Ceramic Product Development and Commercializationaazad/pdf/IJAC.pdfCeramic Product Development and Commercialization *[email protected] r 2006 The American Ceramic Society mechanism

Published by Blackwell Publishing for the American Ceramic Society

TheAmericanCeramicSocietywww.ceramics.org

Ceramic Product Development and Commercialization

International Journal of Applied C

eramic Technology

IJAC

CP - Volum

e 3 · Issue 3 · pp. 177–258 2006

Topical Focus:Advanced Ceramics for Sensor Applications

Novel Structural Modulation in Ceramic Sensors ViaRedox Processing in Gas Buffers

Abdul-Majeed Azad*

Department of Chemical and Environmental Engineering, The University of Toledo, Toledo, Ohio43606-3390

Sheikh A. Akbar

Department of Materials Science and Engineering, The Ohio State University, Columbus, Ohio 43210

High selectivity, enhanced sensitivity, short response time, and long shelf-life are some of the key features sought in thesolid-state ceramic-based chemical sensors. As the sensing mechanism and catalytic activity are predominantly surface-dom-inated, benign surface features in terms of small grain size, large surface area, high aspect ratio and, open and connectedporosity, are required to realize a successful material. In order to incorporate these morphological features, a technique basedon rigorous thermodynamic consideration of the metal/metal oxide coexistence is described. By modulating the oxygen partialpressure across the equilibrium M/MO proximity line, formation and growth of new oxide surface on an atomic/submolecularlevel under conditions of ‘‘oxygen deprivation,’’ with exotic morphological features, has been achieved in potential sensormaterials. This paper describes the methodology and discusses the results obtained in the case of potential semiconductingceramic oxide-based carbon monoxide and hydrogen sensors with enhanced characteristics.

Introduction

One aspect of current interest and great relevance tothe fundamental understanding of the behavior of ma-terials is the role of dimensionality and size on theiroptical, chemical, and mechanical properties for appli-cation in a wide range of devices. Owing to the nano-scale features, one-dimensional systems exhibit novelphysical and chemical properties that can be exploited inoptics, catalysis, and data storage devices. Hence, such

systems are being synthesized and studied in great detail.There is growing interest in introducing such attributesin nanoscale inorganic materials as well. The mostobvious advantage of doing so is the possibility of theirapplication as quantum dots in a host of devices, such asMEMs, lab-on-a-chip sensors/detectors, micro-solar cellelectrodes, micro-fuel cells, photocatalysts, and electro-catalysts, to name a few.

One area where nanofeatures in the materials are ofimmense relevance is the field of solid-state ceramic-based chemical sensors. High selectivity, enhancedsensitivity, and short response time are some of thekey features sought in these devices. As the sensing

Int. J. Appl. Ceram. Technol., 3 [3] 177–192 (2006)

Ceramic Product Development and Commercialization

*[email protected]

r 2006 The American Ceramic Society

mechanism and catalytic activity of ceramics are largelymicrostructure dominated, benign surface features suchas small grain size, large surface area, high aspect ratio,and open/connected porosity are required to realizea successful sensor material.1–4

A novel technique used to impart such attributes bymodifying the microstructural artifacts of ceramic-basedsensor materials is described in this paper. The effects ofthe variation in the ambient oxygen partial pressureacross the metal/metal oxide boundary on the micro-structure and gas-sensing characteristics (viz., enhance-ment of sensitivity and shortening of response time) ofpotential oxides such as WO3, MoO3, and TiO2 werestudied. The outlines of the methodology adopted tobring about the morphological transformation in thesematerials have been recently reported elsewhere;2,5–6

some aspects of the accentuated sensing attributes ofstructurally modified sensor films of tungsten oxide andtitanium oxide were also discussed. This paper discussesthe results pertaining to the benign effect of the micro-structural features on the sensing behavior of molybdenumand titanium oxide toward CO and H2, respectively, inan air-like ambient at high temperatures.

The methodology adopted to bring about the de-sired morphological variations can be briefly stated asfollows: at a given temperature and standard pressure(ambient; 1 atm.), the oxidation of a metal to its oxideor reduction of an oxide to a metal or its suboxide oc-curs at a well-defined finite partial pressure of oxygen. Ifthe two phases (metal/metal oxide or metal suboxide/metal oxide) are in equilibrium, the incumbent oxygenpartial pressure is recognized as the equilibrium partialpressure. According to the Gibbs phase rule, on eitherside of this unique oxygen pressure, at a given temper-ature, one of the two coexisting phases must disappear.This is illustrated in the following by considering a hy-pothetical metal oxidation reaction:

MðsÞ þ x

2O2ðgÞ !MOxðsÞ ð1Þ

for which

K1 ¼aMOx

aM :px=2O2

ð2aÞ

and

DGRð1Þ ¼ �RT lnK1 ð2bÞ

where K1 and DGR(1) is the reaction equilibrium con-stant and the standard Gibbs’ energy change, respec-tively, for reaction (1).

Equations (2a) and (2b) can be combined and sim-plified to give

pO2ðM=MOxÞ ¼ exp2:DGRð1Þ

xRT

� �ð2cÞ

Equation (2c) gives the thermodynamic equilibrium ox-ygen partial pressure for the coexistence of a metal withits adjacent oxide at a given temperature in terms of thestandard Gibbs energy change for reaction (1); in thepresent case, it is also equal to the standard Gibbs energyof formation of the oxide, MOx. A plot of pO2

or ln pO2

versus temperature gives the contour of the path of theM/MOx coexistence.

At a pO2lower than the ‘‘line of coexistence,’’ metal

oxide would be reduced either to a corresponding metalor a suboxide. On the other hand, at pO2

values abovethe line, a metal (or its suboxide) would be completelyoxidized to the corresponding stable oxide. Such a ther-modynamically feasible redox process can well be car-ried out in a reducing atmosphere of hydrogen orcarbon monoxide, followed by oxidation in static ordynamic air. For example, a metal oxide of interest canalso be reduced by H2 or CO to a suboxide or the cor-responding metal. If such a reduced surface is simplyheated in air (static or dynamic; pO2

¼ 0:21atm.), itleads to bulk oxidation, regenerating the parent phasewhose morphological features may or may not be verydifferent from the starting material.

Contrary to this, if the reduced surface is exposed toa well-defined pO2

that is only a few orders of magnitu-des higher than the theoretical value for the M/MO co-existence, interesting processes ensue. As the prevailingoxygen potential is only slightly above that establishedby virtue of thermodynamic equilibrium between Mand MOx or between MOx and MOy, this allows theformation and growth of new oxide surface on an atom-ic/molecular level, under conditions of limited oxygenaccess. Similarly, by exposing the oxide to a preciselycontrolled pO2

regime that is below the theoretical lineof the metal oxide stability, one can modulate the extentof reduction of the said oxide either to a suboxide orultimately to the metal. In any event, it can be envisagedthat such a pO2

manipulation will deplete oxygen in amanner so as to cause atomic or submolecular-levelchemical variations. Hence, upon exposure to an envi-ronment that is only slightly rich in oxygen, new ma-terial build-up takes place layer by layer, therebycreating whole new morphological features that are al-ien to the starting bulk oxide.

178 International Journal of Applied Ceramic Technology—Azad and Akbar Vol. 3, No. 3, 2006

Using reliable Gibbs energy data,7,8 the loci oflog10 pO2

as a function of temperature in the range400–8001C, for various metal/metal oxide couples, areplotted in Fig. 1a. It is, however, worth pointing outthat in the light of Fig. 1a, it appears thermodynami-cally impossible to bring about reduction of TiO2 toeither TiO or Ti metal via the buffer gas technique. Thisis because of the fact that the thermodynamic stability oftitania is quite high and the corresponding dissociationoxygen partial pressure in the Ti/TiO2 system israther low. No ‘‘practical’’ mixture of CO and CO2

could cause reduction of TiO2 to TiO or Ti. Forexample, the equilibrium oxygen partial pressure inthe Ti/TiO2 coexisting mixture ranges from 2.732�10�46 atm. at 873 K to 4.192� 10�29 atm. at 1273 K.In the case of Ti/TiO, the equilibrium pO2

values rangefrom 1.18� 10�52 atm. at 873 K to 2.06� 10�33 atm.at 1273 K. Calculations show that no ‘‘practical’’ mix-ture of CO and CO2 gases would be able to establishsuch low pO2

values. Hence, it is impossible to bringabout a redox reaction of TiO2 by this scheme. Thus, inthis case, microstructural modification by hydrogentreatment was resorted to and significant morphologi-cal architect could be developed by using a H2/N2 mix-ture in the case of TiO2.

In the context of molybdenum oxide-based sensors,a plot of log pO2

versus T for the Mo/MoO3 coexistencein the range 450–8001C is shown in Fig. 1b (solidcurve). An oxygen partial pressure in the vicinity ofMo/MoO3 proximity line could be generated bymanipulating the ratio of two gaseous species in abuffer mixture, such as CO2/CO or H2O/H2. In thepresent case, as was in the case with tungsten oxide re-ported previously,9 a buffer mixture of CO2/CO wasused to create adequate oxygen partial pressure. There-fore, by controlling the ratio of the concentration ofCO2 and CO, it is possible to control the partial pres-sure of oxygen. Mixing CO2 and CO in the ratio thatranges from 10�5 to 105 provides good buffered sys-tems. In this range, the theoretical pO2

varies between10�35 and 10�15 atm. at 6001C and between 10�29 and10�9 at 8001C. These are represented as open trianglesin Fig. 1a.

Adopting this scheme, we have successfully creatednovel microstructures in a host of ceramic oxide systemswith the aim of imparting benign surface features thatare paramount in accentuating their functional behaviorin gas-sensing applications, which will be discussedsubsequently.

Experimental Procedure

Materials, Processing, and Characterization

The materials investigated in this work includedthin foils of Mo and W as well as the powders of MoO3

and WO3 (Alfa Aesar, Ward Hill, MA; 99.8% or bet-ter). The vendor-specified average particle size of theoxides was between 20 and 45 mm. The metal foils wereused to demonstrate the authenticity of the proposedconcept while the oxide powders were used to fabricatethe sensor films whose sensing behavior toward differentlevels of carbon monoxide or hydrogen in a 10% O2-bal. N2 background was monitored before and after

Fig. 1. (a) Temperature dependence of the equilibrium oxygenpartial pressure in: 1, Mo/MoO3; 2, W/WO3; 3, Mo/MoO2; 4, W/WO2; 5, Zn/ZnO; 6, Ti/TiO2; and 7, Ti/TiO. The variation inpO2

by changing the CO2/CO ratio in the range of 10�5–105 at4501C, 6001C, and 8001C is also shown (open triangles).(b) Temperature dependence of the equilibrium oxygen partialpressure in the Mo/MoO3 system (solid line). The variation in pO2

by changing the CO2/CO ratio between 10�2 and 105 at 4501C,6001C, and 8001C is also shown.

www.ceramics.org/ACT Novel Structural Modulation in Ceramic Sensors 179

these films were subjected to the reduction–oxidationprocesses described above. A set of these films were alsoreduced in hydrogen and oxidized in air to differentiatethe morphological attributes obtained in the proposedapproach versus those resulting from the classical redoxmethod. Titania samples were subjected only to reduc-tion in a H2/N2 mixture.

The stoichiometry of MoO3 is known to be signif-icantly dependent upon the method of preparationused.10,11 To discern the effect of the synthesis meth-odology on the microstructural evolution, molybdenumtrioxide was synthesized by two additional methods inaddition to using the as-received powder from Alfa-Aesar. In the first case, commercial MoO3 powder wasused without further processing, except that it was thor-oughly ground, ball-milled, and sieved through a 325mesh screen before printing thick films. In the secondcase, the MoO3 powder was obtained by oxidizing a Mometal foil in air at 6001C for 2 h. In the third instance,MoO3 was prepared by a soft solution chemistry routeusing sodium molybdate dihydrate (Na2MoO4. � 2H2O,AR grade from Alfa-Aesar) as the precursor. 10.5 g ofsodium molybdate crystals were dissolved in 150 mL ofde-ionized water with constant stirring and the solutionwas cooled in an ice-bath to 51C. Twenty-five millilitersof 1M HCl acid was added to the cooled aqueous so-lution, stirred, and the mixture was placed in the freezerovernight so as to maintain the temperature around51C. The frozen solution was thawed and heated slowlyto 801C with continuous stirring for 24 h; a pale bluishprecipitate of molybdic acid (H2MoO4) started to formafter heating. The precipitate was washed several timeswith de-ionized water and centrifuged. In order to re-move all the sodium ions, the conductivity of the filtratewas monitored after each wash until the conductivitywas o0.05O�1.cm�1. The solid thus obtained wasdried, ground well, and calcined for 1 h in air at3001C. The formation of MoO3 from the precipitatedmolybdic acid (H2MoO4) was confirmed by followingthe systematic phase evolution via X-ray diffraction(XRD) and scanning electron microscopy (SEM).

For titania-based systems, the starting material wascommercial anatase TiO2 powder (Alfa Aesar, 99.9%pure) having an average particle size of 32 nm. Batchesof 0.3 g of TiO2 powder were compacted into disksunder a uniaxial pressure of 17.5 MPa in a stainless-steeldie (Carver Inc., Wabash, IN) by single-end compact-ion. The TiO2 disks were placed in a furnace (Lindberg/BlueM, Asheville, NC) and sintered in air using a

51C/min heating rate to a set temperature of 1100–14001C, followed by 0–48 h hold times. The sintereddisks were 0.8 mm thick and 10 mm in diameter. Thesintered TiO2 disks were placed in a horizontal tubefurnace (Lindberg/BlueM) and exposed to a flowing at-mosphere of 5% H2/95% N2 at 600–7801C for 8 h.The ends of the fused silica tube were secured by endcaps with o-rings to ensure a gas-tight seal. Oxygen par-tial pressure measured with an oxygen sensor (CG 1000,Ametek, Paoli, PA) was 3.0–9.1� 10�17 Pa and watervapor partial pressure measured with a humidity sensor(HMP234, Vaisala, Woburn, MA) was 27.5–68.8 Pa.The flow rate of H2/N2 gas was over 500 sccm.

Structural and microstructural examination of theas-received foils, raw powders and the products aftereach of the redox reactions was conducted by XRD,SEM/TEM to corroborate the observed enhancement insensing characteristics of the thick film sensors.

Sensor Fabrication and Testing

The thick film sensors in chemiresistor mode werefabricated as follows: the powders were first ball-milledusing 10 mm spherical zirconia milling media (Tosoh,East Hanover, NJ) in 2-propanol for 8 h, dried, and si-eved through a 325-mesh stainless-steel screen. Each pow-der was mixed with V-006 (an organic-based resin vehiclewith a dispersant, from Heraeus, Conshohocken, PA) anda-terpineol (Alfa-Aesar) in an appropriate weight ratio(B70% solid loading) and stirred well so as to form auniform slurry of adequate rheology. In order to improvethe adhesion of the film to the substrate, 2 wt% of tet-raethoxysilane (TEOS, Alfa-Aesar) was also added to theslurry and homogenized before printing. The film wasscreen printed on high-density a-alumina substrates(14 mm� 14 mm) prefabricated with interdigitatedgold electrodes (12 mm� 12 mm) and contact pads.The films were first dried in an air oven at 1501C, fol-lowed by firing in air at temperatures commensurate withthe physico-chemical properties of the oxides, such as,stability, volatility, phase transformation, etc.) for 1–2 h.Gold lead wires (0.25 mm diameter, Alfa-Aesar) were at-tached to the contact pads via a silver paste that was curedin three different stages between room temperature and3501C so as to form good ohmic contacts.

The sensor was placed on a flat platform in an all-quartz experimental setup that was located in the uni-form temperature zone of a compact horizontal Lindb-erg furnace (MiniMite) and the lead wires were taken


out of the furnace through a twin-bore alumina tube. Atype-K thermocouple was also placed just above thesensor to monitor the temperature and its variation (ifany) during the test. The ends of the gold wires wereconnected to a high impedance Agilent 34220A digitalmultimeter (Agilent Technologies Inc., Palo Alto, CA),which in turn was connected to a desktop PC via HPIBinterface card. Sensor resistance data were acquired anddisplayed in real time with the help of IntuiLink soft-ware.

A gas stream consisting of a 10% O2–90% N2 (v/v)mixture was obtained by blending dry compressed airwith high-purity nitrogen to obtain the background(reference) gas. The sensor was first heated to a select-ed temperature in the background ambient, and allowedto equilibrate at that temperature till a steady baselineresistance (Rb) was established. A given amount of COfrom a CO/N2 tank was then bled in and allowed toblend. The sensitivity of a given film was measured byrecording change in film resistance with respect to Rb

upon introduction of a given amount of CO in thestream. The sensor behavior was monitored both withincreasing and decreasing level of CO in the ambient toconfirm the reversibility of the sensor. The responsetime (t90) was calculated by discerning from the record-ed data the time it took for the signal to attain 90% ofthe difference between the two steady states, viz., in thebackground (Rb) and that after CO was introduced (Rg).

The thick film sensor prepared by a procedure de-scribed above was subjected to reduction and oxidationat pO2

below and above the theoretical line of coexist-

ence of the Mo/MoO3. Such oxygen partial pressureswere created by mixing CO and CO2 in appropriateconcentration from compressed gas cylinders as per Eq.(4a) so that in one case it brings about reduction ofMoO3 to Mo metal and in the other case re-oxidation ofmetallic Mo to MoO3 again. As explained in the nextsection, for example, at 6001C, the oxygen partial pres-sure for the Mo/MoO3 coexistence can be computed tobe 1.2� 10�21 atm. (using reliable thermodynamic dataused in the construction of Fig. 1). However, MoO2 isalso a stable phase in the Mo–O system under differentoxygen partial pressure conditions (1.2� 10�26 atm. at6001C). Thus, a pO2

nearly 5 orders of magnitude lower(B10�30 atm.) than the theoretical value for the Mo/MoO2 was considered adequate to reduce MoO3 toelemental Mo effectively and completely. This was gen-erated by mixing CO and CO2 in the ratio CO2:CO 5

0.00231 (pO2¼ 1� 10�30 atm.) This composition was

used as the reducing mixture. Similarly, a pO2nearly 5

orders of magnitude higher than the theoretical wasgenerated by mixing CO and CO2 in the ratioCO2:CO � 23,000 (pO2

¼ 1� 10�16 atm.). The actu-al mixing ratios used in this work turned out to be 0.2%CO2199.8% CO in the first case and, 43 ppm CO1

balance CO2 in the second case. In order to generate thelater mixture, a gas cylinder containing 100 ppmCO1balance CO2 was used as the CO source.

In the case of TiO2-based sensors, gold lead wireswere attached to the nanofiber-covered surface by paint-ing two strips of conductive gold paste onto the surface.The gold paste was cured at 7001C for 2 h under a 5%

Fig. 2. Microstructural features of a pure Mo foil (L) and that heated in air at 6001C/2 h (R). Extreme fracture and deformation of the foilupon heating in air is because of the formation of MoO3 from metallic Mo, which is accompanied by an B 6.5-fold increase in volume (seetext).


H2/N2 atmosphere. Hydrogen-sensing measurementswere conducted at 4001C in a background gas of 10%O2-bal. N2. Other details can be found in Yoo et al. andKarney et al.12,13

Results and Discussion

The results on the variation in the morphologicalfeatures of tungsten oxide films with change in oxygenpotential across the metal/metal oxide proximity linehave been recently reported6,9 and are not includedhere.

Mo- and MoO3-based Studies

Microstructural Evolution in MoO3 Derived from Mo FoilOxidation: In order to first verify whether modulationof oxygen potential in the vicinity of a metal or its oxide

affects the morphological artifacts, small (25 mm�12.5 mm) coupons cut from a Mo foil were subjectedto several redox schemes and were characterized at theend of each treatment by XRD and SEM.

The comparative morphological features of the as-re-ceived Mo foil and that subjected to oxidation in air at6001C for 2 h are shown in Fig. 2. The mechanicalstress causing the disintegration of the foil upon oxida-tion is because of the phase and structural change fromcubic Mo metal (ICDD 42-1120, unit cell volume31.17 A3) to orthorhombic MoO3 (ICDD 05-0508,unit cell volume 202.99 A3), which is attended by anB85% volume change or an B6.5-fold volume in-crease.

On the other hand, totally different microstructureresults when the Mo foil is oxidized and reduced at6001C for 24 h in manipulative pO2

regimes (B5 ordersof magnitudes across the Mo/MoO3 line in Fig. 1) gen-

Fig. 3. Scanning electron micrographs of Mo foil after (a)oxidation and (b) reduction by CO/CO2 buffer mixtures at 6001C/24 h.

Fig. 4. SEM pictures of (a) Mo foil after redox treatment in CO/CO2 buffer for 24 h followed by air oxidation for 4 h and (b)commercial MoO3 powder.


erated by different ratios of CO/CO2. This is shown inFigs. 3a and b, respectively. The nanoscale features ofthe oxides subsequent to both oxidation and reductiontreatment in CO/CO2 mixture can be clearly seen. Thelack of open porosity and the presence of well-definedgrain boundaries in both the cases are probably becauseof sintering via vapor phase transport (VPT) as a resultof long soak duration (24 h at 6001C).

Interestingly, when the foil oxidized and reduced inthe CO/CO2 buffer mixture (Fig. 3) is heated again inair at 6001C for 4 h, thin oriented platelets of MoO3 oflarge aspect ratios and surface area result. The resultingmorphology is quite similar to that seen in the com-mercial MoO3 powders, albeit with a much smaller andmore uniform grain size and a narrower particle sizedistribution. This is shown in Fig. 4.

The microstructural features developed in films fab-ricated from the powder obtained after Mo foil oxida-

tion in air at 6001C for 2 h foil are compared with thosein the films subjected to redox treatments in appropriatepO2

regimes at 6001C in Fig. 5.

Microstructural Evolution in MoO3-Thick Films Fabricatedfrom the Commercial Powder: The morphological fea-tures in the MoO3 thick films that were prepared byusing the commercial oxide from Alfa-Aesar, and sub-jected to different redox treatments at 6001C, are shownin Fig. 6.

It is apparent that while the grains maintain the uniqueplatelet texture, so characteristic of MoO3, the microstruc-ture lacks the uniformity that was obtained as a conse-quence of identical treatment given to the MoO3 powderderived from foil oxidation (nearly monosized small plate-lets with a narrow plate size distribution; see Fig. 5).

The nearly identical morpohological features after sub-jecting molybdenum oxide films to different redox treat-

Fig. 5. Microstructural evolution in the MoO3 films fabricated from Mo foil oxidation. (a) Film calcined for 1 h. (b) ‘a’ reduced by CO/CO2

for 12 h. (c) ‘b’ oxidized by CO/CO2 for 12 h. (d) ‘c’ oxidized by air for 1 h. Size bar 5 5mm.


ments under various oxygen partial pressure regimes sug-gest that the ultimate chemical states of the oxide were alsorestored. The XRD patterns collected at the end of each ofsuch treatment were identical (orthorhombic MoO3,14),except for some preferential texturing that corroborate andvalidate the stated assumption. One such pattern is shownin Fig. 7.

Microstructural Evolution in MoO3 Films Fabricated Froma Sodium Molybdate Precursor: The morphological fea-tures of the molybdic acid (H2MoO4) precipitated fromammonium molybdate and of the MoO3 powder ob-tained after the calcination of molybdic acid at 3001C/1 h are shown in Figs. 8a and b, respectively. The tri-

clinic habits of the molybdic acid crystals (ICDD 26-1449) are easily discernable as are the plate-like struc-ture characteristic of MoO3 in the fired sample.

The phase evolution in MoO3 upon the molybdicacid dehydration appears to be textured along /020Sdirection rather than the /021S direction as is the casewith the bulk oxide.14 This is shown in Fig. 9.

The microstructural evolution in films prepared fromthis powder and subjected to different redox treatmentis shown in Fig. 10.

It is worth mentioning that while the films undergodrastic morphological variations as result of exposure tothe gas environment of varying oxygen potential, theXRD signatures collected on these samples reveal that

Fig. 6. Microstructural evolution in the MoO3 thick films fabricated from MoO3 commercial powder. (a) Film calcined for 1 h. (b) ‘a’reduced by CO/CO2 for 12 h. (c) ‘b’ oxidized by CO/CO2 for 12 h. (d) c oxidized by air for 1 h.


the phase structures of a, b, and d belong to those ofMoO3 (akin to that shown in Fig. 7 with texturingalong the directions of preferred growth) and that of cbelongs to MoO2.15 It is also interesting to note thatreduction–oxidation by the conventional H2–air com-bination (Fig. 10b) leads to a rather compact and densemicrostructure albeit with a large aspect ratio; in clearcontrast to this, the microstructure, is rather open inFig. 10d when CO/CO2 buffer is used as the redoxmedium. This is suggestive of a layer-by-layer buildupof the parent oxide from the reduced intermediate phasein the case of redox treatment under conditions of re-stricted availability of oxygen. The XRD pattern (notshown here) obtained after reduction of the MoO3 inthe chosen CO2/CO buffer does contain peaks belong-ing to MoO2, which later disappeared when subjectedto oxidation in an appropriate CO2/CO mixture. Thephase change from MoO3 to MoO2 and again to MoO3

in the cyclic redox process is accompanied by concom-itant volume changes, which lead to a porous structure.Similar observations have been made in the case of re-generation of the oxide phase in the WO3 system asdemonstrated previously.6,9 This vividly illustrates thatthe oxygen potential prevalent in the vicinity of an oxidephase has a profound effect on its morphological fea-tures, which could be tailored to accentuate the sensingbehavior of a potential semiconducting oxide.

Preparatory Technique-Microstructure-Gas Sensing PropertyCorrelation in MoO3 Films: The response characteristic(such as the variation in film resistance in the presenceof different levels of CO in the ambient) of thick film

sensors fabricated from MoO3 derived from three dif-ferent preparatory techniques to CO gas (14–100 ppm)at 4501C are shown in Fig. 11. The sensitivity and re-

Fig. 7. Typical X-ray diffraction pattern of the MoO3 thick filmsafter various redox treatments with preferential texturing in the/060S direction.14

Fig. 8. Morphology of (a) H2MoO4 precipitate and (b) H2MoO4

precipitate fired at 3001C/1 h.

Fig. 9. X-ray diffraction pattern of the oxide obtained frommolybdic acid fired at 3001C for 1 h. The pattern conforms to thatof MoO3, with preferred texturing along the /020S direction.


sponse time are compared in Fig. 12 (a, as-prepared; b,sample a reduced and oxidized in CO/CO2 mixtures;and c, sample a reduced in CO/CO2 mixture and ox-idized in air). The response of the sensor toward CO isdefined as the normalized variation of resistance of then-type film, namely (R0�Rg)/R0, where R0 is the steady-state resistance (identified as a time-independent pla-teau) in the background gas (in the present case, 10%O2-bal. N2) and Rg is the steady-state resistance of thesensor film (identified again as a time-independent pla-teau) when exposed to the gas of interest (CO). Theresponse time (t90) was calculated as the time taken toreach 90% of the difference (R0�Rg).

As can be seen, neither the sensitivity nor the responsetime changed significantly as a result of various redoxtreatments given to the films prepared from the com-mercial MoO3 powder. The lack of improvement ineither of the two characteristics is believed to be because

of almost negligible variation in the microstructural fea-tures of the films (see Fig. 6). The film reduced and re-oxidized by CO/CO2 (scheme b) showed somewhatmarginally better sensitivity, while the film reduced byCO/CO2 buffer and re-oxidized by air (scheme c)showed the quickest response (B27 s).

Further, there was no improvement in the sensingbehavior in the case of films prepared from the powdersobtained after oxidation of a Mo foil either. This is un-derstandable in the light of the negligible change in themicrostructural features as a result of the three redoxtreatments (Fig. 5).

On the other hand, the film prepared from MoO3

powders derived via the precipitation technique turnedout to be the most promising in terms of sensitivity andresponse time. As can be seen from Fig. 12, significantimprovement in the sensitivity and the response timewas observed in this case. The film subjected to redox

Fig. 10. Morphology in (a) film fabricated from H2MoO4-derived MoO3 and calcined at 5001C/2 h; (b) ‘a’ heated in 10%H2–N2 mixtureat 5001C/1 h and oxidized in static air; (c) ‘a’ reduced by CO/CO2, and d: c oxidized by CO/CO2 (both at 6001C for 12 h). Note the differentlengths of the scale bars.


treatment by CO/CO2 showed improved sensitivitycompared with the untreated sample. However, thefilm reduced by CO/CO2 followed by air oxidationshowed an increase of B40% in sensitivity and decreaseby B300% in the response time at 100 ppm CO level.This is believed to be because of the unique microstruc-tural artifacts developed in these films (Fig. 10). Clearly,the third method resulted in sensor films with the mostimproved sensitivity to all levels of CO and particularlythe lowest response time for 100 ppm CO.

Microstructural Modulation in MoO3-Containing Compos-ite Structures: Similar morphological changes havebeen observed in the case of sensors prepared with atwo-phase composite mixture, viz., ZnMoO4–MoO3

(MZM). It was found that the rod-like MoO3 grainsin the original mixture are regenerated as highly orient-ed thin platelets upon exposing the MZM film to a gasmixture containing 1% CO at 4501C for 1 h, followedby natural cooling in air. It was also observed that as aresult of this bulk redox reaction, the morphological

Fig. 11. Variation in resistance of MoO3 films prepared by different techniques and subjected to different redox treatments, upon exposure tovarious levels of CO; films in (i) and (ii) were subjected to redox in CO/CO2 mixtures (scheme ‘‘b’’) and (iii) to scheme ‘‘c.’’


features of the major phase (ZnMoO4) have undergonenoticeable variation (from regular near spherical grainsto triclinic habits with well-defined sharp edges) withoutany chemical degradation which was verified from theEDS spectra collected in different pockets of the com-posite after subjecting it to the above-mentioned redoxtreatment. These results are shown in the SEM pictures(Fig. 13) and, the EDS spectra (Fig. 14). The morphol-ogy of regenerated ZnMoO4 grains in Fig. 13c isisostructural with that of the mineral microcline (KAl-Si3O8) that belongs to the triclinic pinacoidal class ofcrystals,16 http://mineral.galleries.com/minerals/symmetry/triclini.htm; see Fig. 13e.

The systematic response of a potentiometric sensorusing zinc molybdate as the active electrode toward avarying level of CO (up to 900 ppm) in air at 5001C isshown in Fig. 15.

TiO2-Based Studies

Recently, Yoo et al.18 developed a novel techniquefor generating TiO2 surfaces containing arrays of nano-fibers by a simple and highly scalable method involvingreaction with noncombustible H2/N2 gas. After nano-carving at 7001C in a H2/N2 gas atmosphere, orientedarrays of single crystal nanofibers were created over theentire surface of polycrystalline TiO2 disks. While mostone-dimensional nanostructures are created by deposi-tion or growth,19–21 the TiO2 nanofibers by the nano-carving technique are created by etching TiO2 grainsduring H2/N2 heat treatment. The nanofibers were onlyformed on the surface and the depth of the nanofiber-formed region ranged from 0.7 to 1 mm. The phase ofnanofibers was determined to be rutile TiO2 by XRD,TEM,18 and XPS,2 although the nanofibers were gen-erated under a reducing environment of H2/N2. Basedon a time sequence SEM experiment18 of a specificgrain, it was determined that the nanofiber formationoccurred by an etching process and not by a growthprocess. In addition, thermogravimetric analysis (TGA)showed that the weight of TiO2 decreased during H2/N2 treatment, which confirmed that the nanofiber for-mation was because of an etching process.12 The etchingdirection was determined to be /001S from TEM ob-servation.18 This etching was selective and anisotropic,which led to aligned arrays of nanofibers. This crystallo-graphic etching process was dubbed ‘‘Nanocarving.’’

The chemical reaction of H2 with TiO2 can berepresented as

TiO2 þ xH2ðgÞ ¼ TiO2�x þ xH2OðgÞ ð3Þ

From TGA, weight loss of TiO2 was observed duringthe H2/N2 treatment and it follows the parabolic ratelaw.13 From mass spectroscopy and ICP analysis of thereaction tube, no volatile Ti species was found. There-fore, during H2/N2 treatment, only O was removedfrom the surface. As surface stoichiometry did notchange and there was large volume change, the excessTi cations should diffuse from the surface to the bulk tomaintain the stoichiometry.

The effective diffusivity, Deff, including both thelattice and the grain boundary diffusion is given by

Deff ¼ Dl þdGðDgb � DlÞ ð4Þ

where Dl and Dgb are the diffusivity for the lattice andthe grain boundary, respectively, and G and d are the

25

20

15

10

35

30

25

20

15

10

5

0

5

014 32 66 100

14 32 66 100

a ab bc

a b c a b ca b c a b c

ca b c a b c

Fig. 12. Comparison of the sensitivity (top) and response time(bottom) of the films derived from three (a, MoO3 from vendor;b, MoO3 from Mo oxidation; and c, MoO3 from sodiummolybdate) and subjected to reduction by CO/CO2 buffer followedby air oxidation. The sensor measurements were carried outat 4501C.


widths of grain (grain size) and the grain boundary, re-spectively. If the grain boundary width d does not varymuch with the sintering temperature, then the effec-tive diffusivity is linearly dependent on 1/(grain size).

Weight-loss experiments on samples with different grainsize support this relationship, suggesting that the weightloss for nanofiber formation is related to grain boundarydiffusion.12

Fig. 13. Microstructural features of the ZnMoO4–MoO3 (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 4501C/1 hfollowed by natural cooling in air are shown in (c) and (d), respectively. (e) Motif of microcline mineral.


It was also observed that the rate of nanofiber for-mation decreased as the sintering time and temperatureincreased. As weight loss is related to grain boundarydiffusion of Ti cation, the grain size increase by sin-tering temperature resulted in decreased diffusivity,which caused the reduction of the etching rate for

nanofiber formation. Sintering time also affected therate of nanofiber formation. The sintering time at12001C was varied for 0, 1, 12, and 48 h, producingaverage grain sizes of 1.12, 2.22, 2.72, and 3.43 mm,respectively. A linear intercept analysis of SEM imageswas used to determine the grain size of sintered TiO2.These samples were then heat treated in H2/N2 at7001C for 8 h. Differences in the etched nanostructuresare evident when comparing the smallest grain sample(Fig. 16a) with the largest grain sample (Fig. 16b). Thesmaller grain size sample exhibited extensive fiber for-

Fig. 14. EDS spectrum of ZnMoO4 grains before and after theredox reaction in 1% CO at 4501C/1 h, followed by naturalcooling in air (top), and that of the regenerated platelets of MoO3

shown in Fig. 13d (bottom).

150

100

50

00 900

496 ppm

902 ppm

650 ppm

308 ppm

198 ppm

110 ppm

1800 2700 3600t /s

E / m

V

airair

Fig. 15. Reversible response of a potentiometric CO sensor usingzinc molybdate as the catalytically active electrode.17

Fig. 16. Scanning electron microscopy images of the surface ofTiO2 after sintering at 12001C for a hold time of (a) 0 h and (b)48 h, followed by exposure to a flowing H2/N2 gas mixture for 8 hat 7001C.


mation on the surface, whereas the larger grain sizesample showed less extensive fiber formation. It shouldbe pointed out that the sample with large grain size(created by longer sintering time) shows very little fiberformation, as the number of grain boundaries is muchless in this sample which slows down the diffusion rateof titanium compared with that in a sample with smallergrain size because of no soak time.

Figure 17 shows the influence of the hydrogen con-centration on the sensitivity of the nanfiber sensor. Thesensitivity is defined as the steady-state resistance (R) ofthe sensor at a given concentration of the target gas di-vided by the resistance in the absence of the target gas(R0). Compared with as-sintered titania that showed noresponse, the enhanced surface area nanofiber-basedsensors exhibited good sensitivity. Sensing tests on oth-er samples with nanostructures produced via the nano-carving process in TiO2 as well as TiO2–SnO2

composite systems show clear correlation between sen-sitivity and enhanced surface area.13

Conclusions

The following conclusions can be drawn:1. Novel microstructural features could be incor-

porated into a given semiconducting oxide via precise

oxygen potential modulation and hydrogen-containinggas-phase etching.

2. A buffer gas mixture (CO/CO2 or H2/H2O) ofan appropriate ratio provides a convenient environmentfor the proposed redox scheme; from the point of view ofexperimental ease, CO/CO2 is a better buffer mixture.

3. The formation and growth of new oxide surfaceon an atomic or submolecular level, under conditions ofoxygen deprivation, appears to be the most likely pathway,as evidenced by the analyses of the XRD patterns andmorphological features of the intermediate products.

4. The temperature–time–pO2correlation can be

effectively used as a useful tool for tailoring microstruc-tural features in oxide ceramics that are amenable to re-dox reactions.

5. Nanocarving in TiO2 is achieved through etch-ing by hydrogen-containing species. While oxygen fromTiO2 is taken out as H2O(g), Ti diffuses from the sur-face to the bulk.

6. The regenerated oxide phase with unusualmicrostructure shows enhanced gas-sensing behavior.

References

1. A.-M. Azad, ‘‘Smart Sensors—Innovative R&D to Tech Transfer: Focus onEnvironmental Remediation,’’ Eng. Horizons, 114 5–8 (1998).

2. S. Yoo, S. A. Akbar, and K. H. Sandhage, ‘‘Nanocarving of Titania (TiO2): ANovel Approach for Fabricating Chemical Sensing Platform,’’ Ceram. Int., 301121–1126 (2004).

3. C. T. Holt, A.-M. Azad, S. L. Swartz, R. R. Rao, and P. K. Dutta, ‘‘CarbonMonoxide Sensor for PEM Fuel Cell Systems,’’ Sensors Actuat. B, 87 414–420(2002).

4. A.-M. Azad and E. D. Wachsman, ‘‘Solid State Chemical Sensors for COGas,’’ Proceedings of the 198th Meeting of the Electrochemical Society: Solid StateIonic Devices II, Ceramic Sensors, Vol. 32. eds., A. Z. Phoenix, E. D. Wachs-man, W. Weppner, E. Traversa, M. Liu, P. Vanysek, and N. Yamazoe, 455–466 October 23–27 2000.

5. A.-M. Azad and M. Hammoud. ‘‘Thermodynamically-driven Reconstitutionof Ceramics to Create Nanoscale Features,’’ Nanoscale Features in Ceramics,’’in the Proceedings of the 4th ASM International Surface Engineering Con-gress & 19th International Conference on Surface Modification Technolo-gies, St. Paul, MN, August 1–3, 2005, in press.

6. A.-M. Azad and S. A. Akbar, ‘‘Ceramic Materials and Nano-structures forChemical Sensing,’’ Proceedings of the Optics East SPIE Conference onSensors for Harsh Environments II, Boston, October 23–26, 2005, in press.

7. O. Kubaschewski and C. B. Alcock, Metallurgical Thermochemistry, 5th Edi-tion, Pergamon, New York, 378–384, 1979.

8. L. B. Pankratz, J. M. Stuve, and N. A. Gockcen, Thermodynamic Data forMineral Technology. U.S. Department of Interior Bureau of Mines, Wash-ington, DC., 261–295, 1984.

9. A.-M. Azad and M. Hammoud, ‘‘Fine-Tuning of Ceramic-Based ChemicalSensors via Novel Microstructural Modification. I: Low Level CO Sensing byTungsten Oxide, WO3,’’ Sensors Actuat. B, (2006), in press.

10. M. B. Robin and P. Day, Advantage in Inorganic and Radiochemistry. H. J.Emelius and A. G. Sharpe, eds., Academic Press, New York, 335–344, 1967.

11. A.-M. Azad, S. G. Mhaisalkar, L. D. Birkefeld, S. A. Akbar, and K. S. Goto,‘‘Behavior of a New ZrO2–MoO3 Sensor for Carbon Monoxide Detection,’’J. Electrochem. Soc., 139 2913–2920 (1992).

1.1

1.0

0.9

0.8

0.7

0.6

0.5

0.4

0.3

0.20 1 2 3 4

Fig. 17. Sensitivity (R/R0) versus H2 concentration ofnanofibrillar titania when exposed to 0–4% H2 gas at 4001C;before conversion into nanofibrillar titania, the specimen wassintered at 12001C for 8 h (see text).


12. S. Yoo, H. Rick, K. H. Sandhage, S. Dregia, and S. A. Akbar, ‘‘EtchingMechanism of TiO2 Nano-Fiber Formation with Hydrogen Gas,’’ J. Mater.Res., (2005) (in press).

13. C. Carney, S. Yoo, and S. A. Akbar, ‘‘TiO2–SnO2 Nanostructures and TheirH2 Sensing Behavior,’’ Sensors Actuat. B, 108 29–33 (2005).

14. ICDD reference card # 05-0408.15. ICDD reference card # 32-1395.16. S. C. Abrahams, ‘‘Crystal Structure of Zinc Molybdate ZnMoO4,’’ J. Chem.

Phys., 46 2052–2063 (1966).17. E. D. Wachsman and A.-M. Azad, ‘‘Solid State Potentiometric CO Sensor,’’

US patent # 6,598,596, B2 (2003).

18. S. Yoo, S. A. Akbar, and K. H. Sandhage, ‘‘Nanocarving of Bulk TitaniaCrystals into Oriented Arrays of Single-Crystal Nanofibers via Reaction withHydrogen-Bearing Gas,’’ Adv. Mater., 16 260 (2004).

19. O. K. Varghese, D. Gong, M. Paulose, C. A. Grimes, and E. C. Dickey,‘‘Crystallization and High-Temperature Structural Stability of Titanium Ox-ide Nanotube Arrays,’’ J. Mater. Res., 18 156 (2003).

20. J.-J. Wu and C.-C. Yu, ‘‘Aligned TiO2 Nanorods and Nanowalls,’’ J. Phys.Chem. B, 108 3377 (2004).

21. D. Li and Y. Xia, ‘‘Fabrication of Titania Nanofibers by Electrospinning,’’Nanoletters, 3 555 (2003).


Documents

Ceramic Product Development and Commercializationaazad/pdf/IJAC.pdfCeramic Product Development and Commercialization *[email protected] r 2006 The American Ceramic Society mechanism