Metal–oxide films with magnetically-modulated
nanoporous architectures
Craig A. Grimesa) and R. Suresh Singh
Department of Electrical Engineering & Materials Research Institute, 204 Materials Research
Laboratory, The Pennsylvania State University, University Park, Pennsylvania 16802
Elizabeth C. Dickey and Oomman K. Varghese
Department of Materials Science and Engineering, The Pennsylvania State University,
University Park, Pennsylvania 16802
(Received 13 November 2000; accepted 22 March 2001)
A magnetically-driven method for controlling nanodimensional porosity in
sol-gel-derived metal–oxide films, including TiO2, Al2O3, and SnO2, coated onto
ferromagnetic amorphous substrates, such as the magnetically-soft Metglas1 alloys, is
described. On the basis of the porous structures observed dependence on external
magnetic field, a model is suggested to explain the phenomena. Under well-defined
conditions it appears that the sol particles coming out of solution, and undergoing
Brownian motion, follow the magnetic field lines oriented perpendicularly to the
substrate surface associated with the magnetic domain walls of the substrate; hence the
porosity developed during solvent evaporation correlates with the magnetic domain size.
I. INTRODUCTION
Metal–oxide films of controlled nanoporosity offer an
exciting opportunity for developing a new class of materials with unique physical, electrical, and magnetic
properties. In recent years porous metal oxide films
have attracted considerable attention for application in
photovoltaic cells,2,3 catalysis,4 –9 gas sensors,10 –12 biotemplates,13–16 and electrochromic displays.17 Sol-gel
self-assembly deposition methods are of great interest
due to their inherent flexibility and low cost. Considerable effort has focused on the ordered alignment of a
template around which the material of interest is assembled. Depending upon the desired pore size, organic polymers,18,19 block copolymers,20,21 latex22–24 and
polystyrene25 spheres, water-in-oil emulsions,26 and water droplets27 have been used as templates. A templatefree method for fabrication of micrometer-sized
honeycomb structures by self-assembly of block copolymers has been described in Refs. 28–30.
We have recently discovered that within certain processing windows TiO2, Al2O3, and SnO2 metal–oxide
films dip-coated onto magnetically-soft ferromagnetic
substrates exhibit unique nanodimensional porous architectures; see Fig. 1. We have not been able to replicate
this porous structure on nonferromagnetic substrates,
prompting us to consider the influence of a magnetic
field on the charged sol particles. The thick-film Metglas
a)
e-mail: cgrimes@engr.psu.edu
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J. Mater. Res., Vol. 16, No. 6, Jun 2001
FIG. 1. SEM image of sol-gel-derived 400-nm-thick TiO2 nanoporous
coating made while exposed to a 0.0-Oe static magnetic field. The
average pore size is approximately 60 nm.
2826MB1 substrate used in this work, Fe40Ni38Mo4B18,
is a magnetically soft alloy made by rapid meltquenching, and while it maintains no long-range order,
it does have short-range order over several atom
lengths;31–34 hence, the as-cast material has magnetic
ordering on the scale of tens of nanometers. It appears
that as the sol layer upon the 2826MB substrate evaporates the moving charged sol particle comes out of solution and, under well-defined conditions (e.g., sol
concentration, pH, drying rate), follows the Lorentz force
© 2001 Materials Research Society
C.A. Grimes et al.: Metal – oxide films with magnetically-modulated nanoporous architectures
arising from the magnetic field lines associated with the
perpendicularly oriented (normal to the surface) magnetization vector of the domain walls. This force is a function of the stray magnetic field strength associated with
the vertically oriented magnetization vector of the domain wall (a function of alloy composition and thermal
processing history), the relative strength of an applied
magnetic field (which determines the size and number
of domains), the charge of the sol particle coming out of
solution (largely a function of sol pH), and the rate at
which the sol particle comes out of solution which is a
function of the drying rate, coating rate, particle mass, sol
concentration, and ambient humidity level.
Using the aforementioned Metglas substrate, we have
found that the presence of the nanoporous structure and
the pore size can be controlled by application of a direct
current (dc) magnetic biasing field which alters the domain structure of the substrate. Figure 1 shows the surface features of a TiO2 film coated onto an as-cast,
demagnetized substrate in a zero magnetic field (earth’s
field was canceled by use of Helmholtz coils). Figure 2
shows an identical film coated upon a magnetically saturated substrate, through application of an 8 Oe dc magnetic field, which contains only one magnetic domain
and no domain walls. As seen in Fig. 2, the resulting film
is smooth on the nanometer length scale indicating uniform deposition of the metal–oxide film and is indistinguishable from films coated using identical process
parameters on nonmagnetic substrates including Si, Al,
Ag, Cu, quartz, and glass.
The origin of this unique nanoporous architecture appears to arise from the interaction of charged sol particles
with the magnetic flux lines of the substrate. A charged
particle moving in a magnetic field is acted upon by the
Lorentz force F, given by F ⳱ Q(v × B), where Q is the
charge of the particle, v the velocity, and B the magnetic
flux density with bold font used to indicate vectors.35
The Lorentz force alters the trajectory of moving charged
particles, forcing them into a spiral about the magnetic
field vector. This same principle is used in magnetron
sputtering where a dc magnetic field, created by placing
magnets under the sputtering target, is used to collect
electrons near the target surface resulting in an enhanced
plasma density.36
For thick ferromagnetic films, the magnetization vector of a domain wall rotates out of plane;37,38 see Fig. 3.
Hence, the moving, charged sol particles collect on the
surface of the substrate at the domain wall, as illustrated
in Fig. 4, replicating the domain wall pattern upward in
sol composition. The electromagnetic dual of this effect
has been used for many decades to image domain walls,
i.e. Bitter patterns.38,39
Our work has shown realization of the nanoarchitectures to be dependent upon the following: (i) magnetic
state of the substrate; (ii) pH of the sol; (iii) drying rate
of the film per relative humidity, temperature, and particle density. Since each of these three variables affects
the nanoporous structure, the effect of each will be discussed separately.
FIG. 3. Illustrative drawing showing the change in the magnetization
vector orientation across domain wall38,39 for a thick film (i.e., a Bloch
wall).
FIG. 2. SEM image of 400-nm-thick sol-gel-derived TiO2 film. Coating was done while the substrate was exposed to a 8-Oe static magnetic field.
FIG. 4. Schematic drawing illustrating the Lorentz force attraction of
moving charged particles to the surface at the domain wall.
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C.A. Grimes et al.: Metal – oxide films with magnetically-modulated nanoporous architectures
II. FABRICATION OF NANOPOROUS METAL
OXIDE FILMS
Nanoporous films of Al2O3, SnO2, and TiO2 have
been deposited using sol-gel. As the steps employed
for obtaining the nanoporous films of these materials
are similar, we discuss here the details of sol preparation and film deposition of TiO2, taking it as a representative case.
The sol for TiO2 films was prepared through controlled hydrolysis and condensation reactions of a metal
alkoxide dissolved in corresponding alcohol in the presence of an acid catalyst.40 The reagents used in the experiment, titanium tetraisopropoxide (TTIP) (99.999%),
2-propanol (99.5%), and nitric acid (70% redistilled),
were procured from Aldrich. A 1 ml volume of TTIP,
0.05 ml of HNO3, 0.1 ml of deionized water, and 32.7 ml
of 2-propanol were used for preparing a 0.1 M TiO2
sol.41 The first step in the sol preparation process was to
dissolve TTIP in 2-propanol. While the solution was
stirred under a nitrogen ambient, deionized water or nitric acid mixed with deionized water was added to it drop
by drop. The sol resulted after stirring the solution for
2 h, which was then covered with parafilm and stored in
a nitrogen environment.
Prior to coating the substrates were cleaned by spraying acetone across the sample, rinsed in deionized water,
and then dried in flowing nitrogen. The substrates were
handled with stainless-steel tweezers throughout the
cleaning and coating process.
The coated substrates were dried in a humidity chamber. Nitrogen was passed through a bubbler at room temperature, with the humidity controlled by the flow rate of
the nitrogen gas (approximately 1.3 cm/s through the
chamber). The humidity and temperature of the chamber
were monitored with a digital hygrometer.
netic field serves to increase both the size of the magnetic
domain and the thickness of the domain wall42 over that
seen in Fig. 1.
We have not yet been able to image the magnetic
domain structures of the as-cast 2826MB substrates; the
very properties that make this substrate interesting to use
make the domain images difficult to obtain. The magnetically soft properties of the substrate are a result of the
alloy having nanodimensional domains, without the grain
boundaries or pinning features associated with magnetically hard materials. However, we have been able to
correlate the topography of the metal oxide structures
seen in Figs. 1, 5, and 6 with the domain structures found
in the magnetically harder amorphous melt-cast Metglas
FIG. 5. SEM image of sol-gel-derived 400-nm-thick Al2O3 nanoporous coating made while exposed to a 0.6-Oe static magnetic field.
The average pore size is approximately 250 nm.
A. Magnetic state of substrate
The influence of the substrates’ magnetic state can be
seen from comparison of Figs. 1, 2, 5, and 6. Figures 5
and 6 show nanoporous features coated onto 2826MB
substrates at applied magnetic field biasing values,
0.6 and 4.2 Oe, respectively, intermediate to those of
Figs. 1 and 2. The 3.0 cm × 1.2 cm as-cast 2826MB ribbons coated in this work have a coercive force of approximately 0.72 Oe. The nanoporous structures are not
dependent upon the particular oxide and have been reproduced for Al2O3, TiO2, and SnO2 when coated on
substrates of similar magnetic state.
The substrate of Fig. 5 was exposed to a 0.6 Oe dc
magnetic field while dip-coated with TiO2. The biasing
field has increased the approximate domain size42 by a
factor of 10 over that seen in Fig. 1, resulting in a pore
size of approximately 250 nm. Application of the mag1688
FIG. 6. SEM image of sol-gel-derived 400-nm-thick TiO2 nanoporous
coating made while exposed to a 4.2-Oe static magnetic field.
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C.A. Grimes et al.: Metal – oxide films with magnetically-modulated nanoporous architectures
2605SC, alloy composition Fe81B13.5Si3.5C21 (upon
which we have not succeeded in realizing the nanoporous
architectures). Figure 7, from Ref. 43, shows the generic
domain structure of a 2605SC film, an alloy similar to the
2826MB composition but magnetically harder and hence
with significantly larger domains.44 – 46 Comparison of
Fig. 7 with our nanoporous metal oxide coatings shows
similarly shaped features which, while not conclusive,
we interpret as evidence that the nanoporous structure is
replicating the magnetic domain structure of the 2826MB
substrate, a magnetically softer alloy and hence with
smaller magnetic domains than its 2605SC counterpart.
The 2826MB substrate of Fig. 6 was coated while exposed to a 4.2-Oe dc magnetic field. At this applied field
value the film is not quite completely saturated, with
scattered domain walls of 1 mm length and 1 micron
thickness still existing between the few remaining domains. As seen in Fig. 6, the nanoporous structure is
established atop the magnetic domain wall; the full structure is approximately 0.6 mm long with an aspect ratio of
approximately 1000. The nanoporous structure atop the
domain wall is indicative that the wall itself, see Fig. 3,
has split into regions having different magnetizationvector orientations. This fractionalization of the domain
wall has been reported43 in Metglas alloy 2605SC;
Fig. 8, a line drawing of a differential phase contrast
domain image from Ref. 43, shows how the domain wall
has split into subregions of different magnetization vector orientation.
The size of a magnetic domain is proportional to the
anisotropy constant K (approximately 4.3 × 104 J/m3 for
Fe50Ni50) and inversely proportional to the saturation
magnetization (0.57 T for 2826MB).44 The thickness of a
domain wall is the result of two competing forces, ex-
change energy and the anisotropy energy.37–39 To a
first approximation the thickness of a Bloch wall is given
by approximately (0.3␲2kTc/4Ka)0.5 37,38 where k is
Boltzman’s constant, Tc the Curie temperature (626 K for
2826MB), and a the nearest neighbor distance (approximately 0.32 nm for Fe50Ni50, assuming a simple cubic
lattice). Calculations show a domain wall thickness of
approximately 45 nm for Fe50Ni50, with the magnetization vector of one-third this thickness, 15 nm, significantly oriented out of the plane of the substrate, which is
in agreement with the wall thicknesses seen in Fig. 1. For
a given substrate the anisotropy constant K, and hence
the domain size and domain wall thickness, can be controlled through by application of a magnetic biasing field
or annealing the substrate in the presence of a saturating
magnetic field.
B. pH of the sol
Sol formation generally occurs through two reaction steps, namely hydrolysis and condensation. In the
case of alkoxide precursors, the hydroxyl ions attach
to the metal atom replacing alkoxy groups. For titanium
isopropoxide the overall hydrolysis and condensation
reaction can be represented as the following:40,47– 49
Ti(OC3H7)4 + 2H2O → TiO2 + 4HOC3H7. The condensation reaction leads to the formation of colloidal particles, which can be polymeric or particulate depending
on the type of precursors and pH of the sol.40,49 The
colloidal particles within the sol undergo Brownian motion and have a tendency to aggregate spontaneously
upon meeting47 to reduce their surface energies.
Colloidal particle aggregation may be inhibited by the
formation of surface charge47,49,50 developed by either
preferential dissociation of one of the lattice ions of the
FIG. 7. Fresnel image and pair of Foucault images (sensitive to vertical and horizontal components of the magnetization vector, respectively) of
two magnetic domain structures seen on Metglas alloy 2605SC. From Ref. 43.
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C.A. Grimes et al.: Metal – oxide films with magnetically-modulated nanoporous architectures
sol particle or preferential adsorption of charged species
from solution. The surface charge is formed either by
protonating, Ti–OH + H+ → Ti–OH2+, or deprotonating,
Ti–OH + OH− → Ti–O− + H2O, the Ti–OH bonds.49 The
pH at which the surface is electrically neutral is called
the point of zero charge (PZC); the surface is negatively
charged at a pH > PZC and positively charged at
pH < PZC.51 For TiO2 the PZC has been reported as
5.250 and 5.5.52
Opposite charges that may be present within the solution tend to accumulate adjacent to the charged surface of
the sol particles and act to screen the charges of the
potential-determining ions40,47,51 forming a double
charge layer. This is illustrated schematically in Fig. 9,
which shows a stationary oxide particle with a surface
positive charge surrounded by diffuse layer of negative
charges in the solution. Figure 10, redrawn from Ref. 40,
shows the potential distribution in the double layer. According to the standard theory, the potential drops linearly through the tightly bound layer of surface charge
and counterions called the Stern layer. Beyond the Helmholtz plane, in the Gouy layer, the counterions diffuse freely40 in the solution. The thermal energy and
electrical energy of counterions in the solution decide the
spatial extent of the diffuse layer.47,53 As the particle
moves, the cloud of opposite charge lags behind; hence,
a moving particle has a net charge associated with it.
If an electric field is applied to a colloid, the charged
particles move toward the electrode of opposite charge
(i.e., electrophoresis).40,54,55 When the particle moves, it
carries along with it the adsorbed layer and part of the
cloud of counterions (negatively charged in this case),
while the more distant portion of the double layer is
drawn toward the opposite electrode.40 The slip plane
(Fig. 10) separates the region of fluid that moves along
with the particle from the region that flows freely. The
rate of movement is controlled by the potential at the slip
plane ␾z56 (see Fig. 10) which is lower for a higher
amount of countercharge screening.
During the drying portion of the dip-coating process
the particles move in random directions, with the continuous evaporation of water and solvent changing the
surface charge and counterion concentration.57,58 In
the magnetic analog of electrophoresis, the Lorentz force
originating with the Bloch domain walls of the magnetic
substrate guides moving charged particles. Hence the
particles agglomerate in a space determined by the magnetic field, leaving solvent in the remaining portions. For
a given domain-wall magnetic field strength, if the domain is too large, the sol particles will be unaffected by
the Lorentz force and deposit within the interior of the
domain walls resulting in a continuous smooth film.
FIG. 8. Drawing of fractionated submicron domain wall structures in
Metglas alloy 2605SC film. The original figure was obtained using
differential phase contrast microscopy.43 Regions within the domain
wall have different magnetization vector orientations.
FIG. 9. Schematic representation of a positively charged sol particle
surrounded by negatively charged ions in the solution.
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FIG. 10. Variation of potential in the double layer. Redrawn from
Ref. 40.
J. Mater. Res., Vol. 16, No. 6, Jun 2001
C.A. Grimes et al.: Metal – oxide films with magnetically-modulated nanoporous architectures
For a given sol concentration and drying rate there
exists a critical sol pH at which the nanoporous architectures are obtained. Figure 11 shows the topologies of the
TiO2 films coated onto 2826MB substrates, in earth’s
field, as a function of acid and water content in the sol.
Beginning with 20 ml of a 0.01 M TiO2 sol, different
amounts of a HNO3:H2O (1:2) solution were added and
the resulting films imaged: (a) 20 ␮l of solution added,
pH ≈ 1.17; (b) 12 ␮l of solution added, pH ≈ 1.29;
(c) 6 ␮l of solution added, pH ≈ 1.37. All films were
dried in a 98% relative humidity environment. As seen in
Fig. 11, the nanoporous structure appears only at specific
acid/water concentrations.
C. Sol particle velocity
The rotation radius of the particle spiraling about the
magnetic field vector due to the Lorentz force, commonly called the Larmor radius,35 is
R=
mv⊥
.
QB
For all sol concentrations, realization of the nanoporous structure requires the films to be dried slowly,
over a period of approximately 15 min, with the films
kept at 98% humidity and 27 °C; rapid drying results in
smooth uniform films like that seen in Fig. 2. The NP
structure is much easier to achieve at low sol concentrations. The 0.01 M sol, of correct acid content, dried in
98% humidity shows the nanoporous structure over
100% of the substrate. With the 0.01 M sol drying at
lower humidity levels the nanoporous structure appears
in isolated islands across the surface of the substrate,
covering approximately 5% of the surface when dried at
77% humidity and approximately 1% of the surface
when dried at 56% humidity. The islands seen at the
lower humidity levels appear to be due to the formation
of slower drying droplets during the coating process. The
nanoporous structures are more difficult to achieve at
higher sol concentrations and result in nanoporous structures with thicker walls; for a 0.1 M sol the nanoporous
structure covers approximately 10% of the surface when
dried at 98% humidity.
(1)
m is the particle mass and vⲚ the particle velocity perpendicular to B. For any given magnetic field strength, a
too heavy, too fast, neutral, or quasi-neutral particle will
not be significantly affected (directed), resulting in deposition of a uniform film. Alternatively too great a sol
particle concentration or too large a stray magnetic field
would result in a cascade of particles upon the surface,
arriving at such a rate that the Lorentz force has little
effect on the particle trajectories resulting in a uniformly
smooth coating.
III. CONCLUSIONS
A new method for fabricating nanoporous metal oxide
thin films is presented. The process appears to involve
the attraction of the moving charged sol particles to the
out-of-plane magnetization vector of the substrate’s domain walls through the associated Lorentz force, thereby
replicating the domain wall structure of the substrate in
the deposited metal oxide film. The substrate used in this
work is a rapid melt-quenched ferromagnetic glass,
FIG. 11. Variation in topology of TiO2 films coated onto 2826MB substrates, in earth’s magnetic field (approximately 0.4 Oe), as a function of
acid content. Beginning with 20 ml of a 0.01 M TiO2 sol, different amounts of a HNO3:H2O (1:2) solution were added to the 20 ml and the
resulting films imaged: (a) 20 ␮l of solution added, pH ⳱ 1.17; (b) 12 ␮l of solution added, pH ⳱ 1.29; (c) 6 ␮l of solution added, pH ⳱ 1.37.
J. Mater. Res., Vol. 16, No. 6, Jun 2001
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C.A. Grimes et al.: Metal – oxide films with magnetically-modulated nanoporous architectures
Metglas 2826MB,1 of composition Fe40Ni38Mo4B18,
which has no long-range order (hence the magnetically
soft properties) but has short-range order over a few atom
lengths. Pore sizes ranging from approximately 40 to
400 nm have been obtained, with the nanoporous structure disappearing completely on magnetically saturated
substrates. Under identical coating conditions we have
seen no evidence of the nanoporous structure when coating nonmagnetic substrates including Al, Ag, Si, Cu,
quartz, and glass.
To date we have proceeded through trial and error,
finding a range of conditions within which we can fabricate nanoporous films of variable pore size. The
nanoporous architectures appear to follow the magnetic
domain structure of the substrate, disappearing when the
substrate is magnetically saturated. We have been unsuccessful within our parameter space in achieving the
nanoporous architectures upon magnetically hard surfaces, i.e., those having large stray fields such as Alnico
and cobalt–samarium magnets, as well as magneticallyharder ferromagnetic glass ribbons of composition
Fe87B3Nb5Si3C2 and Fe81B13.5Si3.5C2 (Metglas alloy
2605SC). It is possible that for the sol concentrations and
drying rates investigated the magnetic attraction is too
strong, leading to such a rapid cascade of particles upon
the substrate that the magnetic template is covered. However for magnetically hard substrates presumably the
right set of templating conditions could be achieved. For
example, it may be possible to fabricate the nanoporous
structures on nonmagnetic substrates, such as silicon, by
adjacent placement of magnetically hard substrates during the coating process. It should also be noted that some
metallic glasses spontaneously form a surface passivation layer,59 which could interfere with the formation of
the nanoporous thin film.
As evidenced in Fig. 11, realization of the nanoporous
architectures is highly dependent upon the acid content
of the sol, which thereby determines the charge of the sol
particle. The nanoporous structures are reliably found
only when dried slowly at high humidity levels; for an
otherwise identical coating process, 100% nanoporous
coverage when dried at 98% humidity (approximately
15 min) goes to approximately 1% coverage when dried
in 56% humidity. The pore size ranges from approximately 60 nm when coated in a 0.0-Oe field to 120 nm
when coated in a 0.4-Oe field and to 250 nm when coated
in a 0.6-Oe field. Though the results of our experiments
show the nature of the porous structure to be dependent
on the magnetic state of the substrates, other factors like
the interaction of alloy composition with sols having
high acid-water concentrations, and their susceptibility to
form oxides, need to be considered. Nanoporous films of
controllable pore size and surface area should find great
utility in catalysis,4 –9 biotemplating,13–16 filtration,60
and sensing61– 63 applications.
1692
ACKNOWLEDGMENTS
This work was supported by the National Science
Foundation under Contracts ECS-9988598, ECS9875104, and NSF DMR-9976851. The authors wish
to thank Professor Marc A. Anderson of the Water
Chemistry and Materials Program, University of
Wisconsin—Madison, for many helpful conversations
and suggestions.
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