Ceramic Materials and Nano-structures for Chemical Sensing
Abdul-Majeed Azad Department of Chemical and Environmental Engineering The University of Toledo, 3052 Nitschke Hall, Toledo, OH 43606-3390.
[email protected]
Sheikh A. Akbar Department of Materials Science and Engineering The Ohio State University, 2041 Watts Hall, College Road, Columbus, OH 43210.
[email protected] ABSTRACT High selectivity, enhanced sensitivity, short response time and long shelf-life are some of the key features sought in the solid-state ceramic-based chemical sensors. Since the sensing mechanism and catalytic activity are predominantly surface-dominated, benign surface features in terms of higher aspect ratio, large surface area and, open and connected porosity, are required to realize a successful material. In order to incorporate these morphological features, a technique based on rigorous thermodynamic consideration of the metal/metal oxide coexistence, is described. By modulating the oxygen partial pressure across the equilibrium M/MO proximity line, formation and growth of new oxide surface on an atomic/ submolecular level under conditions of ‘oxygen deprivation’, with exotic morphological features has been achieved in a number of metal oxides that are potential sensor materials. This paper describes the methodology and discusses the results obtained in the case of two model systems, viz., tungsten oxide (WO3) and titanium oxide (TiO2).
Key words: Semiconducting ceramic oxides, Redox reactions, Microstructure, Chemical sensors
Proceedings of Optics East SPIE Conference on Sensors for Harsh Environments II conference, Boston, October 23-26, 2005, vol. 5998, 1-15 (2005).
1. INTRODUCTION One aspect of current interest and great relevance to the fundamental understanding of the behavior of materials is the role of dimensionality and size on their optical, chemical and mechanical properties for application in a wide range of devices. Owing to the nanoscale features, one-dimensional systems exhibit novel physical and chemical properties that can be exploited in optics, catalysis and data storage devices. Hence, such systems are being synthesized and studied in great details. For instance, polymer nanofibers are used as selective gas separation membranes, filters, biomedical materials (drug carriers and wound dressings), protective clothing, space mirrors, and precursor platforms or scaffolds for the nanotube/nanowire synthesis. Thus, they become model systems to study and correlate the theoretical explanations that are still in progress. Such behavior is almost nonexistent in the bulk material where the particle size is in the micron level. There is growing interest in introducing such attributes in nanoscale inorganic materials as well. The most obvious advantage of doing so is the possibility of their application as quantum dots in a host of devices, such as MEMs, lab-on-a-chip sensors/detectors, structural elements in artificial organs and arteries, reinforced composites, micro solar cell electrodes, micro fuel cells, photocatalysts (splitting of water and in the deactivation of chemical and biological weapons), and electrocatalysts, to name a few. One area where nanofeatures in the materials are of immense relevance is the field of solid-state ceramic-based chemical sensors. High selectivity, enhanced sensitivity and short response time are some of the key features sought in these devices. Since the sensing mechanism and catalytic activity of ceramics are largely microstructure-dominated, benign surface features such as small grain size, large surface area, high aspect ratio and, open/connected porosity are required to realize a successful sensor material1-4. A novel technique employed to impart such attributes by modifying the microstructural artifacts of ceramic-based sensor materials is described in this paper. The effect of the variation in the ambient oxygen partial pressure across the metal/metal oxide boundary on the microstructure and gas sensing characteristics (viz., enhancement of sensitivity and shortening of response time) of some oxides such as WO3, MoO3 and TiO2 were studied. This paper describes the methodology and discusses the results obtained in the case of two model systems, viz., tungsten oxide (WO3) and titanium oxide (TiO2).
2. THEORETICAL RATIONALE The methodology adopted in this work stems from the simple fact that at a given temperature and standard pressure (ambient; 1atm.), the oxidation of a metal to its oxide or reduction of an oxide to metal or its suboxide occurs at a welldefined finite partial pressure of oxygen. If the two phases (metal/metal oxide or metal suboxide/metal oxide) are in equilibrium, the incumbent oxygen partial pressure is recognized as the equilibrium partial pressure. According to the Gibbs phase rule, on either side of this unique oxygen pressure, at a given temperature, one of the two coexisting phases must disappear. This is illustrated in the following by considering a hypothetical metal oxidation reaction:
x M (s ) + O2 ( g ) → MO x (s ) 2 For which K 1 =
a MOx aM . p
x 2 O2
And ∆G R (1) = − RT . ln K 1 where K 1 and
(1)
(2a)
(2b)
∆G R (1) are the reaction equilibrium constant and the standard Gibbs’ energy change, respectively, for
reaction (1). Equations (2a) and (2b) can be combined and simplified to give:
pO2( M
MOx )
⎛ 2.∆G R (1) = exp⎜⎜ ⎝ xRT
⎞ ⎟⎟ ⎠
(2c)
Equation (2c) gives the thermodynamic equilibrium oxygen partial pressure for the coexistence of a metal with its adjacent oxide at a given temperature in terms of the standard Gibbs energy change for reaction (1); in the present case it is also equal to the standard Gibbs energy of formation of the oxide, MOx. A plot of pO2 or ln pO2 vs. temperature, gives the contour of the path of the M/MOx coexistence. Using reliable thermodynamic data5-6, the loci of log 10
pO2 as
a function of temperature in the range 400-800°C, for various metal/metal oxide couples are plotted in Fig. 1. Obviously, in every case, at a
pO2 lower than this ‘curve of coexistence’, metal oxide would be reduced either to
corresponding metal or a suboxide. On the other hand, at
pO2 values above the curve, a metal (or its suboxide) would be
completely oxidized to the corresponding stable oxide. Such a thermodynamically feasible redox process can well be carried out in a reducing atmosphere of hydrogen or carbon monoxide followed by oxidation in static or dynamic air.
Fig. 1.
Temperature dependence of the equilibrium oxygen partial pressure in: 1. Mo/MoO3; 2. Sn/SnO2; 3. W/WO3; 4. Mo/MoO2; 5. W/WO2; 6. Zn/ZnO; 7. Ti/TiO2 and 8. Ti/TiO. The variation in
pO2
by changing the CO2/CO ratio in the range 10-5-105
at 450, 600 and 800°C is also shown (open triangles).
For example, a metal oxide of interest can be reduced by H2 or CO to a suboxide or the corresponding metal. If such a reduced surface is simply heated in air (static or dynamic; pO2 = 0.21 atm.), it leads to bulk oxidation, forming an oxide whose morphological features may or may not be very different from the starting material.
Contrary to that, if the reduced surface is exposed to a well-defined
pO2 that is only a few orders of magnitudes higher
than the theoretical value for the M/MO coexistence, interesting processes ensue. Since the prevailing oxygen potential is only slightly above that established by virtue of thermodynamic equilibrium between M and MOx or between MOx and MOy, this allows the formation and growth of new oxide surface on an atomic/ molecular level, under conditions of ‘oxygen deprivation’. Similarly, by exposing the oxide to a precisely controlled pO2 regime that is below theoretical line of the metal oxide stability, one can modulate the extent of reduction of the said oxide either to a suboxide or ultimately to the metal. In any event, it can be envisaged that such a pO2 manipulation will deplete oxygen in a manner so as to cause atomic or submolecular level chemical variations. Hence, upon exposure to an environment that is only slightly rich in oxygen, new material build-up takes place layer-by-layer, thereby creating whole new morphological features that are alien to the starting bulk oxide. Adopting this scheme, we have successfully created novel microstructures in a host of ceramic oxide systems with a view of imparting benign surface features that are paramount in accentuating their functional behavior in gas sensing applications which will be discussed subsequently. It is, however, worth pointing out that, in the light of Fig. 1, it appears thermodynamically impossible to cause reduction of TiO2 into either TiO or Ti metal via gas buffering technique (discussed in the next section), since the equilibrium dissociation oxygen partial pressure of TiO2 into a suboxide and/or Ti is very small and not achievable practically. On the other hand, substantial morphological architect could be developed by using a H2/N2 mixture in the case of titania.
3. METHOD OF ESTABLISHING THE pO2 DESIRED FOR THE REDOX REACTIONS As can be seen from Fig. 1, the equilibrium oxygen partial pressure in a typical metal oxide is rather low. An oxygen potential in the vicinity of these equilibria could potentially be generated by manipulating the ratio of two gaseous species in a buffer mixture, such as CO2/CO or H2O/H2. At high temperatures, CO and CO2 can exist in equilibrium with traces of oxygen.
1 CO + O2 = CO2 2
(3)
⎛ ⎜ pCO o 2 For which ∆G = − RT ln⎜ 1 2 ⎜⎜ ⋅ p p CO O 2 ⎝ This gives pO2
Where ∆G
0
⎛ pCO2 = ⎜⎜ ⎝ pCO
⎞ ⎟ ⎟ ⎟⎟ ⎠
(4a)
2
⎞ 1 ⎟ ⋅ − 2 ∆G ° ⎟ ⎠ e RT
(J ) = −282400 + 86.81T
(4b)
(4c)
Therefore, by controlling the ratio of the concentration of CO2 and CO, it is possible to control the partial pressure of oxygen. Mixing CO2 and CO in the ratio that ranges from 10-5 to 105 provides good buffered systems. In this range, the theoretical pO2 varies between 10-35 and 10-15 atm. at 600°C and between 10-29 and 10-9 at 800°C. These are represented as open triangles in Fig. 1. Similar pO2 can be generated by an equilibrium established in a mixture of H2/H2O. Such oxygen potentials could also be obtained with relative ease via establishment of the following equilibrium:
1 H 2 + O2 → H 2 O 2
(5)
⎛ ⎜ pH O o 2 For which ∆G = − RT ln⎜ 1 2 ⎜⎜ ⎝ p H 2 pO2
This gives
⎛p p = ⎜⎜ ⎝ p
H 2O
0
⎞ 1 ⎟. ⎟ ⎠ e
(6a)
2
O2
Where ∆G
⎞ ⎟ ⎟ ⎟⎟ ⎠
H2
− 2 ∆G 0 RT
(J ) = −239500 + 8.14T ln T − 9.25T
(6b)
(6c)
Thus by varying the water vapor to hydrogen ratio in the range of 10-5 to 105, the corresponding equilibrium oxygen partial pressure at 600°C can be conveniently varied between 10-34 and 10-14 atm. However, owing to the ease of mixing gaseous components and to eliminate the possibility of water condensation in the cooler section of the experimental setup, a buffer mixture of CO/CO2 was employed in this work.
4. EXPERIMENTAL PROCEDURE The materials investigated in this work included thin foils of W as well as the powders of WO3 and TiO2, all from AlfaAesar (99.8% or better). The W metal foil and WO3 powder were used to demonstrate the authenticity of the proposed concept. In the case of TiO2, sintered pellets (at 1200 °C for 6 h) were reduced in N2/H2 mixture at 700 °C/8h. In each of these cases, the products after redox reactions was also studied by X-ray diffraction (XRD), scanning and transmission electron microscopy (SEM/TEM) to corroborate the observed enhancement in sensing characteristics. The WO3-based thick film sensors in chemiresistor mode were fabricated as follows. The powders were first ball-milled using 10mm spherical zirconia milling media (Tosoh, NJ) in 2-propanol for 8h, dried and sieved through a 325-mesh stainless steel screen. Each powder was mixed with V-006 (an organic-based resinous vehicle with dispersant, from Heraeus, PA) and α-terpineol (Alfa-Aesar) in an appropriate weight ratio (~70% solid loading) and stirred well so as to form a uniform slurry of adequate rheology. In order to improve the adhesion of the film to the substrate, 2 wt% of tetraethoxysilane (TEOS, Alfa-Aesar) was also added to the slurry and homogenized prior to printing. The film was screen-printed on high density α-alumina substrates (14mm x 14mm) pre-fabricated with interdigitated gold electrodes (12mm x 12mm) and contact pads. The films were first dried in an air oven at 150°C followed by firing in the range of 500-900°C, depending upon the physico-chemical properties of the oxides (stability, volatility, phase transformation, etc.) for 1-2h in air. Gold lead wires (0.25 mm diameter, Alfa-Aesar) were attached to the contact pads via silver paste which was cured in three different stages between room temperature and 350°C so as to form good ohmic contacts. The sensor was placed on a flat platform in an all-quartz experimental set-up which was located in the uniform temperature zone of a compact horizontal Lindberg furnace (MiniMite) and the lead wires were taken out of the furnace through a twin-bore alumina tube. A type-K thermocouple was also placed just above the sensor to monitor the temperature and its variation (if any) during the test. The ends of the gold wires were connected to a high impedance Agilent 34220A digital multimeter, which in turn was connected to a desktop PC via HPIB interface card. Sensor resistance data was acquired and displayed in real-time with the help of IntuiLink software. A gas stream consisting of 10% O2-90% N2 (v/v) mixture was obtained by blending dry compressed air with high purity nitrogen to obtain the background (reference) gas. The sensor was first heated to a selected temperature in the
background ambient, allowed to equilibrate at that temperature till a steady baseline resistance (Rb) was established. Given amount of CO from a CO/N2 tank was then bled in and allowed to blend. Sensitivity of a given film was measured by recording change in film resistance with respect to Rb upon introduction of a given amount of CO in the stream. The sensor behavior was monitored both with increasing and decreasing level of CO in the ambient to confirm the reversibility attribute of the sensor. The response time (t90) was calculated by discerning from the recorded data, the time it took for the signal to attain 90% of the difference between the two steady states, viz., in the background (Rb) and that after CO was introduced (Rg). To fabricate a TiO2-based sensor device, gold lead wires were attached to the nanofiber-covered surface by painting two strips of conductive gold paste onto the surface. The gold paste was cured at 700°C for 2 h under a 5% H2/N2 atmosphere. Hydrogen sensing measurements were conducted in a background gas of 10% O2 in a balance N2. A constant hydrogen gas flow rate through the test tube was maintained with a mass flow controller (MKS instruments, Austin, TX). The change in the electrical resistance of the sensor upon exposure to different hydrogen concentrations was measured with a HP digital multimeter and recorded with a computer. Details can be found in ref. 2.
5. RESULTS AND DISCUSSION 5.1 W and WO3-based Studies The above-described hypothesis pertaining to the variation in the morphological artifacts in materials when they are subjected to redox reactions in precisely controlled oxygen potential regimes was first verified in the case of pure W metal. Small (25 mm x 12.5 mm) coupons cut from a tungsten foil were subjected to several redox schemes and were characterized at the end of each treatment by XRD and SEM. Figure 2 compares the XRD signatures of the as-received W foil and the foil subjected to oxidation ( pO2 ~10-15) and reduction ( pO2 ~10-29) by CO/CO2 buffers at 800°C for 24h. It can be readily inferred that elemental W is regenerated after oxidation and reduction in CO/CO2 atmospheres7. The comparative morphological features of these samples are shown in Fig. 3.
Fig. 2.
XRD patterns of pure W foil and that subjected to redox reaction by the CO/CO2 buffer.
a
b
Fig. 3. Electron micrographs of: (a) pristine W foil and (b) the foil oxidized and reduced by a CO/CO2 mixture at 800°C/24h.
The fractured morphology of the surface in Fig. 3b clearly indicates that the material has undergone major phase and structural (metal (cubic) → oxide (triclinic)) changes prior to re-conversion to the metal again. This is corroborated by the EDS spectra collected after each event and is shown in Fig. 4. Fig. 4.
EDS spectra of W foil subjected to various oxidation and reduction treatments (see the text).
As shown in Fig. 5, the XRD signature of a W foil directly oxidized in air at 800°C/2h is identical with that of the foil subjected to an oxidationreduction cycle in CO/CO2 stream at 800°C/36h, followed by oxidation in air. Phase analysis of the two patterns shows that in both cases, WO3 is the dominant phase. This is understandable, since despite the intermediate heat-treatments in low pO2 regimes, the phase evolution is ultimately dictated by the final processing parameters and the ambient conditions which is the same in the two cases.
Fig. 5.
XRD pattern of (a): W foil after bulk oxidation in air at 800°C/2h, and (b): that oxidized and reduced in CO/CO2 mixtures at 800°C/36h followed by air oxidation at 800°C/2h.
Accordingly, as shown in Fig. 6 the microstructural features of the two samples are also similar albeit with some grain growth and compaction clearly visible in the second case, probably due to several heat-treatments leading to some sintering as well.
a
Fig. 6.
b
Microstructural features of W foil after (a) heating in air for 2h at 800°C, and (b) oxidized and reduced for 36h in CO/CO2 mixtures followed by air oxidation for 2h at 800°C.
In contrast to this, the microstructural features seem to undergo drastic changes, when the W foil is oxidized in CO/CO2 mixture in one case and oxidized, reduced and re-oxidized in manipulative pO2 regimes generated by different CO/CO2 ratios. This is shown in Fig. 7.
a
Fig. 7.
b
Scanning electron micrographs of W foil after: (a) single-stage oxidation and, (b) cyclic oxidation-reduction-oxidation by CO/CO2 buffer mixtures at 800°C/24h.
The microstructure shown in Fig. 7a results when pure W foil (Fig. 3a) is heated at 800°C/24h in CO/CO2 mixture in a pO2 range that is above the W/WO3 line in Fig. 1, while the morphology shown in Fig. 7b evolves when the foil in Fig. 3b is subjected to an identical treatment. A comparison of these features with those shown in Fig. 6, lends credibility to the proposed notion that depending upon the location of the equilibrium oxygen potential across the M/MO line of coexistence, changes on microscopic levels are caused in the bulk oxides. This perhaps leads to atomic/submolecular level non-stoichiometry – that is nanoscale in nature and hence undetected by bulk techniques such as XRD. It is these defect sites that act as the nucleation and growth centers for the new oxide phase whose growth, due to the limited access to oxygen (defined by a pO2 that is designed to be only slightly above the theoretical line of M/MO coexistence), is slow. Evidently, the reduced phase is subject to a concentration strain in terms of pO2 . The need for 24h long dwell at 800°C for oxidation in the CO/CO2
buffer compared to 2h for air oxidation supports the diffusion-like controlled buildup of the new oxide phase almost on atomic scale - one monolayer after another. Indeed, in many cases where bulk oxides were subjected to redox treatment by CO/CO2 mixtures for shorter duration, it was rather difficult to discern quantitative phase changes based simply on the XRD signatures. For example, when a WO3 thick film is subjected to the reducing and oxidizing sequence in appropriate oxygen potential regimes created by CO/CO2 buffers in one case, and another film after the above-mentioned treatment is finally heated in air, no difference in their XRD signatures could be discerned, as shown in Fig. 8. All the three patterns could be indexed as those belonging to triclinic WO37. Fig. 8.
Comparative XRD patterns of: (a) WO3 film calcined in air at 800°C/2h, (b) a, subjected to redox in CO/CO2 mixtures at 800°C/12h and (c) b, heated in air at 800°C/2h.
On the contrary, the morphological features are significantly affected – both in terms of shape and size of the particles as well as in terms of the uniform porosity in the sample, as seen from the SEM pictures in Fig. 9.
a
b
c
Fig. 9.
d
Scanning electron micrographs of WO3 thick films: (a) as-prepared, (b) a reduced in CO/CO2, (c) b oxidized in CO/CO2 and (d) b oxidized in air at 800°C, for various periods of time.
This vividly illustrates that the ambient oxygen potential in the vicinity of an oxide phase has profound effect on its morphological features, which could be tailored to accentuate the sensing behavior of a potential semiconducting oxide. The response characteristics (such as resistance, sensitivity and response time) of WO3 thick film sensors to CO gas (14100 ppm) at 450°C are shown in Fig. 10.
Fig. 10. Response of a WO3 thick film sensor to CO at 450°C. a: as-prepared, b: a reduced and oxidized in CO/CO2 mixtures, and c: a reduced in CO/CO2 mixture and oxidized in air.
As can be seen from Fig. 10, surface modification by the proposed scheme has certainly brought about marked changes in the behavior of a WO3-based CO sensor. The sensor could be operated in a cyclic fashion without compromising the signal, with a concomitant enhancement in sensitivity and significant shortening of the response time. In order to examine if such exotic microstructural features could be developed in bulk oxides, WO3 thick films formed by heating the slurry at 600°C/1h in air were subjected to reduction in a H2/N2 mixture at 600°C for ½ h, cooled to room temperature and again heated in air up to 500°C for ½ h. The evolved microstructures are shown in Fig. 11. In such cases also, the regeneration of the parent oxide with different morphological features can be explained in a way similar to that summarized above. The only notable difference is that in the later case, the oxygen potential in the ambient is rather high (0.21 atm.) – a parameter which could kinetically favor the process, leading to the faster attainment of thermodynamically most stable (M/O) stoichiometry upon regeneration.
a
Fig. 11. Evidence of microstructural modification in bulk oxides (a) via H2 reduction-air oxidation (b); b at higher magnification is shown in c.
Similar morphological changes have been observed in the case of sensors made with composite mixtures, viz., ZnMoO4-MoO3 (MZM). It was found that the rod-like MoO3 grains in the original mixture are regenerated as highly oriented thin platelets upon exposing the MZM film to a gas mixture containing 1%CO at 450°C for 1h followed by natural cooling in air. It was also observed that as
b
c
a result of this bulk redox reaction, the morphological features of the major phase (ZnMoO4) have undergone noticeable variation (from regular near spherical grains to triclinic habits with well-defined sharp edges) without any chemical degradation8-9. These results are shown in the SEM pictures (Fig. 12) collected in different pockets of the composite after subjecting it to the above-mentioned redox treatment. The morphology of regenerated ZnMoO4 grains in Fig. 12c is isostructural with that of the mineral microcline (KAlSi3O8) that belongs to the triclinic pinacoidal class of crystals10.
a
b
c
d
Fig. 12. Microstructural features of the MZM composite sensor film: a. as-prepared phase pure ZnMoO4 (ZM), b. as-prepared MZM composite; morphological changes in ZnMoO4 and MoO3 phases after the composite was exposed to 1%CO at 450°C/1h following by natural cooling in air, are shown in c and d, respectively.
5.2 TiO2-based Studies Figure 13 shows scanning electron microscope (SEM) image of the surface of a TiO2 sample before and after the nanocarving. As can be seen, TiO2 grains converted into aligned nanofibers, which had a diameter of around 20 nm and lengths of 1 µm. Before nanocarving (Figure 13a), polycrystalline TiO2 disk samples were prepared by sintering powder compact at 1200 °C for 6 h. The nanofibers were only formed on the surface and the depth of nanofiber formed region ranged from 0.7 to 1 µm. The phase of nanofibers was determined to be rutile TiO2 by XRD, TEM11 and XPS2, although the nanofibers were generated under a reducing environment of H2/N2. Based on a time sequence SEM experiment11 of a specific grain, it was determined that the nanofiber formation was by an etching process and not by a growth process. In addition, thermogravimetric analysis (TGA)12 showed that the weight of TiO2 decreased during H2/N2 treatment, which confirmed that the nanofiber formation was due to an etching process. The etching direction was determined to be from TEM observation11. This etching was selective and anisotropic, which led to aligned arrays of nanofibers. This crystallographic etching process was dubbed “Nanocarving”. To determine the role of nitrogen gas, if any, TiO2 specimen were heat-treated in 100% N2 atmosphere under similar conditions as the H2/N2 nanocarving process (700 °C, 8 hrs) and no evidence of etching was evident. To confirm that only H2 gas plays a role on the nanocarving process, N2
was switched with another inert gas, Ar. In 5% H2/ 95% Ar atmosphere, nanofibers were observed on the surface after the nanocarving process. These results indicate that only hydrogen is the reacting gas species. The chemical reaction of H2 with TiO2 can be represented as: TiO2 + xH 2 ( g ) = TiO2− x + xH 2 O( g ) From TGA, weight loss of TiO2 was observed during the H2/N2 treatment and it follows the parabolic rate law13. From mass spectroscopy and ICP analysis of the reaction tube, no volatile Ti-species was found13. Therefore, during H2/N2 treatment, only O was removed from the surface. Since surface stoichiometry did not change and there was large volume change, the excess Ti cations should diffuse from the surface to the bulk to maintain the stoichiometry.
(a)
(b)
Fig. 13. Scanning electron microscope (SEM) images of TiO2 surface (a) before and (b) after nanocarving.
In Figure 14a, only half of the grain transformed into nanofibers and the region where no fiber was generated had nanoparticles with the diameter of few tenths of nanometers. In addition, the end tips of nanofibers were round and
possessed similar diameters as the nanoparticles, i.e. the nanoparticles appeared to be sitting on the end tips of nanofibers. Previous TEM and EDS results showed that the nanoparticles were enriched in Fe and Ni11. In addition, the nanoparticles were likely to be metallic alloys since the O peak was not observed in the EDS pattern. These nanoparticles appear to have been generated during the H2/N2 gas treatment by the external reduction of the Fe and Ni oxides present as impurities in the starting TiO2 powder. One of the potential applications of the nanofiber arrays is in chemical sensors. Traditionally, metal oxides such as SnO2 and TiO2 have been used for sensors for reducing gas species such as CO and H2. SnO2 sensors utilize low-temperature chemisorption (under 400 °C) of environmental gases on the surface, which results in electron exchange leading to a change in the resistance. However, over 400 °C, SnO2 exhibits poor sensing performance. On the other hand, TiO2 is stable at higher temperatures but the sensitivity of TiO2 sensors is usually not as good as SnO2 sensors. Recently, several attempts to combine the advantages by mixing TiO2 and SnO2 have been reported14-16. In the present study, two TiO2SnO2 mixed phases, solid solution (SS) and spinodal decomposition (SD) were produced and subjected to the H2 heat treatment process to produce nanofibers. The H2 sensing behavior of the TiO2-SnO2 nanofibers was also investigated and was compared with that of pure TiO2 nanofibers17.
(a)
(b)
Fig. 14. SEM micrographs of a TiO2 disk surface after H2/N2 heat treatment, showing nanofibers with nanoparticles; (a) low magnification and (b) high magnification of the area enclosed by a square in (a).
Figure 15 shows the sensitivity of 4 samples (pure TiO2, 2 h and 6 h SS, and 6 h SD TiO2-SnO2) with changing H2 concentration. While the 2 h SS sample practically shows no sensitivity, pure TiO2 and the 6 h SD TiO2-SnO2 show sensitivity higher than the 6 h SS TiO2-SnO2 sample. From SEM micrographs it is evident that pure TiO2 has welldeveloped nanofibers and SD TiO2-SnO2 has grooves over the entire surface. On the contrary, 6 hr SS TiO2-SnO2 sample shows few nanofibers on the surface and 2 hr SS shows no nanofibers on the surface. Thus, the higher sensitivity of pure TiO2 and SD TiO2-SnO2 sensors are most likely due to the increased surface area produced by the nanocarving process. Although the SD TiO2-SnO2 have lower sensitivity to H2 than pure TiO2, in initial tests it was observed that the SD samples were more stable than pure TiO2 over extended testing periods (1 month) and in oxygen atmospheres between 400 and 700oC.
Fig. 15. The H2 gas sensitivity of four samples (pure TiO2, 2 h and 6 h SS, and 6 h SD TiO2-SnO2) highlighting microstructure-sensitivity correlation.
6. CONCLUSIONS The following conclusions can be drawn: 1. 2. 3.
Novel microstructural features could be incorporated in a given semiconducting oxide via precise oxygen potential modulation and hydrogen containing gas-phase etching. A buffer gas mixture (CO/CO2 or H2/H2O) of appropriate ratio provides a convenient environment for the proposed redox scheme; from the point of view of experimental ease, CO/CO2 is a better buffer mixture. The formation and growth of new oxide surface on an atomic or submolecular level, under conditions of oxygen deprivation appears to be the most likely pathway.
4. 5. 6.
Temperature-time-pO2 correlation is a benign microstructural determinant for oxide ceramics that are amenable to redox reactions. Nano-carving in TiO2 is achieved through etching by hydrogen containing species. While oxygen from TiO2 is taken out as H2O (g), Ti diffuses from the surface to the bulk. The regenerated oxide phase with unusual microstructure shows enhanced gas sensing behavior.
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