One-Dimensional Oxygen-Deficient Metal
Oxides
Wei-Qiang Han†
†
Center for Functional Nanomaterials, Brookhaven National Laboratory, Upton, NY 11973
Abstract Metal oxides are usually non-stoichiometric. Non-stoichiometry affects
their physical properties and chemical reactivity and can lead to novel devices and
new applications in renewable energies. This chapter introduces two types of oxygen-deficient metal oxides. The first is non-stoichiometric oxygen-deficient onedimensional nano-ceria (CeO2-x) and non-stoichiometric oxygen-deficient TiO2-x
nanowires with nanocavities; the second is sub-stoichiometric Magnéli phases TinO2n-1 (4≤ n ≤ 10) nanowires and sub-stoichiometric Cr2O2.4 nanobelts with
modulation structures. The application of 1D-nano-ceria for water-gas shifting reaction also is detailed.
11.1 Introduction
Metal oxides contain at least one metal cation and the oxygen anion (O2-). They usually are semiconductors and non-stoichiometric
under typical experimental processing conditions. Oxides can be
classified into two main categories, As listed in Table 11.1, viz. ptype and n-type, and into four subcategories, metal deficient, metal
2
excess, oxygen deficient, and oxygen excess. [1] For example, cerium oxide (CeO2) falls into the oxygen deficient category. For this
oxide, O2- vacancies are the cause of the metal excess; the oxide will
have the real formula CeO2-x to keep the crystal neutral because two
electrons are needed to be incorporated for each oxygen ion that is
removed. A good site for these electrons is the Ce4+cation. If one
electron is associated with one Ce4+cation, it will change from a Ce4+
ion to a Ce3+ ion. Interestingly, titanium oxide falls into both the
metal-excess and oxygen-deficient categories.
Table 11.1 Semiconducting properties of binary metal oxides of nonstoichiometric composition 1
P-type
semicon-
ductor
Deficit of met-
CoO,
n-type
semicon-
ductor
NiO,
FeO,
-
MnO, Cu2O
al
Excess of oxy-
UO2
-
gen
Excess of met-
-
TiO2, ZnO, CdO
Deficit of oxy-
-
TiO2, CeO2, ZrO2,
al
gen
SnO2, Nb2O5, Ta2O5,
WO3, PuO2, Bi2O3,
PbO2
3
Nanostructured objects have attracted wide attention in recent
years because of their (1) new physics phenomena that affect physical properties; (2) unusual quantum effects and structural properties;
and, (3) promising applications in optics, electronics, thermoelectric,
magnetic storage, and renewable energies. One-dimensional (1D)
systems are realized by creating nanostructures that are quantum
confined in one direction. When the dimensionality of the material is
lowered, the new variable of length scale becomes available to control of materials properties. Then, as the system’s size declines and
approaches nanometer length-scales, it is possible to elicit dramatic
differences in the density of electronic states, opening new opportunities to alter physical- and chemical-properties. [2,3] 1D nanostructures, including nanotubes (NTs) and nanowires (NWs), are used as
elements for miniaturized electrical-, nanofluidic- and opticalnanodevices and have played important roles in renewable energies.
[4, 5]
This chapter will introduce and detail the non-stoichiometric oxygen-deficient 1D-nano-ceria (CeO2-x) and its applications in watergas shifting (WGS) reaction. It also will describe non-stoichiometric
oxygen-deficient TiO2-x NWs with nanocavities, sub-stoichiometric
Magnéli-phases TinO2n-1 NWs and sub-stoichiometric Cr2O2.4 nanobelts (NBs) with modulation-structures.
4
11.2 Oxygen-Deficient 1D-Nano-CeO2-x and its
Applications in the WGS Reaction
11. 2. 1. Crystal Structure of Cubic-Ceria
Cerium with a 4f25d06s2 electron configuration can exhibit both +3
and 4+ oxidation states, and intermediate oxides whose composition
is in the range Ce2O3–CeO2 can be formed. The dioxide CeO2 crystallizes in the fluorite structure. It has a face-centered cubic cell (f. c.
c) with space group Fm3m, (a=5.41134 Å, JCPDS 43-1002). In this
structure, each cerium cation is coordinated by eight equivalent
nearest-neighbor oxygen anions at the corner of a cube and each anion is coordinated tetrahedrally by four cations. The structure, illustrated in Figure 1, is considered as a cubic close-packed array of cerium ions with oxygen ions occupying all the tetrahedral holes.
Fig. 1 The crystal structure of CeO2 in the fluorite structure
The cerium oxides, ranging from Ce2O3-CeO2, earlier were treated
using the classic point-defect model of non-stoichiometry, in which
oxygen-vacant sites were considered to occur randomly in the lattice, in conformity with the law of statistical thermodynamics. Later
experiments indicated that non-stoichiometric phases originating
from the fluorite lattice were formed at low-temperature by the removal of oxygen ions and ordering of the vacancies formed. [6]
5
Reduced ceria results from the removal of O2- ions from the CeO2
lattice, so generating a vacant anion site according to the following
equation:
4Ce4+ + O2- → 4Ce4+ + 2e-/□ +0.5O2 → 2Ce4+ + 2Ce3+ + □ +0.5O2
where □ represents an empty position (anion-vacant site) originating from the removal of O2- from the lattice, here represented as an
oxygen tetrahedral site (Ce4O). Electrostatic balance is maintained
by the reduction of two cerium cations from +4 and 3+.
11. 2. 2 Background of the WGS Reaction
The WGS reaction, typically used to generate H2 through the reaction of a gas mixture of CO and H2O (CO + H2O → H2+CO2, Hº298
= -41.1 kJ/mol), is a well-known catalytic process of industrial importance. [7] Based on thermodynamic and kinetic considerations, a
high conversion of CO is obtained with a two-bed operation at low
(180-250 ºC) and high-temperatures (350-500 ºC). In continuously
operating industrial applications, the classical catalysts employed are
Fe2O3-Cr2O3 for the first stage (high-temperature shift (HTS)), and
Cu/ZnO/Al2O3 for subsequent stages (low-temperature shift (LTS)),
to obtain a good performance under steady state conditions.
For proton-exchange membrane (PEM) fuel cells, the anode catalyst usually is Pt/C, chosen as it is more sensitive to CO that those
mentioned above, because the PEM fuel cell operates at lower temperatures at which CO can de-active the Pt. Usually, the CO in the
fuel must be deeply reduced to < 10 ppm. The WGS reaction is a
critical step in fuel processors for preliminary clean up of CO, and
6
the additional generation of hydrogen before the preferential oxidation of CO, or the methanation step. WGS units sited downstream of
the fuel reformer to further lower the CO content, and improve the
H2 yield. To obtain this equilibrium outlet concentration of CO from
the reformate fuel, the WGS catalyst must be active at low temperatures, 200–280 °C, depending on the inlet concentrations of CO in it.
The reaction is moderately exothermic, with low temperatures resulting in low CO levels; however, the kinetic of the reaction is favorable, even at higher temperatures. [8]
Employing Fe–Cr and Cu–ZnO catalysts in fuel processors poses a
series of disadvantages: The low activity of Fe–Cr as HTS catalyst
and its thermodynamic limitations at high temperatures; the sensitivity of the Cu-ZnO catalyst to air or temperature excursions; the
lengthy pre-conditioning of such catalysts for intermittent operation
(pre-reduction/passivation); and, the large reactor volume dictated
by the slow WGS kinetics of the Cu–ZnO catalyst at low temperatures. Therefore, Fe-Cr and Cu-ZnO catalysts are unsuitable for automotive applications, where the need for fast start-ups dictate the
need for using a low volume of a non-pyrophoric catalyst (fewer
stages). Thus, it is critical to develop efficient safe catalysts for the
WGS reaction in the fuel-cell process.
Ceria recently attracted great interest, particularly for reducing the
emissions of CO, NOx and hydrocarbons from automobile exhaust,
to abate soot formation in diesel fuels, and to minimize CO content
in fuel-cells. [9] The key to these applications is that CeO2-x easily
produces oxygen vacancies in an oxygen-deficient environment,
7
shifting some Ce4+ to Ce3+ ions in the stable fluorite structure. [10]
Oxygen vacancies also are crucial for binding catalytically active
species to ceria. Thus, oxygen vacancies in ceria are considered to
play an essential role in catalytic reactions.[11] Ceria-supported noble-metal co-catalysts, such as Pt-, Au-, Pd-loaded ceria, exhibit
very interesting properties for the WGS reaction with fuel cells. [12]
Under some conditions related to H2 production, the WGS reaction
rates were higher on noble-metals loaded ceria than on commercial
catalysts. [13]
This section discusses the use of pure- and Pd-loaded 1D-nanoceria, a mixture of NTs and NWs, as catalysts for the WGS reaction
at low-temperature.
11.2.3 Synthesis of 1D- Ceria
Various methods are used to prepare special ceria morphologies
with enhanced reducibility. Zhou et al. generated ceria nanoparticles
(NPs) by adding an aqueous ammonium hydroxide precipitant into a
solution of cerium nitrate at room-temperature and then introduced
oxygen into the reactor to oxidize Ce3+ to Ce4+. [14] Chen, et al. obtained ceria NWs via adding this same precipitant into cerium nitrate
at 70 °C, and subsequently allowed it to age at 0 °C for one day. [15]
Yu et al. prepared ceria nanocrystals (NCs) in spherical-, wire- and
tadpole-shapes from a nonhydrolytic sol-gel reaction of cerium (III)
nitrate and diphenyl in the presence of surfactants. [16] Natile et al.
synthesized ceria NPs by two different synthetic routes: Precipitation from a basic solution (sizes around 8-15 nm) and microwave-
8
assisted heating hydrolysis (size around (3.3-4.0 nm). [17] They
found that the NPs made by the latter method were more reduced
than those from the former. Methanol oxidation is also favored on
the ceria NPs prepared by the latter method because of their high
specific area and the presence of greater amount of active sites of
Ce3+ cations. Zr4+ and La3+ doped porous ceria NPs with a high BET
surface area of 160 m2/g exhibited a photovoltaic response, directly
derived from the NPs’ size; normal ceria does not show this response. [18] All these results suggest that ceria with high surface area can increase the Ce3+ ratio that leads to high reducibility.
The 1D-nano-CeO2-x used for WGS reaction, described in the next
section, was synthesized by two successive stages: Precipitation and
aging. At the precipitation stage, 1.5 grams of cerium nitrate
(Ce(NO3)3.6H2O) was added to 15 ml de-ionized water and heated at
100 °C. Once a large amount of vapor formed, 10 ml 5% ammonia
hydroxide solution was added. Very fine yellowish precipitates
formed immediately and the mix started boiling. After 3 minutes, the
solution was transferred quickly to a 0 °C refrigerator. [19] Figure 2
displays the powder profile refinement of the as-produced material
using GSAS/EXPGUI code.
Fig. 2 Powder profile refinement of fresh 1D-ceria
Fig. 3 Pure 1D-ceria sample (a) typical morphology with three kinds of
nanostructures: NPs, NWs and NTs; (b) a high-magnification TEM image of a ceria NT, with a wall of about 5.5 nm thick; and, (c) a high-magnification TEM image of a ceria NR
9
A careful inspection reveals that there are two kinds of 1D
nanostructures of CeO2-x. One is a NW with consistent cross-wise
lattice; while the other is the NT with weak contrast in the middle
(Fig 13.3(a)). These characteristics can be seen more clearly in high
resolution images of a CeO2-x NT and a NW, respectively (Fig. 3(bc)). The TEM image of the pure 1D-nano-CeO2-x shows polycrystalline ceria NWs and NTs (~80 %), together with ceria NPs (~20%)
with a diameter similar to that of the 1D-nano-CeO2-x (Figure 3(a)).
Most 1D-nano-cerias are NWs whose diameters typically range from
6-25 nms, and lengths up to a couple of microns. In Figure 3b, the
selected area diffraction patterns (obtained by fast Fourier transform
(FFT) techniques) in the upper right corner correspond to cubic ceria. The direction of the incident electron-beam is along <110>, i. e.,
the exposed crystal plane is (110). The axis of the CeO2-x NT is
along the <110> direction. Two kinds of lattice-fringe directions attributed to (111) and (200) are observed that, respectively, have an
interplanar spacing of 3.1 Å and 2.7 Å. For most NTs, the thickness
of the wall is almost uniform over the tube, though thickness differs
from tube to tube. Fig. 3c shows the CeO2-x NW has the same crystalline features as the NT. For the 1D-ceria, the preferred exposed
crystal planes for both NWs and NTs are {110} and {100. Based on
electron diffraction analyses and high resolution imaging, the CeO2-x
NPs, NWs and NTs were shown to have the same crystal structure, a
cubic fluorite structure, consistent with the x-ray measurements. The
lattice parameters of the CeO2-x NTs vary from 0.54 nm to 0.56 nm
10
depending on their diameters. In general, the lattice parameter increases with decreasing diameter of the NTs. Cerium-nitrate solution
reacts with ammonium hydroxide to form Ce(OH)3 as an intermediate product with a 1D nanostructure that is retained if the pH of the
reaction is higher than 8.[20] Excess ammonium hydroxide was used
in the present experiment so the intermediate Ce3+ oxidized to Ce4+.
Quickly cooling the samples to 0 °C retained the 1D nanostructure.
The precipitates were dehydrated further, and re-crystallized during
the aging time. Prolonging aging time leads to more 1D-like hollow
structure, i.e. NTs.
Fig. 4 EELS spectra showing different M5 peak intensity for CeO2-x NTs with (a)
d=14.6 nm, (b) d=17.3 nm, and, (c) 25.5 nm. The thicknesses of the wall of the
NTs are 5.5, 6.0, and 10.8 nm for (a), (b) and (c), respectively. The spectra are
normalized for the M4 peak
The increase in the lattice parameter of the CeO2-x NTs implies
that the oxidation state of the CeO2-x NTs may differ from that of
bulk CeO2. EELS (electron energy-loss spectrometry) can analyze
the chemical composition of TEM specimens with a lateral resolution down to about one nanometer. The valence of the cerium ions is
determined from the relative intensity of the white lines (M4 and M5)
of the cerium in the EELS spectra. The NPs are almost completely
reduced to CeO1.5 when the diameter < 3 nm. This reduced CeO2-x
has a fluorite structure, the same as that of bulk CeO2. Also, EELS
spectra taken from the edge and center of the NP indicated that for
large NPs the valence reduction of cerium ions occurs mainly at the
11
surface, forming a Ce1.5 layer and leaving the core essentially as
CeO2. The fraction of Ce3+ ions in the NPs rapidly increased with
declining NP’s size. [21] Fig. 4 shows the M4 and M5 edges of the
EELS spectra from three NTs with diameter, d = 14.6, 17.3 and 25.5
nm. It qualitatively illustrates the systematic change in the EELS
spectra is correlated with the diameters of the NTs, that is the intensity of the M5 edge rises with the decrease in the diameter of the
NTs. To determine the relative amounts of cerium ions Ce3+ and
Ce4+, the second derivative method is used to measure the M5/M4 ratio, since it is insensitive to variations in thickness. The M5/M4 ratio
these three NTs, d=14.6, 17.3 and 25.5 nm , respectively, are 1.27,
1.22 and 1.05; based on M5/M4 being 1.31 for Ce3+ and 0.91 for
Ce4+, the fraction of Ce3+ (Ce3+/[Ce3++Ce4+]) therefore is estimated
correspondingly as 0.90, 0.78 and 0.35. Compared with the CeO2-x
NPs of the same diameter, the fraction of Ce3+ in the CeO2-x NTs is
significantly larger. The main reason is that NTs have two surfaces:
The outer surface and the inner one. Actually, the total surface area
depends on the thickness of the wall of the NTs. If the cerium ions in
the CeO2-x NTs follow the same distribution as that of the CeO2-x
NPs, that is, Ce3+ exists on the surface, while Ce4+ inside, the fraction of Ce3+ mainly would be determined by the thickness of the
wall. In fact, the thicknesses of the wall of the NTs for Fig. 4(a-c)
are about 5.5, 6.0 and 10.8 nm, respectively. Oxygen vacancies in
ceria NT combined with their inner and outer surfaces could offer
more functional, effective features and play an essential role in applications, such as catalytic reactions. Techniques to make high-
12
yield ceria NTs with sustainable stability during the WGS reaction
still is challenging and worthy of further effort. [19]
11.2.4 Testing 1D-Ceria for the WGS Reaction
The in-situ time-resolved XRD experiments were performed at
beam line X7B (λ = 0.922 Å) of the National Synchrotron Light
Source (NSLS) at Brookhaven National Laboratory. The in-situ Ce
LIII-edge x-ray absorption near-edge spectra (XANES) and the insitu Pd K-edge data were collected there, at beam line X19A and
X18B, respectively. [22] The products from time-resolved XRD and
XAFS experiments were measured with a 0-100 amu quadruple
mass spectrometer (QMS, Stanford Research Systems). The portion
of the exit gas flow that passed through a leak valve and into the
QMS vacuum chamber provided the relative pressure of the products. [23]
Fig. 5 Pure 1D-ceria sample (a) a 3D plot of in-situ time-resolved XRD patterns
collected during the hydrogen reduction process. (b) H2 and CO2 relative pressure
during the WGS reaction; (d) TEM image of the sample after the WGS reaction;
and, (d) The lattice parameter of the ceria determined from the in situ diffraction
during WGS and H2 reduction conditions as a function of temperature, which
show relative cell expansion of H2 versus WGS
Samples of 1-2 mg were loaded into a 1-mm sapphire capillary
tube attached to a flow system. The 1D-nano-ceria was exposed in
pure H2 up to 400 ºC for activation before the WGS reaction. [24,
13
25] A similar set up to that used for the WGS reaction was employed
for the temperature-programmed reduction and oxidation, for which
pure H2 and 5% O2 in He were used, respectively. The temperature
ramp rate was ~ 2 ºC/min. The in-situ time-resolved XRD patterns
(Fig. 5a) showed the retention of the cubic-fluorite structure, and
peak widths that were nearly constant during the reduction process,
although there were significant changes in the lattice parameter from
thermal expansion and the partial reduction of the cerium oxide, viz.
from 5.43 Å at 25 oC to 5.47 Å at 400 oC. The reduction of pure 1Dnano-ceria in H2 started at 150 oC, a much lower temperature than
those previously reported for bulk or 3D ceria NPs (i.e. NPs with no
preferred growth in any direction).
Once the catalyst was cooled down to ambient temperature, the
gas was switched to 5% CO/He passing through a water bubbler, at
ambient temperature, at a flow rate of 10 ml/min. The H2O versus
CO vapor ratio was ~ 0.35. After equilibration, the WGS reaction
was carried out isothermally at 200, 250, 300 and 350 ºC, with a
holding time of four hours at each temperature. Figure 5b displays
the relative pressure of H2 and CO2 as a function of time. Catalytic
activity increased with increasing temperature, becoming significant
at 300 ºC. This behavior is very different from that of bulk and 3Dceria NPs, which exhibit a negligible catalytic activity for the WGS
reaction. A series of in-situ XRD patterns collected during the WGS
reaction showed no obvious changes; however, further analysis (see
below) revealed alterations in the number of oxygen vacancies in the
ceria structure, and in the cell dimensions. It was proposed that ceria
14
participates in the WGS reaction when the oxygen vacancies formed
by CO reduction facilitate the breakdown of the H2O to form H2 and
O-2 ions. [26] The number of oxygen vacancies can be determined
from the lattice parameters of the ceria under WGS conditions and
H2 reduction conditions because the cell expands when cerium is reduced from Ce+4 to Ce+3. [27] However, the cell also displays thermal expansion; consequently, one can only compare relative oxidation at a chosen temperature from lattice parameters. Figure 5(c) is
the TEM image of the 1D ceria (~60%) after the WGS reaction.
Compared to the pre-reaction sample, the 1D ceria still are crystalline and the preferred exposed crystal planes are {110} and {100},
though the amount of 1D ceria is slightly decreased. Figure 5d
shows the lattice parameter of the ceria determined from the in-situ
diffraction during WGS and H2 reduction conditions as a function of
temperature. The cell expands more during H2 reduction conditions
than under WGS conditions at 350 ºC. This finding reveals that the
presence of the H2O together with the CO partly re-oxidizes the ceria, and is consistent with the hypothesis that the reaction of H2O at
the O vacancy sites produces adsorbed O2- and H2 gas. [26]
To enhance catalytic activities at low-temperature, Pd (1% in
weight) was loaded on the 1D-nano-ceria. [28] It was deposited on
the 1D-nano-ceria through the drop-wise addition of palladium nitrate (Pd(NO3)2.xH2O) (1 wt%) into an aqueous suspension of 1Dnano-ceria. After stirring the solution for one day, it was washed
with ethanol and then distilled water. Examination under TEM of the
Pd-loaded 1D-ceria sample did not find any isolated Pd0 or Pd-ion
15
NPs. EDS also did not detect a Pd signal, signifying that there was
no Pd-rich area. Figure 6a displays the in-situ XRD patterns of the
Pd-loaded 1D-nano-ceria collected during the reduction process with
pure H2 up to 200 ºC. The cubic-fluorite structure of ceria was stable
during the reduction process, similar to pure 1D-nano-ceria. At room
temperature, no other diffraction peaks were apparent, except those
of 1D-nano-ceria. Around 60 ºC, the diffraction peaks of Pd started
to appear. However, only Pd (111) was observed due to the low concentration of the Pd as well as its broad dispersion of Pd on the 1Dnano-ceria. Together, the XRD and TEM results lead to the conclusion that Pd exists in the starting sample as a non-crystalline species.
Fig. 6 Pd-loaded 1D-ceria sample (a) A 3D plot of in-situ time-resolved XRD patterns collected during the reduction process; (b) H2 and CO2 relative pressure
during the 1st pass of the WGS reaction using the catalyst of Pd-loaded ceria. The
catalyst was at first ramped to 200 °C in H2; (c) H2 and CO2 relative pressure
during the 2nd pass of the WGS reaction using the catalyst of Pd-loaded ceria;
and, (d) The lattice parameters of the ceria determined from the in situ diffraction
during WGS first pass and second pass
In addition, the dimension of the ceria cell still was 5.449Å (see
Table 11.2) when the sample was cooled under an H2 flow to room
temperature. This could reflect the stabilization of the reduced structure of ceria in a flow of pure hydrogen or by the storage of hydrogen in the ceria lattice. [29]
16
Table 11.2 Detailed Analysis of the XRD patterns for Pd-loaded 1D-nano-ceria
catalyst
Sample
Lattice paramCeO2
eter (Å)
Pd
CeO2 from 400
°C peak
Starting mate-
5.423
-
-
5.449
-
-
5.427
5.427
4.035
5.435
5.432
3.948
5.433
-
3.938
rials
After reduced
at RT
Before 1st
WF+GS
Before 2nd
WGS
Before 3rd
WGS
Figure 6b displays the resulting H2 and CO2 relative pressure as a
function of WGS reaction time. Some WGS activity was observed at
200 ºC, and catalytic activity increased at the other three temperatures ranges. The corresponding in-situ time-resolved XRD patterns
displayed structures of cubic ceria and metallic Pd with small
changes in the cells. Under identical experimental conditions, the
WGS reaction was repeated for the same sample in two additional
passes; good catalytic activity was observed during both passes.
Figure 13.6c shows H2 and CO2 relative pressure during the 2nd pass
of the WGS reaction using the Pd-loaded 1D-nano-ceria catalyst.
The XRD patterns of the Pd-loaded 1D-nano-ceria sample after the
WGS runs in CO/He/H2O showed Pd (111) had shifted to lower
17
two-theta angle after the first pass. Table 11.2 gives details of the
lattice parameters of the Pd and ceria. The lattice parameters of Pd
decreased after each pass of the WGS reaction. However, after H2
reduction (RT, H2 flow), the lattice parameter of Pd at 4.035 Ǻ is
much larger than that of the measured Pd lattice parameter of 3.889
Ǻ (powder diffraction file 46-1043). Since the lattice parameter of
PdH0.76 is 4.02 Ǻ (Powder diffraction file 19-0951), it is likely that
the Pd is in the hydride form. The in situ time-resolved XRD pattern
of Pd-loaded 1D-nano-ceria during the pretreatment in pure hydrogen, displayed a shift of the diffraction peak of Pd (111) to a lower
two-theta angle as soon as the temperature was decreased to ambient
levels. This suggests that PdHx forms during cooling. [30] Figure 6d
depicts the lattice parameters of ceria determined from the in situ
diffraction during the WGS first pass and second pass; the cell parameters are found less in the second one. This could be related to
the fall in activity after each pass, since the decline in lattice parameter of ceria indicates a less extent of reduction, and thus, fewer oxygen vacancies as in the succeeding of WGS reaction.
Fig. 7 Pd-loaded 1D-ceria sample (a) Pd K-edge XANES spectra during the reduction process; (b) Ce LIII-edge XANES spectra during the reduction process;
and, (c) Ce LIII-edges XANES spectra during the WGS reaction
In-situ XANES measurements were used to study the oxidation
state of the Pd and Ce in 1D-ceria during the catalytic process. Figure 7a displays the Pd K-edge XANES spectra of the Pd-loaded 1Dnano-ceria sample. The absorption edge of the starting material is at
18
a higher energy position than that of the Pd foil, suggesting that the
Pd is ionic. Also, the much greater white-line intensity implies that
Pd is ionic in the starting material. After purging at room temperature in H2 for three hours, the Pd K-edge spectra of Pd-loaded 1D
nano-ceria resembled that of the Pd foil, indicating the Pd species
was reduced, even at ambient temperature, although the spectrum
still displayed slightly higher white-line intensity than did the Pd
foil; this indicated that the Pd ions are partially reduced. Pd hydride
might also be formed since it has a similar XANES pattern to that of
the Pd foil. [31] However, Pd hydride displays a slightly larger
white-line intensity in the Pd-K edge XANES as well as some shift
in peaks ii and iii to relatively lower energy. For a similar cluster of
Pd or Pd hydride, it was observed that (Er-Eb)R2 = constant, where R
is the inter-atomic distance, Er is the energy of resonance, and Eb is
the energy of a bound state at threshold.[32] Thus, this shift essentially reflects an increase of the Pd-Pd bond distance due to an increase in the lattice parameter during the formation of the hydride,
which the XRD pattern confirms. Also apparent in Figure 7a, at 300
°C, is that the Pd K-edge spectrum of the catalyst is very similar to
that of the Pd foil, pointing to a complete reduction of Pd to its metallic form. Figure 7b illustrates Ce LIII-edge XANES spectrum of
the Pd-loaded 1D-ceria sample during the reduction process; this
edge is shifted to a lower energy position in a pure hydrogen atmosphere at ambient temperature, a finding consistent with the results of
the Pd K-edge XANES, where the sample of Pd-loaded 1D-ceria is
partially reduced at ambient temperature. At 60 ºC, the Ce LIII-edge
19
shifted to even lower position, indicating further reduction of the ceria. At 200 ºC, the shift of the Ce LIII-edge does not shift much more.
We also noted that there is a shoulder peak at 5727.4 eV in pristine
Pd-ceria catalyst (highlighted with an arrow), pointing to the existence of some Ce3+ ions. The intensity of this shoulder increased
gradually as the temperature rose, suggesting that the concentration
of the Ce3+ ions increased. In CO/H2O, the Ce LIII-edge spectrum of
the Pd-loaded 1D-nano-ceria was similar to that in H2 atmosphere at
ambient temperature, indicating a slightly oxidation of the catalyst
after it was fully reduced in H2 at 200 ºC (Figure 7c). During the
WGS reaction, the Ce LIII-edge position shifted to lower energy at
350 ºC, while its intensity increased at 5727.4 eV, pointing to the reduction of ceria in the Pd-ceria catalyst under the conditions of the
WGS reaction. [28]
A redox mechanism might be at the basis of the reaction. [13] Using σ to designate an adsorption site, the mechanism can be written
as follows:
CO + σ → COad
H2O + σ → Oad + H2
COad + σad → CO2 +2σ
The catalyst is oxidized by H2O and reduced by CO. The adsorption sites for COad and Oad need not be the same. For Pd/ceria catalysts, CO might adsorb on the Pd and then reduce the ceria near the
Pd interface; subsequently, H2O re-oxidizes the reduced ceria. [13]
To understand more comprehensively the interaction between the
loaded-Pd clusters and the 1D-nano-ceria, the changes in the ceria
20
lattice parameters of pure 1D-nano-ceria, Pd-loaded 1D-nano-ceria
and ceria 3D-NPs during the hydrogen-reduction process were compared (Figure 8). For pure 1D-ceria, the lattice parameter of ceria
barely changed at the beginning of the ramping process, but increased after the onset temperature of 150 ºC. At 350 ºC, the lattice
expanded by ~ 0.04 Å compared with that at 150 ºC, as the consequence of the reduction of Ce4+ to Ce3+ ions, since the radius of the
Ce3+ (1.14 Å) is much larger than that of the Ce4+ (0.97 Å). This
temperature at which reduction occurred was much lower than that
for ceria NPs (490 ºC). [33] During the hydrogen-reduction process,
a dramatic increase of the ceria lattice parameter (0.025 Å) was observed at about 60 ºC for Pd-loaded 1D-nano-ceria, and concurrently, metallic Pd appeared. This finding suggested that the interactions
between the Pd clusters and the 1D-nano-ceria might facilitate the
reduction of Pd-loaded 1D-nano-ceria, which was consistent with
reports that Au (or Cu) can activate the surface of ceria and decrease
its reduction temperature. [34]
The Pd-loaded 1D-nano-ceria could be reduced at temperatures
~100 ºC lower than pure 1D-nano-ceria, wherein large number of
oxygen vacancies might be produced to participate in the formation
of the active sites for WGS reaction. As observed from Figure 8,
pure 1D-nano-ceria was significantly reduced at 300 ºC, leading to
the good catalytic activity observed in Figure 5b, mainly due to the
1D conformation of the ceria, whose preferred exposed crystal
planes are {110} and {100}. The {110}/{100} dominated surface
21
structures are more reactive for CO oxidation than the {111}dominated one.
Fig. 8 Changes in ceria lattice parameters during the temperature ramping of 1Dnano-ceria, Pd-loaded 1D-nano-ceri,a and 3D-ceria NPs
Fig. 9 TEM images of Pd-loaded 1D-ceria reduced pre-treated in hydrogen at 200
°C after WGS reaction (a) NWs; and (b) nanoparticles; TEM images of Pd-loaded
1D-ceria pre-treated in hydrogen at 400 °C after WGS reaction (c) NWs; and (d)
nanoparticles; (e) Schematic image of the evolution of (A) a NW breaking down
into faceted NPs; and (B) subsequently changing to nn-faceted 3D NPs
Long-term stability is a concern with all WGS catalysts and particularly so with ceria-supported precious metals. There is no consensus on what processes are most important. [13] Zalc et al. proposed that deactivation occurs from ceria over-reaction. [35] Goguet
et al. suggested deactivation in the Pt/ceria catalysts is due to formation of carbon during reaction at 400 ºC as a result of CO dissociation. [36] Wang et al. considered that the deactivation of Pt/ceria
and Pd/ceria catalysts under the WGS reaction originated from the
loss of metal dispersions since the WGS rates of a series of Pd/ceria
catalysts increased linearly with surface area of the Pd. [37] To understand the deactivation mechanism, we studied the effect of pretreating the catalysts with hydrogen by combining the activity measurements with subsequent TEM measurements. [28] For Pd-loaded
1D-nano-ceria, the one pre-treated in H2 at 200 ºC had much higher
activity than one treated in H2 at 400 ºC. Thus, the relative WGS
catalytic activity at 250 ºC for samples of Pd-loaded (1%) 1D-nano-
22
ceria pre-treated in H2 up to 400 ºC and 200 ºC are, respectively,
1.83% and 2.96%. For both such samples, 25% of the 1Dnanostructures remained after the WGS reaction. Also, the preferred
exposed crystal planes for these remainding1D-nano-cerias were
{110} and {100} for both samples (Fig. 9a and c). In both samples, a
large number of 1D-ceria nanostructures were broken down into NPs
after a couple of redox cycles. The TEM image, taken after WGS reaction, of the NPs of the sample pre-treated in hydrogen at 200 °C
often resemble polyhedrons with clear crystalline facets, and the preferred exposed surfaces still are {100} and {110} (Fig. 9b). However, many NPs of the sample pre-treated in the hydrogen atmosphere
at 400 °C after WGS reaction do not have clear crystalline facets, or
a preferred crystalline orientation (Fig. 8d). Besides {100} and
{110}, there are many other exposed surfaces, such as (111) that is
an unfavorable surface for CO oxidation. [38] This might explain
why the 1D-nano-ceria sample pre-treated in hydrogen at 200 °C has
a better WGS activity. Another factor in the decreasing of surface
area might be the aggregation of NPs after the reaction. In the present work, a higher reducing temperature introduced more changes
in crystallographic faces to those that do not favor the activity. Figure 13.8e shows a schematic image of the evolution of a NW breaking down into faceted NPs, and subsequently changing to nonfaceted NPs, i.e. 3D NPs. The loss of the effective faceted surfaces
of ceria during the WGS process could be responsible for the loss of
catalytic activity.
23
11.3 Sub-stoichiometric Magnéli Phases 1D-TinO2n-1
Stoichiometric titanium dioxide (TiO2) is one of the most widely
studied transition-metal oxides because of its many potential photoelectrochemical applications. [39] Although many semiconductors
were proposed as photoelectrodes, the general conclusion is that
TiO2 is the one of the best due to its excellent chemical stability and
resistance to corrosion. [40] There are three forms of stoichiometric
TiO2 crystals: Orthorhombic brookite, tetragonal anatase, and rutile.
[41] Electronically, these three TiO2 phases are wide-band-gap nonconducting materials. However, the wide bandgap and high electrical-resistivity of stoichiometric TiO2 are two main factors precluding the realization of the requirements for a highly efficient photoelectrode.
Hydrogen reduction of stoichiometric TiO2 generates nonstoichiometric TiO2-x with significantly better conductivity. [42] Many researchers explored the relation in nonstoichiometric TiO2-x between
its electrochemical applications and its microstructure, composition,
defect formation, conduction mechanism, and optical properties. In
particular, studies focused on the sub-stoichiometric titanium oxides,
TinO2n-1 (i.e., the Magnéli phases, where n is a number between 4
and 10 (i.e. 1.90 ≥ x ≥1.75)). [43, 44] The Magnéli phases comprise
two-dimensional chains of octahedral TiO2, with every nth layer
missing oxygen atoms to accommodate the loss in stoichiometry.
The formed superstructures display recurrent shear planes wherein
the MO6 octahedra exhibit edge-sharing, rather than vertex sharing.
Consequently, there is added disorder in such superstructures com-
24
pared to standard idealized crystallographic phases. Magnéli phases
are insoluble in acid, electrochemically stable, and have high electronic conductivity; accordingly, they served as gas sensors for detecting hydrogen, oxygen, and carbon monoxide, in battery electrodes,
antireflection
coatings
in
solar
systems,
and
in
photoelectrolysis. [45]
In this section, a simple route is described to synthesize Magnéli
phases 1D-TinO2n-1. Meanwhile, we measured their electrical
transport and optical properties.
The synthesis process entails two steps. First, the intermediate
product, H2Ti3O7 NWs, is produced by treating anatase TiO2 particles with NaOH in an autoclave at 160-180 ºC for 2-3 days, and then
subsequently washing the material with nitric acid. X-ray diffraction
(XRD) measurements revealed that the intermediate product is monoclinic H2Ti3O7. [46, 47] Scanning electron microscopy (SEM) and
transmission electron microscopy (TEM) images showed that most
products are straight NWs of 30- 200 nm diameter, and up to 10 μm
long. In the second step, the intermediate product is heated under
hydrogen for 1-4 hours at 850-1050 ºC.. The final products from the
reduction treatment were dark, rather than being white, like the intermediate H2Ti3O7. [48]
Figure 10 (a) shows an XRD spectrum of a sample of the final
product. The characteristic spectral peaks denote that the sample is
triclinic Ti8O17 (JCPDF# 50-0790, a= 5.526 Ǻ, b= 7.133 Ǻ, c=
44.059 Ǻ, α= 66.54 º, β= 57.18º, γ = 108.51 º). Figures 10 (b) and (c)
are, respectively, an SEM image and a low-magnification TEM im-
25
age of the sample. Both show many NWs with diameters typically
ranging from 30 to 200 nm, similarly to the intermediate H2Ti3O7
NWs. Figure 10 (d) is a high-magnification TEM image of a part of
a single-crystal NW. The labeled lattice distance is 6.4 Ǻ, corresponding to the (004) plane of the triclinic crystal. After heating, the
intermediate H2Ti3O7 is heated at 1050 ºC in a hydrogen atmosphere, the XRD spectrum reveals that the material is triclinic Ti4O7
(JCPDF# 50-0787, a= 5.5965 Ǻ, b= 7.11259 Ǻ, c= 12.4647 Ǻ, α=
95.05 º, β= 95.18º, γ = 108.77 º). The TEM image shows that most
products are quasi one-dimensional fibers, with diameters of about 1
micrometer, i.e., much larger than that of the intermediate H2Ti3O7
fibers.
Fig. 10 (a) XRD spectrum of the sample prepared by heating intermediate
H2Ti3O7 at 850 ºC under a hydrogen atmosphere; (b) a SEM image of the sample,
(b) its low-magnification TEM image; and, (c) a high-magnification TEM image
of a part of a NW
Nanocavities embedded within anatase TiO2 NWs were generated
by heating intermediate H2Ti3O7 NWs in an oxidation atmosphere at
650 ºC. During the reaction, monoclinic H2Ti3O7 transforms into tetragonal anatase TiO2 and H2O (H2Ti3O7 → 3TiO2 +H2O). Figure 11
(a) shows numerous polyhedral nanocavities inside the TiO2 NWs;
Fig. 11(b) is a high-resolution TEM image viewed along the [100]
direction. The inset is the fast Fourier transform from the whole area. The nanocavities have polyhedral shape. The boundaries between the nanocavities and the NW body are sharp. The boundaries
26
planes are {01-1} {100}, and {001}, which all are low-index planes
of the anatase crystal and have the lowest surface formation energies
(J/m2) of 0.44, 0.53 and 0.90, respectively. Examining the nearedge-fine structure of the Ti L-edge and O K-edge (see Fig 11 (c)),
revealed that the Ti/O ratio in the area of the nanocavity is found to
be 18% higher than in the “body” of the NW. For comparison, an
average spectrum, taken at lower magnification also is included. In
addition, the Ti L3/L2 ratio increases at the nanocavities, suggesting
a decrease in the Ti valence, which might be due to a decline in the
O stoichiometry of the TiO2 around the nanocavity. [49]
Fig. 11 (a) Low magnification TEM-image of the TiO2 NWs, showing the high
density of nanocavities; (b) High-resolution image viewed along [100] direction,
showing polyhedral nanocavities in the NW; (c) Core-loss EELS spectra from the
center of the nanocavity and the body of the NW
For forming the Magnéli phases, the H2Ti3O7 intermediate NWs
were treated at temperatures higher than 850 ºC under a reduction
environment. The chemical reactions for transforming intermediate
H2Ti3O7 to the final Magnéli phases, Ti8O15 and Ti4O7, are expressed as
8H2Ti3O7 + 3H2 → 3Ti8O15 + 11H2O
4H2Ti3O7 + 3H2 → 3Ti4O7 + 7H2O
During reduction processes at 850oC, monoclinic H2Ti3O7 NWs
change into triclinic Ti8O15 NWs. This latter phase is made up of
two-dimensional chains of TiO2 octahedra, wherein every 8th layer
has missing oxygen atoms to accommodate for the loss in stoichi-
27
ometry. At a higher temperature (1050 ºC), monoclinic H2Ti3O7
NWs are transformed into triclinic Ti4O7 fibers, viz. twodimensional chains of titania octahedra, wherein every 4th layer has
oxygen atoms missing to accommodate the loss in stoichiometry.
The oxygen-deficient density of Ti4O7 is double that of Ti8O15. This
higher level of oxygen-deficiency entails more changes in the materials’ morphology and structure. Intermediate H2Ti3O7 NWs often
are closely packed, and during high-temperature exposures, these
bundles may merge to form large-diameter masses. Nevertheless, areas of the NWs with numerous defects still break off, generating
short thick fibers.
Figure 12 (a) shows the UV-visible diffuse reflectance spectra of
the starting trititanate NWs, the TiO2 NWs with nanocavities and the
Ti8O15 NWs and the Ti4O7 fibers. Very different from the starting
trititanate NWs and TiO2 NWS, whose absorption band covers only
the UV region, the optical absorption bands of Ti8O15 and Ti4O7
cover the full range of visible-light wavelengths and extend into the
near-IR region. This difference is attributed to the different ionization states of oxygen vacancies. [50]
Figures 12 (b) and (c) show the I-V curves of the Ti8O17 and Ti4O7
samples, respectively, measured at room temperature; clearly, both
samples exhibit high electrical conducting behavior at room temperature. The resistances of the Ti8O17 and Ti4O7 samples, respectively,
are 9.74Ω and 0.965Ω at room temperature; their corresponding
electrical conductivities at room temperature are 0.236 S/cm and
10.36 S/cm. The electrical conductivity of the Ti8O15 and Ti4O7
28
samples also were measured by dipping them into liquid nitrogen,
starting from room temperature. The values decreased by about four
orders-of-magnitude to 2.410-5 S/cm and 410-3 S/cm, respectively,
reflecting the lesser thermal excitation of electrons at lower temperature. The electrical resistivity of bulk Ti4O7 was recorded previously.
Ti4O7 shows two conductivity transitions, a semiconductorsemiconductor one in the range of 130-140 K, and a semiconductormetal one at 150 K. [51] For both transitions, there is steep increase
of the electrical conductivity with increasing temperature. In the metallic high-temperature phase, 3d electrons of Ti ions are delocalized
and contribute to the electrical conductivity. Below 150 K, the 3d
electrons are localized to form covalently bonded Ti3+–Ti3+ pairs.
Since pair formation involves lattice displacement, these pairs may
be considered as two-particle polarons, or bipolarons. While the ordered arrangement of the Ti3+ pairs is established in the semiconducting low-temperature phase (T<130 K), this long-range order
vanishes
during
the
intermediate
temperature
phase
(130 K<T<150 K) and then is described as a liquid state of the bipolarons. [52]
Fig. 12 (a)
UV-visible diffuse reflectance spectra of of the starting
H2Ti3O7 NWs, the products of TiO2, Ti8O15 and Ti4O7. Room-temperature I-V
curves of (b) Ti8O15; (c) Ti4O7; and, (d) anatase TiO2
The electrical conductivity of TiO2-x depends strongly on the conditions of treatment that engender different oxygen deficiencies,
29
suggesting that such deficiencies play a crucial part in electrical
conductivity. TinO2n-1 can be considered as being made up of (n-1)
TiO2 octahedra and one TiO octahedral. Rutile (or anatase) TiO2 is a
well-known wide-bandgap semiconductor. In contrast to TiO2, its
crystal structure of TiO is based on a NaCl structure with ordered
vacancies in both the metal- and the oxygen-sublattices; one-sixth of
the titanium atoms, and one-sixth of the oxygen atoms are missing.
The existence of the vacant sites within the TiO structure is thought
to permit sufficient contraction in the lattice such that 3d orbitals on
titanium overlap, thereby broadening the conduction band and allowing electronic conduction. [53] The metallic behavior of TiO is
responsible for the metallic conductivity of TinO2n-1. Ti4O7 has much
higher TiO density than that of other TinO2n-1, including Ti8O15, and
thus, Ti4O7 is the most highly conductive phase. This is consistent
with the above measurement.
11.4 Sub-Stoichiometric Chromium Oxide Nanobelts with
Modulation Structures
The physical and chemical properties of transition metal oxides
can be changed dramatically because the transition metal elements
can have different oxidation states. The Cr atom has a 4s1 3d5 configuration. Chromium, with different valence charges, forms different chromium oxides, such as CrO3, Cr2O5, CrO2, and Cr2O3. Chromium dioxide (CrO2) is a half-metallic ferromagnet with almost
100% spin polarization at the Fermi level such that electron conduc-
30
tion only occurs with one orientation of electron spin; the other spin
direction is insulating. CrO2 has a Curie temperature of 397 K and
thus offers promising applications in magnetic tunneling and spin injection devices. [54] However, CrO2 is a metastable oxide that readily decomposes to thermodynamically stable chromia (Cr2O3), which
is an antiferromagnetic insulator with a Néel temperature of 307 K
and is suitable as a tunnel junction barrier both below and above this
temperature. [55] For antiferromagnetic nanomaterials, the surface
spins increase as the size declines. The effect of uncompensated
spins on their surfaces is an important factor that may affect their
magnetic properties. Néel suggested that very fine particles of an antiferromagnetic material should exhibit particular magnetic properties such as superparamagnetism and weak ferromagnetism. [56]
Size and shape are two important factors affecting the nanostructures’ properties and performances. Chromia NPs and nanopores
were prepared by several methods, including a sol–gel process, gas
condensation, microwave plasma, a sonochemical reaction, laserinduced pyrolysis, hydrazine reduction followed by thermal treatment, mechanochemical processing, urea-assisted homogeneous
precipitation, and a precipitation-gelation reaction. [57] In this section, I introduced a facile way to prepare chromia NBs and NWs,
and describe two types of superlattice-structures related to oxygendeficiency.
The easy synthesis route for the chromia NBs and NWs is as follows: First, a piece of chromium was inserted into the hot zone of a
quartz tube horizontal furnace. Argon was used as the protective gas
31
during the rising temperature stage. Once the temperature reached
800 ºC, an ethanol vapor is introduced while keeping the temperature steady for 1 hour. Next, the flow of ethanol vapor was stopped,
but argon flow was maintained until the sample naturally cools
down. [58]
Figure 13a is the SEM image taken from the as-produced sample;
most of the products are NBs, with some NWs. The lengths of the
NBs are several micrometers and their width ranges from tens nanometers to several hundred nanometers; their thickness is about a
few nanometers to tens of nanometers. In general, the width and
thickness of an individual NB are quite uniform. Interestingly, NBs
with large roots and a folded shape along the axis direction also
were observed. An X-ray spectrum (Fig. 13.13b) of the NBs indicates that peaks can be indexed as rhombohedral Cr2O3 (JCPDF#
38-1479, space group R-3c (167), a= 0.495876 nm and c=1.35942
nm).
Fig. 13 (a) SEM image showing the typical morphology of growth of NBs and
NWs on the Cr substrate; (b) XRD spectrum of rhombohedral Cr2O3 NBs. Cr
(110) peak originated from Cr substrate; (c) High resolution image. The inset is
the corresponding electron diffraction patterns with beam parallel to the [001] direction. The electron diffraction patterns are indexed as rhombohedral Cr2O3
Fig. 13c is the high resolution image of a single-crystalline NB.
The inset is the corresponding electron diffraction patterns from the
NB with the beam parallel to the [001] direction. Tilting experiments
confirm that the basic structure of the NBs is rhombohedral with a=
32
0.496 nm and c=1.359 nm. Many NBs are single crystalline with the
surface normal being the [001] direction.
Fig. 14(a) is a high resolution image from another NB with the
beam along the [-111] direction. The top-left part in the image
shows the basic lattice image of Cr2O3, while the bottom-right part
exhibits additional modulation fringes as indicated by the arrows.
Moiré effects are ruled out since there are no slightly mismatched or
misorientated lattices, as shown in both the high-resolution image
(Fig. 14(a)) and the electron diffraction patterns (Fig. 14(b)). The
diffraction patterns (Fig. 14(b)) shows the satellite spots, indicating
the structure modulation along the [110] direction. The modulation
wave vector was determined as q1=1/3(a*+b*), and the modulation
periodicity in the real space is 7.4 Å (referred to as the modulation
I). Besides this commensurate modulation along [110] direction,
there is another incommensurate modulation along about the [1-23]*
(or [0-51]) direction (referred to as the modulation II), as shown in
Fig. 14(c) and 14(d). The modulation wave vector in Fig. 14(c) is
determined as q2≈0.133a*-0.267b*+0.401c*, and the periodicity in
the real space is 16.4Å. The modulation shown in Fig. 14(d) is similar to that in Fig. 14(c) but with slightly longer periodicity (18.4 Å)
in real space. Most NBs have the second kind of modulation structure, though the modulation periodicity among the different NBs differs slightly.
Fig. 14 High resolution image (a) and electron diffraction pattern (b) from a
NB with modulation I. There is no modulation at the top-left part as shown in
33
the image and the top-left inset which is the Fourier transfer from the top-left
area. The bottom-right part, however, exhibits modulation I as shown in the
image and the bottom-right inset which is the Fourier transfer from the bottom-right area. The superlattices also are present in the diffraction (b); (c)
High resolution image of modulation II with periodicity of 16.4 Å; (d) High
resolution image of modulation II with periodicity of 18.4 Å; (e) EELS spectra
from NBs: A without modulation; B with modulation I; C with modulation II;
and, D referenced micron-sized Cr2O3 particle
To determine the origin of the modulation, we performed EELS
analysis for the different NBs. Fig. 14e (A, B and C) shows EELS
spectra from the Cr2O3 NBs without modulation, with the first kind
of modulation, and second kind of modulation, respectively. For
comparison, a reference EELS spectrum was acquired (Fig. 14e-D)
from a micron-sized Cr2O3 particle (Alfa-Aesar Co.). The spectra
were normalized to equalize their maximum intensities of Cr-L3.
The quantitative calculations show that the atomic ratio of oxygen to
chromium for the NBs without modulation (Fig. 14e-A) is
rO/Cr=1.47, close to that of the micron-sized Cr2O3 particles (~ 1.5).
However, the rO/Cr of the NBs with the modulations usually is markedly lower than that of micron-sized Cr2O3 particles. The rO/Cr of the
NBs with modulation I (Fig. 14e-B) and II (Fig. 14e-C) are 1.18 and
1.22, respectively, indicating the oxygen-deficiency in the nanobelts
with modulation structure. A series of EELS analysis shows that the
composition of the NBs with the modulation is Cr2O3-x with x ranging from 0.25 to 0.70. Interestingly, many NBs with modulation
have a composition close to Cr2O2.4. We also checked the energy
34
range related to C k-edge by EELS and do not find other elements in
the NBs. TEM studies show that about 80% NBs have the modulation structures.
Pop et al. reported an oxygen-deficient Cr2O2.4 powder that they
produced. [59] They first synthesized an intermediate-material
Cr(OH)3-gel using a sol-gel method:
2CrO3 + 3C2H5OH → 2 Cr(OH)3 ↓ + 3CH3COH
Then, the gel was heated at 1250 K in hydrogen atmosphere in
order to stabilize the oxygen vacancies. XRD study showed the crystal structure of the non-stoichiometric Cr2O3-x still is rhombohedral,
the same as stoichiometric Cr2O3, but with a slight differences the
peaks’ relative intensity. The differences arise from the oxygen deficiency in the non-stoichiometric compound due to a modification of
the structural factor ǀFhklǀ2, and a modification of the peak intensity.
The X-ray K edge spectra revealed that the composition of the nonstoichiometric Cr2O3-x is Cr2O2.4. Measurements of the temperature
dependence of the reciprocal susceptibility between 100 and 1200 K
disclosed an anti-ferromagnetic behavior, with the Néel temperature
at 318K. For the non-stoichiometric sample, there is a linear temperature dependence of the reciprocal magnetic susceptibility in the
paramagnetic region. The effective magnetic moment per chromium
ion is μeff = 4.86 μB, a big difference from 3.83 μB for a stoichiometric Cr2O3. [59]
Though the synthesis route, morphology, and microstructure described here is different from those reported by Pop et al., many
nanobelts with modulation in the present work have a composition
35
close to Cr2O2.4, The oxygen vacancies in the nanobelts with modulation structures are ordered, and thus form the superstructures. In
the present work, the chemical reaction for the formation of chromia
nanobelts can be represented as [58]
2Cr + (1.5-0.5x) O2 → Cr2O3-x
Oxygen here mainly comes from the decomposed ethanol at hightemperature. Though the reaction-tube was purged with 99.999% argon before heating, it still may have contained some remnant oxygen. At high-temperature, Cr, like some transition metals such as W,
[60] very easily is oxidized even at reduced atmosphere. The presence of ethanol vapor ensures Cr can be oxidized, but under an oxygen-deficient condition. To understand the role of ethanol, ethanol
vapor is replaced with water vapor or air, no NBs or NWs were
formed signifying that ethanol plays an important role in their formation.
The Cr2O3-x NBs grow in the oxygen-deficient environment, so
that most of them display oxygen deficiency, as confirmed by the
EELS measurements. The ordered superstructures occur in the NBs
that contained oxygen defects, or vacancies, but not in NBs lacking
oxygen defects. Hence, the modulations found in the oxygendeficient NBs reflect the ordering of the oxygen vacancies. This
finding also is consistent with the XRD measurements, since the xray scattering is insensitive to oxygen. The oxygen deficiency chromia NBs with modulation structures and folded shapes described
here are expected to have distinct anti-ferromagnetic and catalytic
properties.
36
11.5 Summaries
1. 1D-nano-cerias (a mixture of NTs and NWs) were prepared by a
mild hydrothermal reaction route, with cerium nitrate and ammonia
hydroxide as reactants. High precipitation temperature and prolonged aging-times were keys for the formation of tubular structure.
The fraction of Ce3+/Ce4+ ions increased with decreasing NTs diameters. High catalytic activity of pure and Pd-loaded 1D-nano-ceria
were observed in the WGS reaction at low temperature. Both ceria
and Pd played important roles during the WGS reaction process.
While 1D-nano-ceria was reduced easily at low temperature to produce oxygen vacancies, Pd can activate the reduction of ceria. The
special 1D feature could increase the catalytic activity of nano-ceria
efficiently because the effective surface area is extended by mitigating the problem of aggregation of NPs and whilst benefiting from
the double surfaces of NTs. Furthermore, the preferred exposed
crystal planes for 1D-nano-ceria are {100}/{110}, this is, the favored surfaces for CO oxidation. The loss the effective faceted surfaces of ceria during the WGS process could be responsible for the
loss of catalytic activity.
2. Ti8O15 NWs and Ti4O7 fibers were prepared by reducing
H2Ti3O7 NWs at different temperatures. Both samples show high
electrical conducting behavior at room temperature and their absorption bands cover the full visible-light region and extend into the near
37
IR region. Nanocavities embedded within anatase TiO2 NWs were
generated by heating H2Ti3O7 NWs in an oxidation atmosphere at
650 ºC. The Ti L3/L2 ratio increases at the nanocavities, suggesting a
decrease in the Ti valence that might be due to the decline in O stoichiometry of the TiO2 around the nanocavity.
3. Two types of unknown superlattice-structures were found, related to oxygen-deficiency in chromium oxide NBs. The first, is a
commensurate modulation along the [1 1 0] direction, and the second one is an incommensurate modulation along about the [1 -2 3]*
(or [0 -5 1]) direction. Cr2O2.4 NBs are produced by simply heating a
piece of chromium under an ethanol atmosphere.
Acknowledgements
The author is grateful to Drs. L. J. Wu, W. Wen, J. C. Hanson, X.W.
Teng J. A. Rodriguez, Y. M. Zhu in Brookhaven National Laboratory
and Mr. Y. Zhang in SUNY at Stony Brook for their cooperation and
helpful discussions. This work is supported by the U. S. DOE under
contract DE-AC02-98CH10886 and Laboratory Directed Research
and Development Fund of Brookhaven National Laboratory.
38
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