Microporous and Mesoporous Materials 232 (2016) 196e204

Contents lists available at ScienceDirect

Microporous and Mesoporous Materials
journal homepage: www.elsevier.com/locate/micromeso

Asymmetric polysilazane-derived ceramic structures with multiscalar
porosity for membrane applications
Thomas Konegger a, b, *, Chen-Chih Tsai a, Herwig Peterlik c, Stephen E. Creager d,
Rajendra K. Bordia a
a

Clemson University, Department of Materials Science and Engineering, 161 Sirrine Hall, Clemson, SC 29634, USA
TU Wien, Institute of Chemical Technologies and Analytics, Getreidemarkt 9/164-CT, 1060 Vienna, Austria
University of Vienna, Faculty of Physics, Boltzmanngasse 5, 1090 Vienna, Austria
d
Clemson University, Department of Chemistry, 219 Hunter Laboratories, Clemson, SC 29634, USA
b
c

a r t i c l e i n f o

a b s t r a c t

Article history:
Received 15 October 2015
Received in revised form
10 June 2016
Accepted 14 June 2016
Available online 16 June 2016

Novel ceramic structures with multi-scalar porosity were developed using a single preceramic poly(vinyl)silazane to generate asymmetric Si-C-N-based membranes through pyrolytic conversion. Macroporous supports in planar-disc configuration were prepared through a sacrificial filler approach,
intermediate structures and microporous layers were deposited via dip-coating. Microporosity in the
selective layer was generated through a controlled thermal decomposition of the precursor component
in nitrogen atmosphere at temperatures up to 600  C, resulting in micropores with average pore sizes of
0.8 nm, as investigated by nitrogen adsorption and small-angle X-ray scattering (SAXS). The general
feasibility of the single-precursor approach towards selective permeation of gaseous species was
demonstrated by the investigation of gas permeances of the generated structures using single-gas permeance testing of He, N2, Ar, C2H6, and CO2. By variation of the deposition sequence during preparation
of the selective layer by dip-coating, asymmetric structures with ideal permselectivities exceeding
predicted Knudsen values were obtained. At 500  C, He/N2 and He/CO2 permselectivities of up to 3.1 and
4.1 were found, respectively, at He permeances up to 3 Â 10À8 mol mÀ2 PaÀ1 sÀ1. The new single-material
system is a first step towards the potential establishment of new, alternative membrane materials systems, circumventing thermal and chemical incompatibilities between constituents, and increasing material performance due to the applicability under extreme operating conditions.
© 2016 The Author(s). Published by Elsevier Inc. This is an open access article under the CC BY license
( />
Keywords:
Polymer-derived ceramics
Polysilazane
Microporosity
Macroporosity
Membranes

1. Introduction
Recent global challenges leading to the development of more
sustainable processes, calls for energy conservation as well as the
increased utilization of renewable energy sources have resulted in
increasing interest in the application of membranes for filtration
and separation processes.
Ceramic membranes present an alternative to conventional
polymer-based membranes due to superior thermal, chemical, and
mechanical properties, facilitating applications at temperatures

beyond 300  C or in harsh chemical environments. In contrast to
dense membranes, porous ceramic membranes exhibit different
mechanisms governing the transport of species through the

* Corresponding author. TU Wien, Institute of Chemical Technologies and Analytics, Getreidemarkt 9/164-CT, 1060 Vienna, Austria.
E-mail address: (T. Konegger).

membrane depending on the pore size. In macro- and mesoporous
materials, the main gas transport mechanisms include molecular
diffusion, viscous flow, and Knudsen diffusion [1,2]. While offering
high permeabilities, membranes operating under the Knudsen flow
regime generally show low selectivities. Therefore, there is interest
in developing microporous membrane materials with pore sizes
<2 nm, exhibiting increased selectivities. In this case, micropore
diffusion is the predominant form of gas transport [3]. A variety of
microporous materials have been under investigation for this
purpose, including SiO2 [4], zeolites [5,6], carbon-based compounds [7,8], metal-organic frameworks [9,10], and titanosilicates
and related materials [11e13]. Common problems limiting the
application of these materials are structural instability, especially in
presence of humid atmospheres in case of SiO2, high costs and
reproducibility problems in case of zeolites, and poor mechanical
properties and sensitivity towards oxidative environments in case
of carbon-based structures [14,15].

/>1387-1811/© 2016 The Author(s). Published by Elsevier Inc. This is an open access article under the CC BY license ( />

T. Konegger et al. / Microporous and Mesoporous Materials 232 (2016) 196e204

The generation of microporous ceramic compounds from preceramic polymers has gained increasing scientific interest in the
recent past. Polymer-derived ceramics (PDCs) are generally obtained by a controlled thermal conversion (pyrolysis) of mainly

organosilicon preceramic polymers into Si-based ceramic materials
[16e18]. Owing to the polymeric nature of the precursors and the
resulting ability to use processing techniques otherwise not
feasible for ceramics, the PDC route represents an alternative to
traditional powder-technology based ceramic processing. Preceramic polymers have been used for obtaining a wide variety of
complex ceramic structures for energy-related applications,
including coatings, fibers, bulk materials, and composites [19].
Scientific interest in polymer-derived cellular ceramics has
demonstrated the potential for obtaining macroporous materials
with uniform, graded, or hierarchical pore systems [20e22].
Microporosity in PDCs is generally produced during the
polymer-to-ceramic conversion of the precursor [23,24]. Both the
evolution of gaseous by-products as a result of the decomposition
of organic groups as well as molecular rearrangements within the
three-dimensional network structure lead to the development of
an intrinsic micropore network [25]. However, further increase in
temperature results in a closing of micropores, most likely due to
viscous flow processes driven by a reduction in surface energy
[26,27]. For instance, microporous compounds derived from preceramic polysiloxanes show a significant reduction in microporosity starting at temperatures as low as 450  C, with an almost
complete collapse of the micropore network at temperatures
beyond 700  C [24].
A significant problem of oxygen-rich Si-based PDCs is the
limited stability at high temperatures; furthermore, Si-O-C bonds
are considered to be hydrothermally unstable [28]. Non-oxide PDCs
such as polysilazane-derived silicon carbonitrides exhibit improved
resistance against oxidation and creep at elevated temperatures
[29]. As a result, the fabrication of these types of microporous nonoxide ceramics for a variety of fields, including catalysis and
membrane applications, has gained increasing attention [23,27,30].
Dismukes and co-workers [23] reported on a variety of methods to
generate microporous polysilazane-derived non-oxide ceramics,

including the controlled thermal decomposition in a reactive atmosphere, in this case, ammonia (NH3), the use of inert, sub-mm
sized fillers, the formation of metal- or cermet-ceramic composites
from polysilazane-stabilized metal colloids, or the thermal conversion of polysilazane mixed with metal-organic compounds.
Using these methods, materials with specific surface areas
>500 m2 gÀ1 and micropore volumes >0.2 cm3 gÀ1 were generated.
A further increase in thermal stability of the micropore structure,
i.e. the shift of the pore collapse onset towards higher temperatures, was reported by chemical modification of the preceramic
polymer, e.g. by the addition of Ni [27,31], or by using a reactive gas
atmosphere such as NH3 during thermal decomposition [30].
Owing to the straightforward method of generating and controlling
their micropore structure, microporous PDC-based SiO2 [32], Si(O)-C [28,33e38], Si-C-N [39e41], and Si-B-C-N [42,43] have been
investigated as selective layer structures in membrane systems.
In addition to the micro- or mesoporous selective layer structure, porous ceramic membranes generally require a macroporous
support structure, responsible for the structural integrity of the
component, resulting in an asymmetric structure with multi-scale
porosity. Often, additional intermediate layers are present to suppress formation of pinhole defects in the selective layer. The most
commonly used material for macroporous supports in PDC-based
gas-separation membranes is Al2O3 [2], but the use of other ceramics such as SiC has also been reported [44,45]. The choice of
mesoporous intermediate layer depends on the materials involved.
g-Al2O3, obtained through sol-gel techniques, is a common material

197

used on a-Al2O3-based supports [46]. An alternative method for the
application of intermediate layers is the combination of preceramic
polymers (in this instance, acting as a binder) with particulate
fillers, e.g. SiC or Si3N4 [28,39]. However, the use of different materials within a single membrane leads to a variety of problems. A
mismatch of coefficients of thermal expansion between support
and deposited layers decreases the ability to withstand repeated
heating and cooling cycles without stress-induced damage.

Furthermore, increased chemical reactivity between the constituents at elevated temperatures places a limit on the maximum
operation temperature of the membrane systems.
Recently, we reported on a novel method for the preparation of
porous ceramic support structures based on preceramic polymers,
including polysilazane [47]. The supports, comprising a welldefined macropore structure obtained through a sacrificial template approach, were found to exhibit tailorable strength and
permeability characteristics [48], rendering them potentially suitable as membrane supports. Based on these findings, we propose a
new strategy for the generation of completely novel structures
derived from a single preceramic polymer, exhibiting pore sizes
ranging from the nanometer- to the micrometer-range.
The specific objectives of this work are the development of
asymmetric ceramic structures with multi-scalar porosity,
composed solely of a polymer precursor-derived non-oxide
ceramic, as well as the proof-of-concept for potential applications
in membrane processes, in particular an evaluation of the selective
permeation of a variety of gaseous species at elevated temperatures, with the aim of obtaining permselectivities exceeding selectivities determined by Knudsen flow. The preceramic polymer
used has to fulfill a variety of requirements, including the potential
for the production of both macro- and microporous ceramics, as
well as adequate properties of the derived material itself, including
sufficient stability at high temperatures in oxidizing and reducing
atmospheres and hydrothermal stability. To fulfill these requirements, we chose a commercially available poly(vinyl)silazane
(PVS) as precursor.
With respect to the objectives of this work, we first present a
methodology for the generation of microporous ceramics with pore
sizes <1 nm, thus being, e.g., in the range of the kinetic diameter of
gases typically encountered in energy production. In the second
part of this contribution, we report on the generation of novel
asymmetric porous structures potentially suitable for membrane
applications, all sub-structures of which are derived from a single
polysilazane-based precursor, including a microporous top layer, a
macro-/mesoporous intermediate layer, as well as a macroporous

support. Finally, we present the proof-of-concept of our singleprecursor approach, demonstrating the general feasibility of these
novel structures for the permeation of gases with selectivities
above values limited by Knudsen diffusion, at temperatures up to
500  C.
By following our novel single-precursor approach, we expect to
circumvent potentially detrimental effects caused by thermal or
chemical incompatibilities between constituents, thus leading to
the development of completely new membrane or nanofiltration
systems suitable for harsh chemical and thermal environments.
2. Experimental
2.1. Preceramic polymer
As precursor material, a commercially available poly(vinyl)silazane (PVS; HTT1800, AZ Electronic Materials, USA) was used. The
pure, liquid polymer has a viscosity of 22 mPa s at 25  C [49], and is
sensitive towards hydrolytic decomposition in ambient atmosphere. As a result, processing was conducted inside a high-purity


198

T. Konegger et al. / Microporous and Mesoporous Materials 232 (2016) 196e204

N2 atmosphere glove box. Dicumyl peroxide (99% purity, Acros
Organics, USA) was dissolved in the preceramic polymer as a radical
initiator to reduce the curing temperature (1 wt% with respect to
the polymer). After addition of the radical initiator, the PVS compound was thoroughly degassed in vacuum before further use. All
pyrolysis treatments of the PVS materials were conducted in an
alumina tube furnace (Lindberg Blue M HTF5534C, Thermo Scientific, USA) in flowing high-purity N2 atmosphere (0.3 l minÀ1; N2 4.8
grade, Airgas, USA), using a gas purification unit to further reduce
H2O and O2 content in the furnace atmosphere (GF-20A, MTI Corp.,
USA).
2.2. Preparation and characterization of microporous materials

Before preparing asymmetric porous structures, the micropore
structure evolution of the PVS compound during the controlled
thermal decomposition in high-purity N2 atmosphere was evaluated. PVS was cross-linked at 105  C for 16 h in high-purity N2 and
subsequently pyrolyzed at temperatures of 400, 500, 600, 700, and
800  C for 4 h, following tailored heating profiles (Fig. 1). Heating
ramps and holding times were chosen with respect to typical parameters required for the preparation of crack-free support specimens. After pyrolytic conversion, the obtained materials were
crushed to powders and stored in inert atmosphere for further
investigations.
The specific surface area and micro-/mesopore size distributions
were assessed by nitrogen adsorption at 77 K (ASAP 2010, Micromeritics Instrument Corporation, USA). The powdered samples
(one sample per pyrolytic conversion temperature) were evacuated
and degassed at 250  Ce300  C for 12 h before adsorption testing.
Specific surface areas were determined by the Brunauer-EmmettTeller method (BET). The total pore volume was calculated from
the amount of N2 adsorbed at a relative pressure of p/p0 ¼ 0.995.
The micropore size distribution was obtained from the adsorption
branch of the nitrogen adsorption/desorption isotherms recorded
in the relative pressure range p/p0 of 10À5 to 0.995. The non-local
density functional theory method (NLDFT), using the 2D-NLDFT
model for heterogeneous surfaces included in the SAIEUS software
package (Micromeritics Instrument Corporation, USA) was used to
calculate the micropore size distribution. From NLDFT results, the
micro- and mesopore volumes were also calculated.
As complementary technique, small-angle X-ray scattering
(SAXS) was used to obtain and validate the micropore structure of

Fig. 1. Heat treatment profiles used for the generation of PVS-derived materials (“400”
through “800” for evaluation of micropore evolution, “800” for support structures and
intermediate layers, “600” for selective layers).

each sample. SAXS was performed with Cu Ka radiation generated

by a microfocus source (Incoatec High Brilliance, Germany) using a
setup with a pinhole camera and an area detector (Nanostar and
Våntec-2000, Bruker AXS, Germany). The samples were attached
between two stripes of a Scotch tape and measured in vacuum at
two different sample to detector distances (108 cm and 13 cm) to
cover a wide range of the q-vector from 0.1 to 20 nmÀ1. The measurement time was 1500 and 900 s for the respective distance. The
data were then radially averaged to obtain the scattering intensity
depending on the scattering vector q ¼ 4p/l sin(q) with 2q being
the scattering angle and l ¼ 0.1542 nm the X-ray wave length.
Finally, the data were merged in the overlap region and corrected
for background scattering. The unified equation for the scattering
intensity as described by Beaucage [50] was used, which combines
Guinier’s and Porod’s law. An additional interference factor had to
be introduced to describe the interaction of the pores within the
pore network: For this, a structure factor from a hard sphere model
[51,52] was chosen. Furthermore, a qÀ4 power law arising from
large objects or grains within the sample was added.
2.3. Preparation of asymmetric polymer-derived ceramic
membranes
The planar, macroporous support structures were prepared
through a sacrificial filler approach recently reported [47,48]. PVS
was mixed with ultra-high molecular weight polyethylene
(UHMW-PE) particulates (Mipelon PM-200, Mitsui Chemicals
America, USA; d50 ¼ 10 mm) at a volume fraction of 30 vol%, cast into
polydimethyl siloxane (PDMS) molds, and cross-linked in N2 atmosphere at 105  C for 16 h. After demolding and an additional
post-curing step, the surface of the cylindrical specimens was
ground to a 2000 grit finish, yielding planar discs with a diameter of
18 mm and a height of 3e4 mm. The cured specimens were pyrolytically converted into ceramics in high-purity flowing N2 atmosphere at 800  C for 4 h following a tailored heating profile (Fig. 1,
“800”), thereby thermally decomposing the sacrificial filler compound as well as converting the preceramic PVS into a ceramic. A
more detailed description of this technique for the preparation of

free-standing polysilazane-derived support structures is described
elsewhere [47].
The intermediate layer was applied via dip-coating. The coating
slurry consisted of a mixture of uncured PVS (8 wt%), pyrolyzed
PVS-derived particulates (32 wt%), and toluene (60 wt%; anhydrous, 99.8% purity, Alfa Aesar, USA). The pyrolyzed particulates
were obtained through pyrolytic conversion of the preceramic
polymer in high-purity nitrogen at 800  C for 4 h. Milling in
cyclohexane with Si3N4 milling media (8000M Mixer/Mill, Spex
SamplePrep, USA) yielded a final average particle size of
0.6e0.7 mm, measured by laser scattering particle size analysis (LA910, Horiba, USA). The coating slurry was homogenized by milling
and ultrasonication. Before coating, the macroporous supports
were cleaned by ultrasonication in cyclohexane and dried at 60  C.
Surfaces not to be coated were masked with PTFE tape. Dip-coating
was conducted at a withdrawal speed of 100 mm minÀ1 after an
immersion for 1 s. After evaporation of the solvent, removal of the
masking tape, and wiping off excess material from the specimen
surface, the specimens were heat-treated at 800  C, following the
standard heating profile (Fig. 1, “800”).
Subsequently, the selective layer was deposited on top of the
intermediate layer by dip-coating. PVS was dissolved in toluene
(anhydrous, 99.8% purity, Alfa Aesar, USA), yielding the coating
solution. After masking of the specimens’ side and bottom with
PTFE tape, samples were immersed in the coating solution (1 s) and
withdrawn at a speed of 20 mm minÀ1. The coated samples were
heat-treated at 600  C, following the heating schedule as listed


T. Konegger et al. / Microporous and Mesoporous Materials 232 (2016) 196e204

above (Fig. 1). A series of coating/pyrolysis steps was conducted to

minimize the number of defects and pinholes in the membranes,
starting with a coating solution containing 50 vol% PVS in toluene
(SL-50), and using a concentration of 5 vol% for subsequent coatings. A maximum number of four coatings were deposited (SL-505-5-5). Two specimens were coated with a fourth layer using a
coating solution containing 20 vol% PVS instead (SL-50-5-5-20).
The microstructure of specimens at intermediate stages as well
as of final membranes was evaluated by scanning electron microscopy (SEM; S4800, Hitachi, Japan).
2.4. Evaluation of permeance properties
A screening of the gas permeance characteristics of the prepared
planar structures was conducted using a custom-built single gas
constant-volume variable-pressure setup (Fig. 2) at temperatures
up to 500  C, using one sample per processing variation. Reliability
of the permeance characteristics was checked by testing of an
additional, independently fabricated specimen of SL-50-5-5-20. He,
N2, Ar, C2H6, and CO2 were used as test gases (HP grades, Airgas,
USA). The permeation area of the planar specimens tested was
about 0.8 cm2. Before testing and between gas changes, the system
was evacuated overnight. After a leak test, the feed side of the
membranes was exposed to a constant gas flow of 0.1 L minÀ1 at
ambient pressure while the permeate side was kept evacuated.
After reaching stationary pressure values, the cut-off valve to the
vacuum system was closed, and the slope of the pressure increase
on the permeate side was recorded in the range of 133e666 Pa
(1e5 torr) using data processing software (LabView environment).
The pressure rise was used to calculate the permeance Pi for each
gas i (mol mÀ2 sÀ2 PaÀ1). The permselectivities, i.e., the ideal separation factors, were obtained from the permeance ratios of two
given gases, e.g., PHe/PAr. A thorough description of the testing
procedure and data evaluation, including the corrections applied,
can be found in the Supporting Information.
3. Results and discussion
3.1. Processing and tailoring of the micropore structure

The preparation of an asymmetric polymer-derived ceramic
membrane requires the ability to generate a microporous selective
layer, on a macroporous support. Processing and tailoring of the
micropore structure of the PVS-derived materials are critically
important as both aspects affect the performance of the material. In
this work, the polymer-to-ceramic conversion process itself is used
to generate microporous PVS-derived materials. After thermal

199

treatment of cross-linked PVS at temperatures between 400  C and
800  C in flowing N2, a strong increase in the specific surface area
(SSA) of PVS-derived compounds was found, with a maximum SSA
of 440 m2 gÀ1 after pyrolysis at 500  C (Fig. 3a). Further increase in
temperature resulted in a precipitous reduction in surface area,
reaching values < 5 m2 gÀ1 after treatment above 700  C. In this
case, the material can be considered unsuitable for membrane
applications.
Since the pyrolysis temperatures must exceed the intended
operation temperatures in order to avoid ongoing conversion processes in the material, samples pyrolyzed at 400  C are potentially
unstable at temperatures above 250e300  C. Therefore, PVSderived compounds after conversion at 500  C and 600  C can be
anticipated to be most promising for a potential use as membrane
materials at elevated temperatures. For these two materials, the
micropore structure was investigated in-depth using N2 adsorption.
In both cases, the adsorption/desorption isotherms can be identified as Type I according to IUPAC classification [53], indicating
microporous compounds (Fig. 3b).
Using a non-localized density functional theory (NLDFT) model,
the pore size distribution was calculated from the adsorption
branches of the recorded adsorption/desorption isotherms (Fig. 3c).
As expected, the major pore volume consists of micropores with

pore diameters well below 2 nm. Even though the total micropore
volume is significantly reduced by increasing the conversion temperature to 600  C, the shape of the pore size distribution remains
comparable, with maxima of the differential PSD curves in the
range of 0.8e0.9 nm. This finding indicates the increasing limitation of access to pores smaller than 3 nm with increasing process
temperature, without significantly changing the structure and size
of pores themselves in the temperature range investigated. The
total pore volume of 0.25 cm3 gÀ1 after conversion at 500  C decreases to 0.10 cm3 gÀ1 after conversion at 600  C, the total pore
volume primarily consisting of microporosity (Fig. 3d).
Fig. 4 shows the experimental SAXS data (symbols) together
with the fits from the model (lines). The main parameters arising
from the model are measures for the pore size (radius of gyration
Rg), the fractal dimension D (describing the density of the network),
the hard sphere radius (twice the hard sphere radius is the typical
distance of the pores), and the hard sphere volume factor h (the
packing density of the pores). The numerical fit values are found in
Table 1.
The mean pore size, described by the radius of gyration, increases from 400  C to 500  C and is then constant. The typical
interaction distance of the pores increases slightly from 0.87 nm at
400  C to 1.3 nm at 800  C. At the same time the hard sphere
volume fraction h decreases from 0.3 to 0.14. The fractal dimension

Fig. 2. Schematic of the experimental setup for single-gas permeance testing of planar asymmetric membranes at temperatures up to 500  C, using the constant-volume variablepressure technique.


200

T. Konegger et al. / Microporous and Mesoporous Materials 232 (2016) 196e204

Fig. 3. Results of N2 adsorption experiments, showing (a) the specific surface area of PVS-derived materials as a function of pyrolysis temperature, (b) N2 adsorption/desorption
isotherms of two promising candidate materials (PVS pyrolyzed at 500 and 600  C), (c) subsequent determination of the micro-/mesopore size distribution by NLDFT, and (d) the

amount of micro- and mesoporosity with respect to the total pore volume of the samples investigated.

this temperature. From the tendency of the fit parameters in Table 1
one can conclude that the pore network is not very sensitive to the
temperature treatment in terms of changes in pore size: The overall
size of pores remains roughly the same, only the decreasing hard
sphere volume fraction (decreasing packing density) and the
increasing hard sphere radius indicate a slightly more homogeneous distribution with less interaction of the pores at higher
temperatures. This lower interaction could influence the connectivity of the network and supports the observation of NLDFT, i.e. the
shape of the pore size distribution is not affected, but after higher
heat treatment temperatures of the material the access to the pores
is reduced. This, in turn, leads to sharply decreased SSA values and,
as a consequence, places an intrinsic temperature limit on the
applicability of the microporous materials.
As demonstrated by N2 adsorption studies and SAXS investigations, microporous PVS-derived materials were successfully
prepared, using intrinsic gas evolution and molecular rearrangement processes during the thermal conversion process.
Fig. 4. Experimental SAXS profiles (symbols) and fitting curves (lines) of PVS-derived
materials as a function of pyrolysis temperature. The insert shows the data in linear
scales for a selected q-range.

D is rather constant close to 4, which corresponds to a smooth
surface in a two-phase model, and decreases to 3.2 only at 800  C,
which is an indication for the onset of an increased roughness at

Table 1
Obtained parameters from the fit model to the SAXS data: Radius of gyration Rg, hard
sphere radius Rhs, fractal dimension D, and hard sphere volume fraction h.
Temperature/ C

2Rg/nm


D

2Rhs/nm

h

400
500
600
700
800

0.55
0.79
0.82
0.76
0.75

4
4
3.8
3.7
3.2

0.87
0.87
0.93
1.1
1.3


0.3
0.25
0.13
0.12
0.14

3.2. Generation of asymmetric porous structures
A combination of shaping and coating techniques was used to
prepare asymmetric ceramic membrane structures in planar
configuration derived from a single preceramic polymer. In previous reports, we described the development of novel, free-standing
open-porosity polymer-derived ceramic structures with potential
use for membrane and catalysis applications [47,48]. In this present
work, these structures were used as supports in asymmetric PVSderived ceramic membranes. The use of UHMW-PE fillers with a
particle size of 10 mm, incorporated into the preceramic polymer
before subsequent casting, cross-linking, and pyrolysis, allowed for
a tailoring of the pore structure of the final ceramic material. After
pyrolytic conversion, planar supports with a diameter of 13.5 mm
and a height of 2.5 mm, exhibiting an interconnected open pore
network, were obtained (Fig. 5). Preliminary permeance tests of the


T. Konegger et al. / Microporous and Mesoporous Materials 232 (2016) 196e204

Fig. 5. Photograph (a) and cross-sectional SEM image (b) of planar PVS-derived
ceramic membrane supports.

planar supports showed helium permeances exceeding
10À5 mol mÀ2 PaÀ1 sÀ1 at room temperature, thus effectively surpassing the upper limits of the recording speed of the pressure rise
rate determined by the permeance test setup. These values are in

general accordance with air permeability characteristics reported
previously for these materials [47].
The deposition of defect-free layers on top of the porous support
structures proved to be a major challenge during the preparation of
the asymmetric PVS-derived membrane structures. Dip-coating of
the untreated supports with a polymer precursor solution led to an

201

infiltration of the base structure, thus yielding inadequate results.
The deposition of an intermediate layer was found to be essential
for the successful application of the selective layer. The intermediate layer serves the role of planarizing the support layer. A slurry
consisting of uncured PVS in combination with pyrolyzed PVSderived powders dispersed in toluene was applied on the surface
of the support structure. By immediately removing excess slurry
from the support surface (using lab wipes) before cross-linking and
thermal conversion of the intermediate layer, the surface pores
could be completely filled with the coating slurry. This resulted in a
smooth surface with an effective reduction in surface pore size
from around 7 mm down to the sub-mm size range after pyrolytic
conversion (Fig. 6a). This smooth surface allowed for the subsequent deposition of the selective layer.
As discussed in the previous section, the highest amount of
microporosity in PVS-derived materials was achieved after a pyrolytic conversion at 500  C. However, since gas permeances were
investigated at temperatures up to 500  C (see Section 3.3), pyrolysis of selective layers was conducted at 600  C to avoid
temperature-induced structural changes during permeance testing.
Even though the total amount of microporosity is reduced after
pyrolysis at 600  C, the size of micropores remains the same, according to adsorption and SAXS results presented in Figs. 3 and 4. A
sequence of dip-coating/pyrolysis cycles, using coating solutions of
various concentrations of PVS in toluene, led to continuous selective layer structures on top of the planar supports. While a single
coating was found insufficient for the production of a dense
structure due to residual infiltration of the coating solution into the

intermediate layer (Fig. 6b), subsequent coatings were found to
gradually increase the thickness of the top layer, reaching approximately 5 mm after the fourth coating (Fig. 6c and d). Even though
pyrolytic conversion typically goes along with significant shrinkage
of the material, selective layers did not exhibit shrinkage-caused
cracking, which can be attributed to the porous structure of the
intermediate layer and the mechanical interlocking through partial
infiltration by the selective layer, in combination with low layer
thicknesses.
Due to both the increasing layer thickness and a reduction in the
number of critical defects such as cracks and pinholes in the top

Fig. 6. Cross-sectional SEM images of PVS-derived ceramic layer structures with multiscale porosity after application and thermal conversion of (a) intermediate layer (S ỵ I), (b)
one selective layer (SL-50), (c) four selective layers (SL-50-5-5-5), and (d) four selective layers using a more concentrated final coating solution (SL-50-5-5-20).


202

T. Konegger et al. / Microporous and Mesoporous Materials 232 (2016) 196e204

layer, He and N2 permeances of the membrane specimens progressively decreased with an increasing number of coating cycles
(Fig. 7). After the fourth coating/pyrolysis sequence, permeances
were around two orders of magnitude below the permeances
before application of selective layers (S ỵ I). The use of a more
concentrated coating solution for the final coating step resulted in
even lower permeances; however, the ratio of permeances PHe/PN2
was higher than in previous steps (S ỵ I: 1.8; SL-50: 2.1, SL-50-5-5-5:
2.2; SL-50-5-5-20: 3.1), indicating improved selectivity.

the selective layer structures are the most probable cause for this
behavior. Viscous flow through defects such as cracks and pinholes

hence cannot be ruled out, as the test setup did not allow for variations in differential pressure between gas feed and permeate
sides to clearly identify a non-linear relationship between flow and
pressure drop characteristic for a viscous flow regime.
By increasing the concentration of PVS in the coating liquid for
the final coating step during selective layer preparation (SL-50-5-5-

3.3. Gas permeance characteristics
A screening of gas permeance characteristics of asymmetric
PVS-derived ceramic membranes was conducted at temperatures
up to 500  C using a variety of gases with distinct kinetic diameters
(He: 0.26 nm, N2: 0.364 nm, Ar: 0.34 nm, C2H6: 0.444 nm, and CO2:
0.33 nm [54,55]) in order to determine and evaluate the general
feasibility of the proposed processing approach for membranebased separation processes. A summary of results is shown in
Fig. 8 and Table 2, with detailed listings of permeances and
permselectivities in the Supplementary Information. For screening
purposes, permeance data were initially obtained for single samples for each processing variation. Subsequently, for the more
promising coating sequence (SL-50-5-5-20), a second, independently prepared specimen was tested to evaluate the general
reproducibility of the permeance behavior. As a consequence, the
results can be considered as indicators for the general suitability of
the structures for membrane-based separation applications, in
accordance with the aims of this work.
A comparison between membranes showed significant differences depending on the coating sequence (Fig. 8). After the initial
coating procedure (SL-50-5-5-5), a steady decrease in gas permeance with increasing temperature can be observed, independent
of the gas investigated. This, together with a linear relationship
between permeances and the inverse square root of molecular
weights of gases (Fig. 8b), indicates that Knudsen diffusion is the
dominating transport mechanism. With rising temperature,
permselectivities approach Knudsen values of the corresponding
gas pairs (Table 2), which are equal to the ratio of the inverse square
root of molecular weights of gas pairs investigated [1]. Since

Knudsen flow is one of the predominant transport mechanisms in
macro- and mesoporous materials, and taking into account the
microporous nature of the selective layer material itself, defects in

Fig. 7. Helium and nitrogen permeances of membrane specimens at various stages
during the generation of asymmetric pore structures. Permeance data is shown for (a)
support with intermediate layer before deposition of selective layer (S ỵ I), as well as
for specimens after deposition of (b) one selective layer (SL-50), (c) four selective layers
(SL-50-5-5-5), and (d) four selective layers using a more concentrated coating solution
(SL-50-5-5-20). Permeances were determined for specimens of comparable thickness
at room temperature.

Fig. 8. Gas permeance characteristics of asymmetric PVS-derived ceramic membranes
SL-50-5-5-5 (a,b) and SL-50-5-5-20, Test 1 (c,d), permeances being shown as a function
of test gas, temperature (a,c), and the inverse square root of the molecular weight of
gases (b,d). The initial coating procedure (SL-50-5-5-5) during selective layer preparation resulted in permeance characteristics following the Knudsen diffusion regime.
While reducing the overall permeance, the use of a coating solution with a higher PVS
concentration for the final coating step (SL-50-5-5-20) resulted in permeance characteristics slightly deviating from Knudsen diffusion, with permselectivities higher
than anticipated by pure Knudsen flow.


T. Konegger et al. / Microporous and Mesoporous Materials 232 (2016) 196e204

203

Table 2
Helium permselectivities (aHe/X) of SL-50-5-5-5 and SL-50-5-5-20 membranes.
Membrane

Temperature/ C


He/N2

He/Ar

He/C2H6

He/CO2

SL-50-5-5-5

20
300
500
20
300
500
20
300
500

2.15
2.37
2.49
3.15
2.96
2.97
3.00
3.13
3.11

2.65

2.39
2.79
2.94
3.42
3.47
3.44
3.44
3.75
3.61
3.16

1.92
2.39
2.52
3.30
4.14
3.67
2.53
3.44
3.34
2.74

2.24
2.86
3.06
4.19
4.44
4.10

2.39
3.63
3.73
3.32

SL-50-5-5-20

Test 1

Test 2

Theoretical Knudsen permselectivities:

20), permselectivities of all gas pairs were increased, exceeding
Knudsen values over a wide temperature range. The apparent
reduction of the number of defects on the membrane leads to a
decrease in overall permeance values. Permselectivities of N2 and
Ar with respect to He were found to remain relatively stable over
the full temperature range. In contrast, C2H6 and CO2 show a slight
deviation in permeance characteristics, both showing an increase in
permselectivity with respect to He in the lower temperature range.
A potential explanation for the deviating behavior of C2H6 and CO2
permeances is a combination of Knudsen flow with other effects
such as adsorption and surface diffusion for C2H6 and CO2 at temperatures below 200  C, resulting in increased permeances and
hence decreased permselectivities with respect to He (and other
non-interacting gases such as N2 and Ar) at these temperatures, an
effect which has been reported, e.g., for silica-based membranes
[56]. At higher temperatures, these effects diminish. Permeance
and selectivity characteristics were found to be comparable between independently prepared samples of SL-50-5-5-20, especially
for non-interacting gases such as He, N2, and Ar (Table 2).

Observable differences, in particular in the permeance behavior of
C2H6 and CO2 at lower temperatures, may be a result of small
variations in processing conditions, leading to variations in the
chemical setup of pore surfaces and hence to an altered impact of
adsorption and surface diffusion effects.
As described in Section 3.1, the evaluation of the micropore
structure of PVS-derived materials was carried out on powdered
samples, as a direct determination of deposited layers was not
feasible. As such, differences in the micropore structure between
powdered and layered PVS-derived structures cannot be ruled out
completely, even though the processing sequence was kept similar
wherever possible. However, the determination of selectivities
exceeding Knudsen selectivities indicate that microporosity is
indeed present in deposited layers.
These results demonstrate that it is possible to generate asymmetric ceramic structures from a single Si-based precursor material
exhibiting gas permeation at elevated temperatures with permselectivities exceeding Knudsen selectivities. Even though the permeances and permselectivities currently achieved are well below
performances of state-of-the-art inorganic gas separation membranes in combination with conventional support structures, with
ideal selectivities He/CO2 of around 40 at He permeances of
10À7 mol mÀ2 PaÀ1 sÀ1 reported for polymer-derived silica systems
[46], or even higher H2/CO2 ideal selectivities exceeding 100 at H2
permeances of up to 10À6 mol mÀ2 PaÀ1 sÀ1 in case of conventional
microporous zeolite- or silica-based systems [2], our approach is a
first step towards the potential establishment of novel PDC-based
membrane systems, with wide possibilities for future improvements. By further optimizing the morphology of both support and
intermediate layer, as well as parameter optimization during the
deposition of the selective layer structure, an improved performance in terms of permeance and selectivity can be expected. In

the methodology reported here, microporosity within the selective
layer was generated solely by pyrolytic conversion of pure PVS in
inert atmosphere. Recent reports of significant improvements in

hydrothermal and thermal stability of microporous polysilazanederived ceramic materials, both by chemical modification of preceramic compounds [27,31] as well as by decomposition of Si-based
precursors in reactive atmospheres, e.g., ammonia [30], represent
promising tools to facilitate the development of single-material
ceramic membrane systems for more demanding operating environments in the future.
4. Conclusions
The generation of asymmetric non-oxide ceramic structures
with multi-scalar porosity from a single-source preceramic polymer has potential benefits for applications in membrane systems,
owing to both thermal and chemical compatibility between the
components comprising the micro- and macropore structure, and
is, to our best knowledge, a completely novel concept. A variety of
methods were used to create asymmetric ceramic structures consisting of a poly(vinyl)silzane-derived Si-C-N-based material in
planar-disc configuration. The structures generated included a
macroporous support with mm-sized pores obtained through a
sacrificial filler approach, an intermediate structure consisting of
precursor-bonded particulates, as well as selective layer structures
with micropores smaller than 1 nm, obtained through a controlled
thermal decomposition of the polymer precursor. While the
micropore size was shown to be relatively stable up to 800  C, the
interconnectivity of the pore network was found to collapse at
temperatures above 600  C, thus effectively limiting the operating
temperature of the membrane structures.
By evaluating single gas permeances at elevated temperatures
up to 500  C, we showed that the generation of asymmetric nonoxide ceramic membranes from of a single Si-based polymer precursor is feasible, and a permeation of gaseous species with selectivities above ideal selectivities solely governed by Knudsen flow
can be achieved. Even though there is significant room for
improvement of the performance in terms of permeances and
selectivity, the work presented here can be considered as a first step
towards the establishment of new membrane materials systems for
extreme operating conditions based on the wide variety of preceramic polymers available, either commercially available or
obtainable by targeted synthesis, with potential applicability in the
fields of gas separation or nanofiltration.

Acknowledgments
T.K. gratefully acknowledges funding by the Austrian Science
Fund (FWF): J3422-N28. RKB gratefully acknowledges partial support for this work from the US National Science Foundation under
grant No. DMR MWN 1008600. Furthermore, the authors thank L.F.


204

T. Konegger et al. / Microporous and Mesoporous Materials 232 (2016) 196e204

Williams and M. Greenough for support with permeance measurements, J. Shetzline for support with adsorption experiments, D.
Field and Dr. C.E. Sosolik for assistance with leak testing of the
permeance setup, and Dr. H.J. Rack for discussions regarding test
cell construction. SEM investigations were conducted at the Electron Microscopy Laboratory at Clemson University.
Appendix A. Supplementary data
Supplementary data related to this article can be found at http://
dx.doi.org/10.1016/j.micromeso.2016.06.027.
References
[1] W.J. Koros, G.K. Fleming, J. Membr. Sci. 83 (1993) 1e80.
[2] H. Verweij, Y.S. Lin, J. Dong, MRS Bull. 31 (2006) 756e764.
[3] K. Li, Ceramic Membranes for Separation and Reaction, Wiley, Chichester,
2007.
[4] R.M. de Vos, H. Verweij, Science 279 (1998) 1710e1711.
[5] J. Caro, M. Noack, Microporous Mesoporous Mater. 115 (2008) 215e233.
[6] D. Fedosov, A. Smirnov, E. Knyazeva, I. Ivanova, Pet. Chem. 51 (2011)
657e667.
[7] E.N. Hoffman, G. Yushin, B.G. Wendler, M.W. Barsoum, Y. Gogotsi, Mater.
Chem. Phys. 112 (2008) 587e591.
[8] Y.J. Fu, K.S. Liao, C.C. Hu, K.R. Lee, J.Y. Lai, Microporous Mesoporous Mater. 143
(2011) 78e86.

[9] J.-R. Li, R.J. Kuppler, H.-C. Zhou, Chem. Soc. Rev. 38 (2009) 1477e1504.
[10] R. Adams, C. Carson, J. Ward, R. Tannenbaum, W. Koros, Microporous Mesoporous Mater. 131 (2010) 13e20.
[11] V. Sebastian, Z. Lin, J. Rocha, C. Tellez, J. Santamaria, J. Coronas, Chem. Commun. (2005) 3036e3037.
llez, J. Coronas, Sep. Purif. Technol.
[12] I. Tiscornia, I. Kumakiri, R. Bredesen, C. Te
73 (2010) 8e12.
[13] X. Li, C. Zhou, Z. Lin, J. Rocha, P.F. Lito, A.S. Santiago, C.M. Silva, Microporous
Mesoporous Mater. 137 (2011) 43e48.
[14] P. Bernardo, E. Drioli, G. Golemme, Ind. Eng. Chem. Res. 48 (2009) 4638e4663.
[15] G.Q. Lu, J.C. Diniz da Costa, M. Duke, S. Giessler, R. Socolow, R.H. Williams,
T. Kreutz, J. Colloid Interface Sci. 314 (2007) 589e603.
[16] R. Riedel, G. Mera, R. Hauser, A. Klonczynski, J. Ceram. Soc. Jpn. 114 (2006)
425e444.
[17] P. Colombo, G. Mera, R. Riedel, G.D. Sorarù, J. Am. Ceram. Soc. 93 (2010)
1805e1837.
[18] P. Colombo, R. Riedel, G.D. Sorarù, H.-J. Kleebe, Polymer Derived Ceramics:
from Nano-structure to Applications, DEStech Publications, Inc., Lancaster,
2010.
[19] T. Konegger, J. Torrey, O. Flores, T. Fey, B. Ceron-Nicolat, G. Motz, F. Scheffler,
M. Scheffler, P. Greil, R.K. Bordia, in: A.K. Agarwal, A. Pandey, A.K. Gupta,
S.K. Aggarwal, A. Kushari (Eds.), Novel Combustion Concepts for Sustainable
Energy Development, Springer, India, 2014, pp. 501e533.
[20] J. Zeschky, F. Goetz-Neunhoeffer, J. Neubauer, S.H. Jason Lo, B. Kummer,
M. Scheffler, P. Greil, Compos. Sci. Technol. 63 (2003) 2361e2370.
[21] P. Colombo, J. Eur. Ceram. Soc. 28 (2008) 1389e1395.
[22] P. Colombo, C. Vakifahmetoglu, S. Costacurta, J. Mater. Sci. 45 (2010)
5425e5455.

[23] J.P. Dismukes, J.W. Johnson, J.S. Bradley, J.M. Millar, Chem. Mater. 9 (1997)
699e706.

[24] H. Schmidt, D. Koch, G. Grathwohl, P. Colombo, J. Am. Ceram. Soc. 84 (2001)
2252e2255.
[25] M. Wilhelm, C. Soltmann, D. Koch, G. Grathwohl, J. Eur. Ceram. Soc. 25 (2005)
271e276.
nez, J.A. Downs, R. Raj, J. Am. Ceram. Soc. 93 (2010)
[26] P.E. S
anchez-Jime
2567e2570.
[27] M.S. Bazarjani, H.-J. Kleebe, M.M. Müller, C. Fasel, M. Baghaie Yazdi, A. Gurlo,
R. Riedel, Chem. Mater. 23 (2011) 4112e4123.
[28] B. Elyassi, M. Sahimi, T.T. Tsotsis, J. Membr. Sci. 316 (2008) 73e79.
[29] E. Kroke, Y.-L. Li, C. Konetschny, E. Lecomte, C. Fasel, R. Riedel, Mater. Sci. Eng.
R. 26 (2000) 97e199.
[30] C. Schitco, M.S. Bazarjani, R. Riedel, A. Gurlo, J. Mater. Chem. A 3 (2015)
805e818.
[31] M.S. Bazarjani, M.M. Müller, H.-J. Kleebe, Y. Jüttke, I. Voigt, M. Baghaie Yazdi,
L. Alff, R. Riedel, A. Gurlo, ACS Appl. Mater. Interfaces 6 (2014) 12270e12278.
[32] Y. Iwamoto, K. Sato, T. Kato, T. Inada, Y. Kubo, J. Eur. Ceram. Soc. 25 (2005)
257e264.
[33] K. Kusakabe, Z. Yan Li, H. Maeda, S. Morooka, J. Membr. Sci. 103 (1995)
175e180.
[34] L.-L. Lee, D.-S. Tsai, J. Am. Ceram. Soc. 82 (1999) 2796e2800.
[35] H.M. Williams, E.A. Dawson, P.A. Barnes, B. Rand, R.M.D. Brydson, A.R. Brough,
J. Mater. Chem. 12 (2002) 3754e3760.
[36] H. Suda, H. Yamauchi, Y. Uchimaru, I. Fujiwara, K. Haraya, Desalination 193
(2006) 252e255.
[37] R.A. Wach, M. Sugimoto, M. Yoshikawa, J. Am. Ceram. Soc. 90 (2007) 275e278.
[38] B. Elyassi, M. Sahimi, T.T. Tsotsis, J. Membr. Sci. 288 (2007) 290e297.
[39] H. Mori, S. Mase, N. Yoshimura, T. Hotta, K. Ayama, J.I. Tsubaki, J. Membr. Sci.
147 (1998) 23e33.

€lger, R. Hauser, E. Kroke, R. Riedel, Y.H. Ikuhara, Y. Iwamoto, J. Ceram.
[40] K.W. Vo
Soc. Jpn. 114 (2006) 567e570.
[41] Y. Jüttke, H. Richter, I. Voigt, R.M. Prasad, M.S. Bazarjani, A. Gurlo, R. Riedel,
Chem. Eng. Trans. 32 (2013) 1891e1896.
[42] R. Hauser, S. Nahar-Borchard, R. Riedel, Y.H. Ikuhara, Y. Iwamoto, J. Ceram.
Soc. Jpn. 114 (2006) 524e528.
[43] R.M. Prasad, Y. Iwamoto, R. Riedel, A. Gurlo, Adv. Eng. Mater. 12 (2010)
522e528.
[44] W. Deng, X. Yu, M. Sahimi, T.T. Tsotsis, J. Membr. Sci. 451 (2014) 192e204.
[45] Y. Zhou, M. Fukushima, H. Miyazaki, Y.-i. Yoshizawa, K. Hirao, Y. Iwamoto,
K. Sato, J. Membr. Sci. 369 (2011) 112e118.
[46] Y. Iwamoto, J. Ceram. Soc. Jpn. 115 (2007) 947e954.
[47] T. Konegger, R. Patidar, R.K. Bordia, J. Eur. Ceram. Soc. 35 (2015) 2679e2683.
[48] T. Konegger, L.F. Williams, R.K. Bordia, J. Am. Ceram. Soc. 98 (2015)
3047e3053.
[49] T. Konegger, C.-C. Tsai, R.K. Bordia, Mater. Sci. Forum 825e826 (2015)
645e652.
[50] G. Beaucage, J. Appl. Crystallogr. 28 (1995) 717e728.
[51] D.J. Kinning, E.L. Thomas, Macromolecules 17 (1984) 1712e1718.
[52] J.S. Pedersen, Adv. Colloid Interface Sci. 70 (1997) 171e210.
[53] K.S. Sing, Pure Appl. Chem. 57 (1985) 603e619.
[54] D.W. Breck, Zeolite Molecular Sieves: Structure, Chemistry, and Use, Wiley,
New York, 1974.
[55] S. Sircar, A.L. Myers, in: S.M. Anesbach, K.A. Carrado, P.K. Dutta (Eds.), Handbook of Zeolite Science and Technology, Marcel Dekker Inc., New York, 2003,
pp. 1063e1104.
[56] D. Lee, S.T. Oyama, J. Membr. Sci. 210 (2002) 291e306.