sorption behaviour of the silica material is strongly influenced by the concentration of hydroxyl groups due to gas molecule interactions with hydroxyls at the silica surface [49]. The lower (surface) occupation leads to a lower CO2 permeance. Due to the lower concentration of hydroxyls, the Si(600) material is also more hydrophobic compared to the Si(400) material. The apparent energy of activation E, for H2 and CO2 is lower than for the Si(400) membranes, while that of 02 is approximately the same. Since E,=E,,-Q~., the lower E,, may be explained by either a smaller E,, or a higher Q.~.,or a combination of both. For the time being it is not possible to confirm which of these hypotheses is valid. When taking into account the fact that the Si(600) is much denser than the Si(400) structure a larger Em and a smaller Q~t are both possible for Si(600). In Figure 16 values for E,, are presented, that have been obtained under the assumption that the Q.,., is the same for both the Si(400) and Si(600) materials. With increasing kinetic diameter of the gas molecule again an increase in Em is observed. The value of E,, is lower for Si(600) than for Si(400), suggesting that molecules are more mobile in this material. This is not sensible, taking into account the denser structure of the material and it is obvious that the lack of experimental data on Q~., is prohibits a proper understanding of the temperature dependence of the permeance behaviour of both silica materials discussed here. MeSi(400) Compared to the non-hydrophobic silica membranes the permeance data for MeSi(400) shows a remarkably different behaviour. The fluxes are quite high while the permselectivity is quite low. This
366 suggests that the membranes are either not microporous or contain many defects. Since the extremely low permeance of SF6 (dk=0.55 nm) and n-C4H10 (dk=0.5 nm) does not correspond to a large concentration of defects, it is expected that the MeSi(400) membranes have a larger average pore size or broader pore size distribution. This is in agreement with the Ar sorption measurements described in the section on characterisation of this chapter, where it was found that the unsupported Si(400) and MeSi(400) have approximately the same average pore size but a different pore size distribution. Comparison with open literature data The permeance and F~ values obtained compare favourable with literature results. De Lange et al. [57] previously reported values of Fro=7.4 - 10-6 (mol/mLs.Pa) and F~=3.9 and 24.7 for H2/CO2 and H2/CH4 respectively. Although the H2 permeance reported by the Lange is higher that that of our membranes, the selectivity is significantly lower. An F~ =235 for H2/CH4 at 150~ was published by Hassan et al. [46] for a silica hollow fiber membrane, but no single gas permeance values were specified. Commercially available CVI silica membranes* onpage1 are reported to have a F~ = 27 for H2/CO2 with a H2 permeance of 1.7x10 7 (mol/m2.s.Pa) at 400~ *.
*Measured in our lab on membranes deliveredby Media and Process TechnologyInc. (MPT), Pittsburgh, PA, USA.
367
le-5
S~(9
le-6
~
2
002
S" "7
13.. le-7 E 0 E ,,....._, LL le-8
q~ Si(600)
eSi(400)
E)
le-9
~3.8 CH 4
~,~so-C4H10
,
,
,
3
4
5
"o.5
SF 6
6
a k [A]
Figure 17:
Permeance of species with differentkinetic diameter, for Si(600), Si(4000 and MeSi(400).
Separation Experimental gas separation factors reported in [ 1] are of the same order of magnitude as permselectivities calculated from single-gas permeance experiments. This indicates that in transport of mixtures the molecules have only limited mutual influence, in agreement with the expected Henry sorption behaviour.
Conclusions
Homogeneous and clean TEOS-based sol-gel synthesis leads to reproducible high quality silica membranes with good separation characteristics. Si(400) membranes have a homogeneous 30 nm thick silica layer, which results in a permeance of Ha, corrected for the resistance of the support, Fm= 4.4.10 .6 mol/m2-s-l.pa1 at 200~
and F~(H2/CH4)>500. It is now possible to prepare intermediate layers at a
temperature as high as 1000~ with on top very thin high quality silica top layers fired at 400-800~ The membrane firing temperatures have an outspoken influence on the membrane properties. Increas-
368 ing the sintering temperature from 400 to 800~
results in a much denser membrane structure with
smaller pores. From the relation between kinetic diameter and permeance results, the Si(600) membranes are expected to have pores with 3.6<0<3.8 A and the Si(400) membranes to have pores with 3.8< O <5.5 A. These results are in agreement with physical adsorption experiments, performed on unsupported material, which reveal a pore size of~5A for Si(400) while the Si(600) samples were N2dense. All permeances are lower in the Si(600) membranes, e.g. H2 is 2.4-10 -7 (mol/m2.s.Pa), while permselectivities with respect to larger gasses are significantly higher. The best choice of the firing temperature depends on the envisaged application. In addition one should take into account that for high temperature applications the firing temperature of the membrane should be higher than the highest temperature during operation'. With the use of cleanroom conditions we were also able to reduce the number of layers for the intermediate as well as the top layer to just one layer each without any large defects. This reduces the preparation time for both layer types by a factor of two, which may save several days of processing time. By incorporation of methyl groups in the silica microstructure the surface and microstructural properties of the microporous silica membranes change significantly. 'MeSi(400) membranes' are 10• more hydrophobic and show much less water sorption than state-of-the-art silica 'Si(400) membranes'. MeSi(400) membranes have larger micropores with a wider pore size distribution than Si(400) membranes which influences their transport properties. Further development of hydrophobic silica membranes should focus on systematic optimisation of sol-gel and calcination procedures, studies of transport properties including water vapour and the introduction of hydrophobicity in the support structure. The separation characteristics of microporous silica membranes can be described very well for practical situations if accurate values of unconditional diffusion mobilities, b0, are available from single-gas permeation experiments. The effective activation energy, Ea, of membrane permeation is the result of counteracting thermal activation of diffusion mobility and sorption. This often results in relatively small values for Ea. This, in turn, makes that weak, non-Arrhenius-type temperature dependencies in the sorption equilibria can no longer be ignored. In addition it is difficult to quantify the diffusion mobility energy of gasses through the microporous structure if accurate Q,,-values are unavailable. Addressing these problems is the subject of future studies. A major problem that remains at this moment is the stability of silica membranes towards hot steam. Steam is used is many processes, in process industry (reforming, gasification, etc.) as well as in the medical and food industry (sterilisation). It would thus be very favourable if meso- and microporous membranes could be developed which are resistant towards hot steam environments. The steam-
This condition must be fulfilled if one has the objectiveof constantmembrane operationduring the complete lifetime of the process. The membrane can also be fired in-situ at the highest temperatureof the process. In any case the membrane properties can be expectedto change slowlywith time.
369 reforming conditions as used in our stability tests are among the harshest conditions one can think of. Hence if the membranes are stable under these conditions, they most probably can be used in any other steam-containing atmospheres as well. A large step forward has been made with solving the problem of the blistering of the y-alumina layer and the reduction of pore-growth and the formation of macroscopic defects in this layer during steam exposure. This result enables us now to test the stability of the silica layer proper against harsh steam-rich environments. The behaviour of the silica layer under steam-reforming conditions will therefore receive considerable attention in the near future. Improvements in high temperature resistance can possibly be achieved by doping with polyvalent metal ions or removal of silica surface hydroxyl groups. Additionally, the hydrophobic silica membranes that were developed recently may very well be more steam resistant than normal silica membranes since they have less affinity towards water molecules. The major part of H2 transport resistance in our state-of-the-art supported silica membranes is currently in the a-alumina support. This makes that more attention will have to be paid to the development of high-quality course porous membrane supports. The centrifugal casting technique equipped with an injection system might be very helpful to develop coarse porous multilayer supports with a very smooth inside surface on which highly selective membranes can be coated. The development of even thinner intermediate and top layers becomes worthwhile only when the transport resistance of the support is not limiting anymore. If membranes can be made with a thickness-1 nm, H2 permeances of > 10-4 (mol/ma.s.Pa) can be realised provided no support limitations or surface transfer rate limitations occur.
Upscaling of the present silica membrane technology into membrane modules is needed to increase the surface to volume ratio for practical applications. The processes of the silica membrane preparation by CVI techniques or by the wet-chemical methods presented here may very well be suitable for production scale manufacture of membranes on hollow fibre supports. The present results suggest that application of silica membranes in for instance natural gas purification, molecular air filtration, selective CO2 removal, industrial H2 purification and conversion enhancement in membrane reactors are feasible indeed.
Acknowledgement The authors are indebted to H. Kruidhof for technical support and supervision, C. Huiskes for technical assistance with the preparation of the tubular supports, prof. dr W.F. Maier and coworkers of the Max Planck Institut, Mtilheim an der Ruhr, Germany for infra-red and hydrophobicity measurements and dr R. Bredesen, SINTEF, Oslo for steam reforming treatments of the membranes.
370
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