Advances and New Directions in Gas-sensing Devices

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Acta Materialia 61 (2013) 974–1000

Advances and new directions in gas-sensing devices

Il-Doo Kim a, Avner Rothschild b, Harry L. Tuller c,⇑

a Department of Materials Science and Engineering, Korea Advanced Institute of Science and Technology, 291 Daehak-ro,

Yuseong-gu, Daejeon 305-701, Republic of Koreab Department of Materials Science and Engineering, Technion—Israel Institute of Technology, Haifa 32000, Israel

c Department of Materials Science and Engineering, Massachusetts Institute of Technology, Cambridge, MA 02139, USA

Abstract

Gas sensors are employed in many applications including detection of toxic and combustible gases, monitoring emissions from vehi-cles and other combustion processes, breath analysis for medical diagnosis, and quality control in the chemicals, food and cosmeticsindustries. Many of these applications employ miniaturized solid-state devices, whose electrical properties change in response to theintroduction of chemical analytes into the surrounding gas phase. Key challenges remain as to how to optimize sensor sensitivity, selec-tivity, speed of response and stability. The principles of operation of such devices vary and a brief review of operating principles based onpotentiometric/amperometric, chemisorptive, redox, field effect and nanobalance approaches is presented. Due to simplicity of designand ability to stand up to harsh environments, metal oxide-based chemoresistive devices are commonly selected for these purposesand are therefore the focus of this review. While many studies have been published on the operation of such devices, an understandingof the underlying physicochemical principles behind their operation have trailed behind their technological development. In this article, adetailed review is provided which serves to update progress made along these lines. The introduction of nanodimensioned materials hashad a particularly striking impact on the field over the past decade. Advances in materials processing has enabled the fabrication of tai-lored structures and morphologies offering, at times, orders of magnitude improvements in sensitivity, while high-resolution analyticalmethods have enabled a much improved examination of the structure and chemistry of these materials. Selected examples, illustrating thetype of nanostructured devices being fabricated and tested, are discussed. This review concludes by highlighting trends suggesting direc-tions for future progress.� 2012 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: Gas sensors; Operating principles; Response mechanisms; Nanostructured architectures; Semiconducting oxides

1. Introduction

Society has benefited greatly, over the past century,from advances that have come about in the energy, trans-portation, communication and medical fields. Morerecently, society has become increasingly concerned withthe unintended consequences of these advances, includingglobal warming, pollution of air and water, and destructionof the ozone layer and forests. Many of these destructiveforces can be tied to increased emissions from powerplants, home and factory heating units, and vehicles, which

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⇑ Corresponding author.E-mail address: [email protected] (H.L. Tuller).

derive nearly 90% of their energy from fossil-fuel combus-tion processes. Add to this various often toxic and/or com-bustible gaseous and liquid products generated at chemicaland materials processing plants, and the need to insuresecurity at airports and other public sites, and it becomesobvious that the means for tracking and controlling suchemissions or chemical analytes are required. This articleconcerns itself, therefore, with gas sensors, reviewing theirprinciples of operation, the progress that has been ongoingin refining their operation and the trends defining whereprogress is likely to take us in the future.

Sensitivity is the primary property that comes to mindwhen discussing sensors. This follows from the fact thatcertain chemical species, even at ppm or lower levels, can

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I.-D. Kim et al. / Acta Materialia 61 (2013) 974–1000 975

be toxic to humans, contribute to corrosion of critical com-ponents (e.g. nuclear reactors) and/or poison catalystsessential in emissions control or to the chemicals industry.This brings into play another key sensor property, selectiv-ity, which reflects the often-enormous challenge of selec-tively detecting small numbers of a specified moleculesuspended in a sea of other chemical species, e.g. the sur-rounding atmosphere. Another important sensor parame-ter is speed of response. For example, an automotiveexhaust sensor must respond within the order of 10 ms toa change in gas composition in order to enable feedbackcontrol of the air-to-fuel ratio needed for proper operationof the catalytic convertor. The last of the four key proper-ties is stability, without which reliable sensor readingsbecome impossible. This latter property is becoming morechallenging to achieve, as we increasingly require sensors tooperate under harsh temperature and environmental condi-tions. The analysis of the four S’s—sensitivity, selectivity,speed and stability—is thus essential in any discussion ofchemical sensor development.

For many years, high sensitivity/selectivity sensor sys-tems were limited to the laboratory. More recently, thetrend has been away from such large stand-alone analyticalchemistry systems (mass spectrometers, chromatographs,IR spectrometers, etc.) that lack portability, require skilledoperators and are costly, towards miniature devices, oftenembedded as part of a sensor array. These lower-costdevices are portable, draw considerably reduced power,and when integrated with appropriate software, provide alevel of selectivity impossible with single-sensor-baseddevices. These advances have been made possible by lever-aging corresponding advances in microelectronic andmicroelectromechanical (MEMS) processing. At the sametime, it must be remembered that while silicon-based chipsnormally operate at or near room temperature, wrapped inpackaging designed to isolate the device from the environ-ment, chemical sensors, on the other hand, commonlyoperate at elevated temperatures to accelerate kinetic pro-cesses and in often harsh chemical environments. Thishas required the integration of materials not common inthe microelectronics field and the modification of the sub-strates and metallizations capable of operating under suchconditions. These efforts have been aided by the need forhigh-power (electric vehicles, power grid controls) andhigh-temperature electronics (automotive and jet enginecontrols) where wide band gap materials such as SiC andAlGaN have been introduced and continue to be refined.

Many means for detecting chemical analytes are possibleand have been investigated, including those based on chem-ically induced modulation of the electrical, electrochemical,electromechanical or optical properties of materials. Exam-ples of each will be given here. The bulk of this review,however, will focus on progress being made on so-calledchemoresistive sensors, which are particularly attractive,given their conceptually simple structure, ease of fabrica-tion and low cost, coupled with high sensitivity. While bothorganic and inorganic materials are being investigated as

the basis of chemical sensors, in this article we focus oninorganic refractory materials, given that many of theapplications that we consider require exposure to harshenvironments. Under those circumstances such materials,most commonly semiconducting metal oxides (SMOs), sat-isfy the requirement of stability.

The introduction of nanodimensioned materials has hada particularly striking impact on the field over the past dec-ade. This follows both from the high surface areas of suchstructures, but often more importantly, the matching of themodulation depth induced by the adsorbed chemical ana-lytes, with the cross-sectional dimensions of the nanosizedparticles or nanowires (NWs) that make up the device.Advances in materials processing has enabled the fabrica-tion of tailored structures and morphologies offering, attimes, orders of magnitude improvements in sensitivity,while high-resolution analytical methods have enabled amuch improved examination of the surface structure andchemistry of these materials. Progress along these lines isreviewed in this article.

The operation of chemical sensors inevitably relies on anunderstanding of a number of disciplines. The bulktransport properties of, for example, SMOs requires anunderstanding of both solid-state physics and defect ther-modynamics and kinetics. The chemisorption of molecularand atomic species on the surface of such semiconductorsrelies on charge transfer processes that involve an under-standing of semiconductor junction physics, electrochemis-try and catalysis. It is no wonder that a detailedunderstanding of the operation of chemical sensorsremains, in many cases, lacking. In this paper, we reviewprogress being made in understanding these phenomena.

Finally, we end this article by describing the directionsthat research is likely to take in the coming years withrespect to new sensor materials platforms, advanced pro-cessing and characterization approaches, light and electricfield modulation, and the modeling of sensor materialsand their operation.

2. Solid-state gas-sensitive devices: brief review of operating

principles

2.1. Electrochemical devices

There are, in principle, many ways to detect chemicalspecies in the environment. Most commonly, the sensingdevice takes the form of a chemical to electrical transducer.Classically, this would be in the form of electrochemicalcells, operating either in the potentiometric or amperomet-ric mode. Indeed, the sensors installed in tens of millions ofnew automobiles per year, for the purpose of monitoringthe oxygen partial pressure, pO2

, of the exhaust gas, arepotentiometric devices utilizing the oxygen ion solid elec-trolyte, yttria-stabilized zirconia (YSZ). An electromotiveforce (EMF) is generated across the electrolyte due to thegradient in pO2

between that in the exhaust manifold andthe air reference [1]. This oxygen activity gradient tends

976 I.-D. Kim et al. / Acta Materialia 61 (2013) 974–1000

to drive oxygen ions by diffusion from the high to the lowpO2

side. The redistribution of charge sets up an electricfield just large enough to counteract the diffusive flux withthe open-circuit EMF induced across the electrolyte givenby:

E ¼ kT4q

Z pIIO2

pIO2

tiond ln pO2; ð1Þ

where tion is the ionic transference number and the otherterms have their usual meaning. This reduces to the well-known Nernst EMF, EN, when the transference numberis unity:

EN ¼kT4q

lnpII

O2

pIO2

!; ð2Þ

where pIIO2

and pIO2

are the oxygen partial pressures on thetwo sides of the electrolyte. This type of sensor, based onthermodynamic principles, is characterized by good repro-ducibility, insensitivity to geometry and morphology andextremely wide dynamic range. Key limitations include arelatively weak dependence on partial pressure (logarith-mic), the need for a high-temperature seal to isolate the un-known from the reference gas, and high specificity for asingle gas. For conventional engines, the pO2

changes by or-ders of magnitude around the desired operating pO2

andsensor sensitivity is therefore more than adequate. Anotherpotentiometric sensor with promise is one designed to mon-itor the CO2 concentration in the atmosphere. Here, a Naion conducting electrolyte is used to monitor the Na activ-ity gradient across the electrolye. A multiphase electrode isused to fix the Na activity on the reference side while a sec-ond electrode, containing Na2CO3, has its Na activity fixedby the partial pressure of CO2 in the gas phase [2]. Designsin which such cells are miniaturized have recently beenpublished [3].

As mentioned above, potentiometric oxygen sensors areinstalled in millions of gasoline vehicles every year to opti-mize the combustion process and reduce deleterious emis-sions. Gasoline engines operate close to thestoichiometric air-to-fuel ratio where the pO2

changes byorders of magnitude upon traversing the stoichiometricratio. Therefore, potentiometric sensors provide strong sig-nals of several hundred mV despite their intrinsically lowsensitivity (Eq. (2)). For lean-burn engines, this is no longerthe case and alternative sensors are required. Here a YSZsensor, operating in the amperometric mode, is utilizedinstead. This relies on the fact that a limiting current, pro-portional to the partial pressure of the gas, but independentof applied voltage, is reached when mass transfer across theelectrolyte becomes limited by diffusive flux through theelectrode [1,4]. Sensors based on this principle are nowbeing applied in diesel-engined cars operating under lean-burn conditions [1,5]. A key challenge for such sensors isthat the limiting current depends on the morphology and

catalytic activity of the electrodes, which change with time,under operation.

2.2. Chemoresistive devices

One can distinguish two types of chemoresistive device.First we discuss those operating at sufficiently high temper-atures such that oxygen in the gas phase exchanges withoxygen ions in the lattice, leading in turn to conductivitychanges taking the form:

r / pmO2

exp � E0

kT

� �: ð3Þ

In addition to higher sensitivity, such resistive sensorsoffer low-cost fabrication, potential for miniaturizationand no need for seals and reference atmospheres, asrequired by zirconia-based sensors. Cross-sensitivity totemperature variation, via the material’s exponentialdependence on temperature, however, presents a majorchallenge. The perovskite with composition SrTi0.65

Fe0.35O3�d was found to have an unusual characteristicfor a nonstoichiometric SMO, namely having a near-zerotemperature coefficient of resistance (i.e. E0 � 0 eV) [6,7].The source of this unusual behavior has been attributedto the nearly equal but opposite dependence of mobilityand carrier generation on temperature due to the creationof an Fe-derived impurity band within the band gap ofSrTiO3 [8,9].

Next we discuss devices operating at lower tempera-tures, for which resistance changes come about fromcharge transfer from or to the SMO, resulting from chemi-sorption of gas molecules onto the surface of the semicon-ductor. This allows for the detection of a much broadergroup of gas molecules, distinguished by their relative oxi-dizing or reducing characteristics. There has been an explo-sion of interest in this category of chemical sensors over thepast decade, given the simplicity of these devices and theunusually high sensitivities achieved with the advent ofnanostructured materials. As a consequence, this reviewfocuses on this category of devices, particularly givenrecent progress in understanding device operation.

2.3. Resonant devices

The resonant frequency of piezoelectrically driven crys-tals, such as quartz, is highly sensitive to changes in resona-tor mass. For small relative mass changes Dm (typically lessthan several per cent), a measurable shift in resonant fre-quency Df, proportional to the mass change, is obtainedaccording to the Sauerbrey equation [10]:

Df ¼ �f 2

q Dm

NqS; ð4Þ

where fq is the fundamental resonant frequency of the crys-tal, e.g. quartz, N is the frequency constant of the specific

Fig. 1. Plausible reactions occurring upon interaction of TiO2 with oxygenfrom the surrounding gas phase. Reprinted with permission from Ref. [26].Copyright (2011), American Chemical Society.


crystal cut (NAT = 1.67 � 105 Hz cm), q is the density ofquartz, 2.65 g cm–3, and S is the surface area covered bythe mass-sensitive film. Given the ability to readily detectfrequency changes at the ppm level and below in resonatorsoperating at MHz and higher frequencies, this leads to thepossibility of detecting exceptionally small mass changes.When configured as microresonators, this leads to the abil-ity to detect mass changes on the picogram level or lower[11,12]. Such resonators become sensor platforms oncethey are coated with materials that selectively absorb or ad-sorb target chemical analytes. The electrically detectedmass change can then be correlated with the concentrationof analyte in the gas phase. One of the authors recentlydemonstrated this by detecting NO2 as it reacted with highsurface area BaCO3 deposited onto a quartz resonator byan ink-jet process [13]. The resonator exhibited reversiblefrequency shifts following exposure to NO2 and subsequentrecovery under CO/CO2 at 400 �C, as per the operation ofthe lean NOx trap [14]. Considerable progress has beenmade within the past decade in identifying alternative pie-zoelectric materials with high-temperature capabilities,allowing for in situ operation of such devices under harshenvironmental conditions, including the use of MEMS-based miniature devices [15–20]. As with the high-tempera-ture resistive devices, the key challenge for resonant-basedsensors is cross-sensitivity to temperature. This can morereadily be addressed with microresonator devices [20].

3. Response mechanisms of chemisorptive SMO gas sensors:

advances in understanding

SMOs such as SnO2 and TiO2 often display remarkablechanges in their electrical properties, e.g. their work func-tion and electrical conductivity, upon exposure to O2,CO, NO2 and other reactive gases. This phenomenonunderlies their application as gas and oxygen sensors.Gas sensors are typically operated in air at temperaturesbetween 100 and 400 �C. Under these conditions, the sur-face of the sensing layer, or the particles within the sensinglayer, in the case of porous materials, are covered withadsorbates. Of particular interest are the chemisorbed spe-cies which trap or transfer electronic carriers to the under-lying oxide, thereby modifying its electrical properties. Inparticular, chemisorbed oxygen or superoxide adions(O�ðadsÞor O�2;ðadsÞ, respectively) play an essential role in thesensing mechanism by modulating the interaction with tar-get gases such as H2, CO and various hydrocarbons [21–24]. As a result, the work function and surface conductivityare sensitive to these gases, as described below.

At much higher temperatures, for example, whenattempting to monitor the partial pressure of oxygen, pO2

,in the exhaust gas of automobiles, essential for controllingthe combustion process and reducing gas emissions, theinteraction involves an actual exchange of oxygen betweenthe metal oxide lattice and the surrounding gas phase [8].At these high temperatures (typically >700 �C), the oxygenexchange reaction modifies the concentrations of electrically

charged point defects (e.g. oxygen vacancies) and the com-pensating electronic charge carriers (electrons and holes)within the bulk of the sensor. As a result, the electrical con-ductivity becomes sensitive to changes in pO2

, as expressed inEq. (3).The interaction with oxygen involves various steps,representing different reactions, as exemplified in Eqs.(5.1)–(5.8) for an n-type SMO:

O2ðgasÞ þ V ðadsÞ�O2;ðadsÞ ð5:1ÞO2;ðadsÞ þ e��O�2;ðadsÞ ð5:2Þ

O�2;ðadsÞ þ e��O�22;ðadsÞ ð5:3Þ

O�22;ðadsÞ þ V ðadsÞ� 2O�ðadsÞ ð5:4Þ

2O�ðadsÞ þ e��O�ðadsÞ þO�2ðadsÞ ð5:5Þ

O�ðadsÞ þO�2ðadsÞ þ e�� 2O�2

ðadsÞ ð5:6Þ

2O�2ðadsÞ þ V þþO �OO þO�2

ðadsÞ þ V ðadsÞ ð5:7Þ

O�2ðadsÞ þ V þþO �OO þ V ðadsÞ: ð5:8Þ

[](ads) designates an adsorption site at the surface, V(ads) is avacant adsorption site, e� is an electron transferring fromthe sensor to the adsorbate, []O designates a regular oxideanion site in the bulk or subsurface region, V þþO is a doublyionized oxygen vacancy, and OO is an oxide anion (O2-) ina regular oxide anion site. Instead of using + or � to assigndefects with positive or negative charges, it is more com-mon, especially in the solid-state physical chemistry com-munity, to use � or / to assign positive or negativeeffective charges, respectively, in accordance with the Kro-ger–Vink notation of point defects [25]. These reactions areillustrated in Fig. 1 for the interaction, for example, be-tween TiO2 and oxygen [26].

There are several things to note about the interactionwith oxygen. First, the reactions described in Eqs. (5.1)–(5.8) represent one out of several possible reaction path-ways. For instance, the following pathway is also possible:

O2ðgasÞ þ V ðadsÞ�O2;ðadsÞ ð6:1ÞO2;ðadsÞ þ V ðadsÞ� 2OðadsÞ ð6:2Þ2OðadsÞ þ e��O�ðadsÞ þOðadsÞ ð6:3ÞO�ðadsÞ þOðadsÞ þ e�� 2O�ðadsÞ ð6:4Þ

Fig. 2. Characteristic thermal desorption spectrum of oxygen from ZnOð10�10Þ surface. Reprinted from Ref. [34], Copyright (1985), withpermission from Elsevier.


2O�ðadsÞ þ e��O�ðadsÞ þO�2ðadsÞ ð6:5Þ

O�ðadsÞ þO�2ðadsÞ þ e�� 2O�2

ðadsÞ ð6:6Þ

2O�2ðadsÞ þ V þþO �OO þO�2

ðadsÞ þ V ðadsÞ ð6:7Þ

O�2ðadsÞ þ V þþO �OO þ V ðadsÞ ð6:8Þ

The difference between the two reaction pathways repre-sented by Eqs. (5) and (6) is in the sequence of dissociationfrom adsorbed dioxygen molecule to oxygen adatom and inthe electron charge transfer to the respective species, i.e.steps 2–4. The rest of the steps remain the same. Othersequences, representing other reaction pathways, may alsobe possible, and it is often quite difficult to distinguish onepathway from the other. It is noteworthy that differentpathways, having the same reactant and products but dif-ferent sequences and different intermediate states, are indis-tinguishable as far as the overall reaction is concerned. Forinstance, the reaction pathways described by Eqs. (5) and(6) account for the same overall reaction:

O2ðgasÞ þ 2V ��O þ 4e=� 2OO; ð7Þ

using the Kroger–Vink notation. Eq. (7) represents the ex-change reaction of oxygen between the gas phase and themetal oxide lattice, in which the predominant defects areoxygen vacancies. Other bulk defects, adsorbed speciesand subreactions may prevail, depending on the defectequilibria of the metal oxide, the sensor temperature andthe partial pressure of oxygen in the surrounding gas phase[27,28]. In general, oxygen adsorption and incorporationinto the lattice decreases the free carrier (electron) concen-tration in n-type SMOs, in accordance with Eq. (7), but in-creases the carrier (hole) concentration in p-type SMOs,according to the reaction:

O2ðgasÞ þ 2V ��O� 2OO þ 4h�: ð8Þ

Second, the oxygen exchange reaction may terminate atdifferent steps, depending on the operating conditions ofthe sensor, as well as on its material properties: microstruc-ture, surface structure, defects and impurities. In particu-lar, the availability of free electrons that can betransferred to oxygen adsorbates may limit the amountof chemisorbed oxygen or superoxide adions, O�ðadsÞ orO�2;ðadsÞ, respectively. Thus, the surface concentrations ofdifferent adsorbed species depend on these parameters.For instance, electron-rich n-type SMOs, such as donor-doped or reduced metal oxides, can adsorb considerablymore adions than their electron-poor counterparts, i.e.undoped (intrinsic) or p-type SMOs. The latter adsorbmostly adatoms or admolecules, i.e. O(ads) or O2,(ads),respectively, and the electron transfer steps to these speciesoften limit the rate and the extent of the overall exchangereaction [27–30]. The slow kinetics of one of the reactions,the so-called rate-limiting step, may stop the overallexchange reaction from proceeding to completion. This isoften the case for gas sensors operated at elevated

temperatures, typically below 500 �C, for which the incor-poration reaction of the adsorbed oxygen adions into thelattice is very slow [31–33].

Third, considering the oxygen exchange reaction merelyas a series of chemical reactions overlooks microscopicdetails such as the structure of the adsorption sites at thesurface and the chemophysical nature of the interactionbetween the adsorbent and the adsorbed species. Thesedetails are important for understanding the sensing mecha-nism and for interpreting experimental results. The interac-tion between the sensor and the gas phase may involvephysical adsorption (physisorption), chemical adsorption(chemisorption), and formation or annihilation of surfaceand bulk defects [22]. The prevailing effect, or dominantreaction, depends on temperature, as illustrated in Fig. 2for the interaction between oxygen and the ð10�10Þ surfaceof ZnO [34]. This characteristic behavior is typical formany SMOs, with the transition from one regime toanother occurring at different temperatures for differentmaterials. For most SMO gas sensors operating at elevatedtemperatures, typically between 100 and 500 �C, chemi-sorption is the dominant effect controlling the sensingmechanism, while for resistive oxygen sensors operatingat higher temperatures (typically above 700 �C), oxygenincorporation into the lattice is the dominant mechanism.

At high temperatures, sufficiently high to overcome thekinetic barriers for all the reactions in the oxygen exchangereaction, oxygen from the gas phase interacts with bulkdefects such as oxygen vacancies, thereby changing theconcentration of electronic charge carriers—electrons andholes—as per Eqs. (7) and (8), respectively. The two reac-tions represented by these equations are linked to eachother by the intrinsic generation of electrons and holesvia thermal excitation of electrons from the valence bandinto the conduction band:

evb þ hcb� e=cb þ h�vb; ð9Þoften written simply as:

0� e= þ h�: ð10Þ


In general, SMOs lose oxygen to the surrounding gasatmosphere at high temperatures and low pO2

, and absorboxygen at lower temperatures and higher pO2

. Therefore,for every SMO, one can define the intrinsic partial pressureof oxygen, ~pO2

for which n = p and Ef = Ei, where n and p

are the electron and hole concentrations, respectively, Ef isthe Fermi energy, and Ei the intrinsic Fermi energy [8,35].At lower partial pressures of oxygen n > p and Ef > Ei, themetal oxide becomes n-type, while at higher pO2

n < p andEf < Ei, the metal oxide becomes p-type. The ~pO2

has differ-ent values for different SMOs, and the value changes withtemperature and as a function of the balance betweenacceptor and donor impurities. As a result, some SMOsare n-type under typical experimental conditions, othersare typically p-type, while some SMOs display a transitionfrom p-type behavior at high pO2

to n-type behavior at lowpO2

. This is the case, for instance, for undoped TiO2 (rutile)which undergoes a p to n transition upon reducing the pO2

from approximately 1 atm to low pressures, with the tran-sition occurring at lower pressures with decreasing temper-atures, as shown in Fig. 3 [36]. At high pO2

, above theintrinsic pressure, the electrical resistance decreases withincreasing pO2

because the material is p-type and the holeconcentration increases with increasing pO2

, in accordancewith Eq. (8). Below that pressure the material is n-typeand the resistance decreases with decreasing pO2

, in accor-dance with Eq. (7).

The amphoteric behavior of TiO2 around the intrinsicpartial pressure of oxygen ð~pO2

Þ gives rise to low sensitivitywith small changes in resistance and ambiguous responsearound this pressure. Moreover, the resistance varies withtemperature with greater sensitivity than it responds tochanges in pO2

. Therefore, very accurate and stable temper-ature control is necessary in order to attenuate false signals

Fig. 3. Typical results for the dependence of resistance of TiO2 (rutile)layers on partial pressure of oxygen in the temperature range 350–850 �C.Reprinted from Ref. [36], Copyright (1980), with permission from Wiley.

coming from temperature fluctuations, rather than trueones coming from changes in pO2

. These deficiencies limitthe accuracy of TiO2 and many other SMO oxygen sensors,with the notable exception of SrTi0.65Fe0.35O3 for which theelectrical conductivity is nearly independent of temperaturebetween 750 and 950 �C in the pO2

range between 10�5 and1 atm, as shown in Fig. 4 [7]. The origin of the exceptionaltemperature-independent conductivity of SrTi0.65Fe0.35O3

derives from a balancing of contributions with opposingtemperature dependencies related to free carrier (hole) gen-eration, on the one hand, and mobility mechanisms, on theother [8,7] .

At intermediate temperatures, typically below 500 �C,several of the above-described reactions (Eqs. (5) and (6))become sufficiently slow that the oxygen exchange reactionis effectively blocked [27,28,30]. In nanocrystalline SnO2, aprototypical SMO gas sensor material, the surface reactionis too slow, at temperatures below 500 �C, to complete thesequence of processes involved in the oxygen exchangereaction [32]. DFT calculations of the SnO2 (101) surfacesupport this conclusion, suggesting a self-limiting mecha-nism for the oxygen exchange reaction at temperaturesbelow 400 �C [33]. Up to this temperature, the oxygenexchange reaction critically depends on the surface termi-nation, which is affected by the partial pressure of oxygenin the surrounding gas phase [33,37]. At low pO2

, the sur-face adopts non-stoichiometric terminations that enableoxygen exchange. However, at high pO2

, as in mostgas-sensing applications where the sensor operates inatmospheric air, the surface quickly oxidizes and adoptsa stoichiometric termination that blocks the incorporationof oxygen into the bulk [33]. As a result, the oxygen vacan-cies in the bulk are unable to equilibrate with oxygen fromthe gas phase.

Fig. 4. The pO2dependence of the electrical conductivity of Fe-doped

SrTiO3 (SrTi0.99Fe0.01O3), SrTi0.65Fe0.35O3 (STF35) and SrFeO3 thickfilms at different temperatures from 750 to 1000 �C. Reprinted from Ref.[8], Copyright (2007), with permission from Nova Science Publishers.


Another possible cause for the self-limited interactionbetween SMOs and oxygen from the gas phase arises fromthe finite availability of electrons to be transferred to oxy-gen adions [27–30]. Numerical calculations of chemisorp-tion isotherms based on Volkenstein’s electronic theoryof chemisorption on semiconductors [38] show that the sur-face concentration of chemisorbed oxygen adions is self-limiting because of the surface charge that builds up duringthis process [39–41]. The surface charge of the chemisorbedadions gives rise to energy band bending in the spacecharge layer adjacent to the surface, as illustrated inFig. 5b. The occupation probability of the chemisorbedsurface states depends on the surface binding energy. Thus,the probability of chemisorption in the neutral stateðO0ðadsÞÞ, so-called the weak form of chemisorption [38], is

reflected in its relatively weak binding energy to the surface(q0, typically of the order of 0.1 eV), while chemisorption inthe charged state ðO�ðadsÞÞ, so-called the strong form ofchemisorption [38], is reflected in its strong binding energy(of the order of 1 eV). These values depend on the differ-ence between the Fermi level Ef and the surface stateenergy level EA. Ef � EA changes with the surface bandbending, qVs, as a function of the surface charge density.This, in turn, is proportional to the surface coverage H�

of the strongly chemisorbed adions.At low activity of the chemisorbed species in the gas

phase, i.e. at low pO2in the case of oxygen, the surface

coverage H of chemisorbed species is small and the occu-pation probability of the chemisorbed surface states ishigh because of the large energy gain in transferring anelectron from the semiconductor to the surface state (seeFig. 5a).

Consequently, most of the chemisorbed species are inthe strongly chemisorbed state and H � H� � H0. Thesurface coverage of chemisorbed oxygen adions increaseswith increasing pO2

up to the point of saturation whenthe Fermi level is aligned with the surface state energy level,i.e. Ef � EA = 0 (see Fig. 5b). At this point the occupationprobability of the surface states becomes small and theweakly chemisorbed oxygen adatoms start to dominate.As a result, the surface coverage of strongly chemisorbedadions saturates and the Fermi level is pinned at the surfacestate energy level.

Fig. 5. Energy band diagrams for dissociative chemisorption of oxygen on an nstrongly chemisorbed oxygen adions. Redrawn after Fig. 1 in Ref. [41].

The transition from strong to weak chemisorption, uponincreasing activity of chemisorbed species in the gas phase,is demonstrated in Fig. 6, showing calculated isotherms (at300 K) of oxygen chemisorption on CdS with doping levelsof 1014, 1016 and 1018 donors per cm3. The surface coverageof strongly chemisorbed adions approaches saturation atatmospheric pressure, reaching approximately 2 � 10�3

(i.e. 0.2%) of a monolayer for the highest doping level(ND = 1018 cm�3). The calculations show that the surfacecoverage of the strongly chemisorbed adions is propor-tional to the square root of the doping level, i.e.H� /

ffiffiffiffiffiffiffiNDp

. These results are in agreement with the so-called Weisz limit of chemisorption on semiconductors [42].

The strongly chemisorbed adions trap electrons from thesemiconductor, localizing them in surface states. The spacecharge region adjacent to the surface becomes depleted offree delocalized charge carriers and consequently the con-ductivity in this region is lower than the bulk conductivity(i.e. where the energy bands are flat). The low conductivityin the depletion region has an impact on the overall resis-tance of the sensor, especially in thin films or porousnanomaterials of high surface to volume ratio. In suchmaterials, small changes in the surface coverage of stronglychemisorbed adions typically lead to significant changes inresistance. These changes underlie the operation mecha-nism of SMO gas sensors used for the detection of traceconcentrations of target gases such as CO, hydrocarbons(HCs), volatile organic species (VOCs) and other reactivegases in air.

In air, the surfaces of the sensor become saturated withstrongly chemisorbed oxygen adions. The adion types andtheir microscopic character may differ between differentsensors, depending on the material’s bulk properties, sur-face orientation, termination and structure, intrinsic (e.g.oxygen vacancies) and extrinsic point defects (impurities)both within the bulk and at the surface, surface adsorptionof foreign species, as well as operation conditions. In par-ticular, temperature and humidity are known to affect thetype of adions that participate in the gas-sensing mecha-nism [21–23]. Without overlooking the importance of theseeffects for understanding the gas-sensing mechanism at themicroscopic level, we do not discuss them further, butrather refer to the classic textbook on surface science of

-type semiconductor: (a) in flat band conditions; (b) upon saturation of the

Fig. 6. Calculated chemisorption isotherms at ambient temperature(300 K) for oxygen chemisorption on CdS with doping levels of 1014,1016 and 1018 donors cm�3. The surface coverage, expressed as a fractionof monolayer (ML), of strongly (H�), weakly (H0) and total(H = H� + H0) chemisorbed species is plotted as a function of the partialpressure of oxygen. Reprinted with permission from Ref. [41], Copyright(2002), American Institute of Physics.


metal oxides [43] and more recent review articles on thesurface science of SnO2 and TiO2 [44,45].

For a phenomenological understanding of the sensingmechanism it is essential to consider the interactions of adi-ons with reducing gases such as CO, HCs and VOCs asexemplified by the following reaction:

COðgasÞ þO�ðadsÞ ! CO2;ðgasÞ þ V ðadsÞ þ e�: ð11Þ

As a result of this reaction, the concentration of freeelectrons increases and consequently the resistancedecreases for n-type or increases for p-type SMOs [21].Exposure to oxidizing gases such as NO2 or Cl2 results inan inverse effect on sensor resistance with electrons trappedat surface states [21]. Furthermore, surface reactions withtarget gases may lead to bidirectional changes in the resis-tance of nearly intrinsic semiconductors, with amphotericconductivity of both electrons and holes [46].

While chemisorbed oxygen adions clearly play a centralrole in the response to reducing gases in air [21–24,34],chemisorbed hydroxyl adions may also be involved in theresponse mechanism when operating in humid air [47].By combining the surface chemical reaction between thereducing gas and the chemisorbed adion (e.g. Eq. (11))together with the adsorption/desorption reaction of thatadion (e.g. Eqs. (5.1)–(5.4) or (6.1)–(6.4), which can besummed up to give 1=2O2;ðgasÞ þ V ðadsÞ þ e��O�ðadsÞ), onecan derive a power law relationship between the resistanceR and the concentration Cgas of the reducing gas in air[21,22]:

R / C�agas: ð12Þ

Different power exponents (a), usually smaller than 1,have been observed empirically [48]. The power exponent acan, in principle, be derived from the mass action laws andrate constants of the respective surface reactions and therelationship between the sensor resistance and the trapped

charge density of chemisorbed adions [21,22,49–51]. In prac-tice, however, this derivation is complicated, because therelationship between the resistance and gas concentrationdepends not only on the surface chemical reaction, but alsoon the morphology of the sensor, e.g. compact (non-porous)thin film or porous thick layer, and on its microstructure(grain size, neck area between partially sintered grains,etc.) [52,53]. Furthermore, inhomogeneous distribution ofgrain size, pore size and other microstructural features, oftenthe case rather than the exception, may modify the correla-tion between microstructure and transport propertiesbeyond the idealized and oversimplified scheme of identicalstructural units. Such disorder effects have been found tohave an effect on the relationship between the sensor resis-tance and gas concentration, resulting in modified powerexponents with respect to that for idealized, perfectlyordered, microstructures [54,55].

Fig. 7 illustrates the convoluted effect of the surfacechemical reaction with gas analytes and the microstructureof the sensor on the electrical response to the gas [56]. Thesurface chemical reaction can be referred to as the receptorfunction of the sensor, because it converts a chemical reac-tion with the gas analyte into a charge transfer event. Thistriggers an electrical signal, which is subsequently amplifiedby the charge transport mechanism through the sensor, theso-called transducer function of the sensor [57–59]. Thesensor response is a convolution of the receptor and trans-ducer functions, as illustrated in Fig. 7.

Fig. 8 illustrates the transducer function of sensors ofdifferent morphologies and microstructures [60]. Explicitexpressions for the relationship between the sensor’s elec-trical response, often expressed as the relative change inthe conductance (G = R�1) upon exposure to the gas ana-lyte, normalized to the baseline conductance in air, andthe gas concentration, have been developed for differentsensor morphologies [21,22,49,51–53]. Approximated ana-lytical expressions, usually power laws with different powerexponents, have been derived for some prototypical sensormorphologies. These include compact thin films, in whichthe interaction with the gas occurs only at the surface ofthe film [51,53,61–63] and polycrystalline porous layers,in which the interaction occurs at the surface of the grains,as well as at grain boundaries and interfaces betweendomains or agglomerates of crystallites [52,53,62]. Formicrocrystalline grains, the predominant effect underlyingthe resistance changes upon exposure to the gas analyteis modulation of the grain boundary potential barrier,ugb. This barrier is proportional to the square of the chargedensity (per unit area) trapped at the grain boundary [64](i.e. ugb / ðN�t Þ

2Þ. The trapped charge density N�t ismodulated by the interaction with the gas that changesthe density of adions at the grain boundaries [62]. Thismodulation changes the sensor resistance, which dependsexponentially on the grain boundary potential barrier [62],

R / expqugb

kT

� �: ð13Þ

Fig. 7. Receptor and transducer functions of SMO gas sensors: (a) chemisorption and reaction between reducing gases (CO) and oxygen adions (O�) atthe surface give rise to the receptor function; (b) electronic charge transport through the grains and across grain boundaries gives rise to the transducerfunction. The latter depends on the microstructure of the sensing layer, e.g. on the grain size and pore size. (c) The sensor element comprises of the sensinglayer, electrodes for electrical measurements, substrate and integrated microheater. Reprinted from Ref. [56], Copyright (2005), with permission fromWiley.

Fig. 8. Schematic illustration of the transducer functions of SMO gas sensors of different morphologies and microstructures. Left to right: compact thinfilms (surface effect), porous microcrystalline layers (grain boundary effect), porous nanocrystalline layers (surface traps effect), and metal–semiconductorjunctions (Schottky junction effect). Reprinted from Ref. [60], Copyright (2007), with permission from Elsevier.


This is, essentially, the transducer function of microcrys-talline porous sensors.

The transducer function of nanocrystalline SMO gassensors is quite different than that of their microcrystallinecounterparts, leading potentially to much higher sensitivi-ties. The exceptionally high sensitivity of nanocrystallineSnO2 gas sensors was first reported by Xu et al. in 1991[65], and subsequently by other researchers (and for othermetal oxides) [56]. The empirical observations weresuccessfully explained by Rothschild et al. [66] with an ana-lytical model that takes into account the unique character-istics of the space charge region in nanocrystallinesemiconductors. Unlike their microcrystalline counter-parts, nanocrystalline semiconductors do not have poten-tial barriers at the grain boundaries [64]. This is aconsequence of the fact that when the radius of the grains

is smaller than the Debye length, typically of the order of10 nm, the space charge region at one face of the crystalliteoverlaps with that at the opposite face, such that the entirevolume of the crystallite becomes depleted of free chargecarriers. Consequently, the energy bands are flat acrossthe entire crystallite and the potential barriers at the grainboundaries disappear [64]. The position of the energybands relative to the Fermi level shifts in the presence ofsurface traps with respect to their position in the trap-freestate, such that the free carrier concentration becomessmaller across the entire crystallite. Essentially, the chargecarriers distribute between localized states at the surface ofthe crystallite and delocalized states in the bulk. This distri-bution gives rise to the following relationship between thefree carrier concentration, n, and the trapped chargedensity at the surface of the crystallite, N�t :


n ¼ n0 � N�t �SV; ð14Þ

where n0 is the free carrier concentration in the trap-freestate (i.e. when N�t ¼ 0Þ and S

V is the surface to volume ratioof the crystallite [67]. This apparently simple relationshipis, in fact, more complicated than it may look at firstglance. The salient point is the fact that N�t scales withthe size of the crystallite, D, because of the finite numberof charge carriers in the crystallite. N�t can be numericallycalculated using Poisson’s equation, the electroneutralitycondition (Eq. (14)) and the occupation probability ofthe surface states which follows the Fermi–Dirac distribu-tion function [66,67]. Numerical calculations, based on thismodel, show that the normalized differential sensitivity tochanges in the trapped charge density at the surface,s ¼ �@ðngas=nairÞ=@ðN�t;gas=N�t;airÞ, is proportional to the sur-face-to-volume ratio of the crystallites [66,67]. The latter isinversely proportional to the crystallite size D (e.g. forspherical crystallites S

V ¼ 6=DÞ, and therefore the sensitivityincreases with decreasing crystallite size.

Fig. 9. The effect of crystallite size (D) on the sensitivity of nanocrystalline SnO[66]. (b) Experimental results, reprinted from Ref. [65], Copyright (1991), with

Fig. 9a shows the normalized differential sensitivity s onthe left y-axis and the sensitivity s = ngas/nair = Rair/Rgas toa change of 1 ppm (1 � 10�6) in the trapped charge densityat the surface on the right y-axis, calculated for nanocrystal-line SnO2 with a doping level (ND = n0) of 1 � 1017 cm�3,surface state density (Nt,air) of 1 � 1012 cm�2, surface stateenergy level of 1 eV below the conduction band edge, andtemperature of 600 K [66]. The functional relationshipbetween the sensitivity and the crystallite size, s / D�1, isin good agreement with experimental observations, as dem-onstrated in Fig. 9b, which shows the response of a nano-crystalline SnO2 sensor upon exposure to 800 ppm of H2

or CO in air, measured at 300 �C [65]. However, thereappears to be a large difference in the magnitude of theresponse between the model calculations and the experi-mental results, because the calculations for a small changeof 1 ppm in the trapped charge density at the surface(Fig. 9a, right y-axis) produce quantitatively similar resultsto the experimental results measured at much higher gasconcentration (800 ppm). This apparent discrepancy may

2 gas sensors. (a) Model calculations, redrawn from data presented in Ref.permission from Elsevier.


result from inexact parameters used for the calculations,because no attempt to fit the calculations to the experimen-tal results was carried out. But it may also suggest that themodel calculations do not capture the entire complexity ofthe sensing mechanism. Indeed, the model accounts forthe response of one crystallite, whereas the macroscopicresponse of the entire sensor may not necessarily be astraightforward extrapolation of the response of a singlecrystallite.

Realistic modeling of the sensor response must take intoaccount not only the individual response of a single nano-particle within the sensor, but also the collective responseof the entire ensemble of many nanoparticles that are notidentical in size, shape, orientation and connectivity totheir close neighbors. These microstructural variations giverise to a random distribution of grain boundary barriers,which has a strong effect on sensor resistance (see Eq.(13)). The effect of disorder on the sensor response has beenstudied by Sukharev [68] using percolation theory [69] witha random resistor network that accounts for the variationsin grain boundary barriers. With this model, Sukharev wasable to reproduce most of the experimental observationsfor both the transient and steady-state response of poly-crystalline metal oxide gas sensors. Interestingly, the modelsuggests that, near the percolation limit, the gas-inducedchanges in grain boundary barriers may give rise to excep-tionally high sensitivity [68].

Another interesting work on the effect of microstruc-tural disorder on sensor response was reported by Williamset al. [54,55], who looked at the effect of grain size and poresize distribution, such as often obtained by agglomerationof nanocrystallites into larger agglomerates and clusters.Unlike Sukharev’s model which considered the effect ofvariations in the grain boundary barriers, Williamset al.’s model explores the effect of the low conductivityin the depletion region adjacent to the surface of nanopar-ticles and agglomerates that are directly exposed to the gasphase with respect to the conductivity in regions that arenot exposed to the gas phase. Thus, Williams et al.’s workexamines the effect of the active volume in porous sensors.One of the most interesting conclusions of this work relatesto the effect of porosity on sensitivity, suggesting that thereis a critical point at which sensitivity passes through a max-imum [54]. The optimal porosity depends on the degree ofcoalescence of nanoparticles to larger agglomerates and onthe ratio between surface and bulk conductivities. Theunderlying reason for this trend is related to the way sensi-tivity is defined by Williams (and many other researchers),s = DG/G0 (where G is the conductance, G = R�1). DG

increases with increasing surface area exposed to the gasphase. The latter varies parabolically with porosity, reach-ing a maximum at a given porosity value. On the otherhand, the baseline conductance, G0, decreases monotoni-cally with increasing porosity. Thus, the ratio DG/G0

reaches a maximum at some porosity value. It is also note-worthy that both Sukharev’s and Williams et al.’s studiesclearly show that the exponent in the ubiquitous

power-law relationship between the sensitivity and gas con-centration (Eq. (12)) is strongly affected by structural disor-der, giving rise to exponents that may be quite differentfrom the ones expected from the stoichiometry of the sur-face reaction [54,55,68].

Besides microstructural disorder, another importanteffect on the collective response of the ensemble is the gascomposition profile across the sensor. A gradient in gasconcentration may result from the consumption of thegas (reactant) by the surface reaction, converting it to anon-reactive product (see, e.g., Eq. (11)). Consequently,the gas concentration at the bottom section of the sensoris smaller than at the top exposed surface, as illustratedin Fig. 10. This may lead to a gradient in local resistance,RðzÞ / ½CgasðzÞ��a (Eq. (12)), across the sensor, therebyaffecting the overall resistance, R ¼ 1

H

R H0

RðzÞdz (where H

is the thickness of the sensing layer). The influence of thiseffect on sensor response was pioneered by Brailsford andLogothetis [70] and subsequently elaborated by Gardner[71–73] and other researchers [74–76] using a diffusion–reaction model in porous media. The extent of this effectdepends on the thickness, porosity and surface area ofthe sensing layer, the position of the electrical contacts(top or bottom electrodes) and the ratio

ffiffiffiffiffiffiffiffiffiD=k

pin which

D is the diffusion coefficient of the gas inside the poroussensing layer and k is the rate constant of the respectivesurface reaction (e.g. Eq. (11)). Fig. 11 shows calculatedgas concentration profiles, at steady state, as a functionof distance from the surface, for different


pratios

[75]. The penetration depth of the gas into the sensor scaleswith


p, while the rest of the layer remains inactive.

This reduces the sensitivity, but on the other hand, it opensup opportunities for tailoring the sensitivity to differentgases by using multisensor arrays of different thicknessesor by modulating the operating temperature leading tovariations in the rate constants (k) for different gases[77,78].

4. Novel nanostructured architectures

The application of novel nanostructured architectures, asmentioned above, has contributed to significant progress inthe development of highly sensitive and selective chemicalsensors. These architectures include one-dimensional (1-D)and quasi-1-D building blocks, carefully assembled by acombination of top-down controlled chemical etching andvarious bottom-up nanofabrication processes [79–81],including hollow spheres [82–84] and hemispheres [85–88],with thin-walled nanostructures, synthesized using hardand soft templates, and their hierarchically assembled nano-building blocks with enhanced surface properties [89–92].These unique structural features, possessing high surfacearea, high porosity and effective surface depletion modula-tion, offer unparalleled sensing characteristics. In thisreview, we provide a comprehensive overview on recentstate-of-the-art research activities, focusing on advancedand unique processing routes for obtaining 1D nanobuilding

Fig. 10. Schematic illustration of the key microstructural effects controlling the sensitivity of SMO gas sensors: chemisorption-induced changes in thetrapped charge density at the surface of the crystallites (illustrated on the right-hand side) and decreasing gas concentration inside the sensing layer due toreaction and conversion to non-reacting products (illustrated on the left-hand side). D and L are the grain size and the thickness of the depletion layer,respectively.

Fig. 11. Calculated gas concentration profiles inside a porous sensinglayer for different


pratios. Reprinted from Ref. [75], Copyright

(2001), with permission from Elsevier.


blocks such as single-crystalline metal oxide NWs [79,93–96], nanotubes [97–100] and polycrystalline nanofibers[101–105], nanostructured macroporous hollow spheres[82–84] and hollow hemispheres [85–88], as well as hierarchi-cal nanostructures consisting of higher-dimensional build-ing blocks assembled from lower-dimensional buildingblocks [89–92,106–109]. Some of these nanostructuredbuilding blocks are further modified by surface decorationby nanocatalysts [110–113]. The chemical sensing propertiesand applications of these structures are then reviewed withthe integration of chemical sensors into field effect transis-tors highlighted at the end.

4.1. One-dimensional nanobuilding blocks (wires/tubes/

fibers)

4.1.1. NanowiresSMOs with 1-D structures have attracted much atten-

tion within the last decade due to their unique structural

and electrical features. A number of synthetic approacheshave been suggested to provide high-aspect-ratio 1-DSMOs. They include catalyst-promoted vapor phase[114,115], template-assisted [116,117], hydrothermal [118],solution [119] and electrospinning methods [120,121]. Inparticular, single-crystalline n-type (ZnO, Nb2O5, In2O3,Fe2O3, WO3, SnO2) and p-type (Co3O4, NiO, CuO) semi-conducting NWs, successfully synthesized via vapor–liquid–solid (VLS) processes, further modified by thermalevaporation, pulsed laser ablation and chemical vapordeposition, have been extensively integrated in the formof single NWs or NW networks. In the case of vapor-phaseapproaches, metal precursor sources such as Sn, Zn and Insalts are evaporated at high temperature in tube furnaces,and reaction gases, i.e. O2, are introduced with chemicalreactions occurring on the catalytic Au-coated substrates.By the formation of eutectic structure between the targetmetal precursor and catalyst, Au, supersaturation and con-densation are generated at temperatures far below the melt-ing temperature of the target metal precursors, resulting inthe vertical growth of metal oxide NWs [122,123]. A num-ber of metal oxides can be easily synthesized in this mannerand interconnected and/or assembled into sensing layers.Because the lengths of the metal oxide NWs are in therange of 10–60 lm, they can be interconnected for chargetransport in multiple NW networks, and coated onto pre-patterned Au catalyst-coated electrodes [124]. The collectedmetal oxide NWs can also be mixed with binders to form apaste and printed onto sensor substrates to form sensinglayers. An additional calcination step gives rise to ran-domly interconnected networks. These structures facilitategas transport between and towards the individual NWs.

Sensors, even when sufficiently sensitive, often sufferfrom (a) poor selectivity, (b) slow response and (c) aging.All three of these features are related to the inadequate cat-alytic performance of sensor surfaces. The surface proper-ties of metal oxide NWs can be further modulated by


contact with catalytic layers. In general, catalysts have twoimportant functions, i.e. chemical sensitization and elec-tronic sensitization, in catalyst-loaded metal oxide sensors.Catalysts offer activation or spill-over of target gases aswell as electron donor or acceptor character, leading tochanges in adsorbed oxygen concentration and variationsin the oxidation states of the catalysts. Fig. 12a shows ascanning electron microscopy (SEM) image of representa-tive SnO2 NW networks grown by thermal evaporationusing Sn metal powder, typical of the VLS approach[125]. The single-crystalline SnO2 NWs exhibited diametersof 50–100 nm and lengths of several tens of micrometers.As shown in Fig. 12c, the response (Ra/Rg) of the pristineSnO2 NW network sensor to 100 ppm C2H5OH at 450 �Cwas 61.7. Further improvements in sensitivity wereachieved by decorating the surface of the SnO2 NWs withcatalytic Ag and Ag2O deposited by e-beam evaporationfollowed by subsequent calcination at 450 �C in air. ThinAg films can be easily agglomerated into nanoclustersand partially oxidized at elevated temperature, resultingin the formation of Ag2O particles (Fig. 12b). In the caseof 5 nm thick Ag-coated SnO2 NW network sensors, theRa/Rg value dramatically increased to 228.1 upon exposureto 100 ppm C2H5OH. On the contrary, the responseof 50 nm thick Ag-coated SnO2 sensors completely

Fig. 12. (a) SEM images of the pristine SnO2 NW networks; (b) TEM image ofand Ag-decorated SnO2 NWs with different Ag-coated thicknesses to 100 ppCopyright (2011), American Chemical Society.

disappeared due to conduction through the relatively thickand highly conductive Ag particles, regardless of chemore-sistive variation (Fig. 12c). This example clearly demon-strates that the size, concentration and interconnectivityof the decorative catalytic nanoparticles are very importantfor enhancing the gas-sensing characteristics of SnO2 NWs.

Gas-sensing properties can be further improved by theformation of n-type/p-type SMO junctions on the surfacesof NWs due to the extension of depletion zones. In general,the individual n-type or p-type NWs comprise two distinctregions. The inner core of the NWs forms a conductivezone and the outer layer serves as a gas-sensitive surfacedepletion zone, typically <10 nm thick. Oxygen gas mole-cules are absorbed on the outer surface, resulting in the for-mation of a surface depletion layer due to the adsorbedoxygen adions, e.g. O�2;ðadsÞ species. In the case of pn junc-tions formed between SMOs, additional depletion layersare created, leading to more efficient resistance modulation.

Fig. 13a and b show representative examples of pristinen-type ZnO NWs and pn junctions between lenticularp-type Co3O4 islands decorated on ZnO NWs, respectively[126]. The ZnO NWs were grown directly onto an aluminasubstrate with two gold electrodes serving as catalyst. The30–70 nm thick ZnO NWs show high crystallinity withð0 1�10Þ fringes separated by 2.8 A (inset of Fig. 13a).

Ag-decorated SnO2 NWs; (c) responses (Ra/Rg) of the pristine SnO2 NWsm of C2H5OH at 450 �C. Reproduced with permission from Ref. [125].

Fig. 13. (a) SEM image of highly crystallized ZnO NWs grown directly onto an alumina substrate; (b) SEM image of Co3O4-decorated ZnO NWs;Reprinted with permission from Refs. [126]. Copyright (2011), Royal Society of Chemistry. (c) TEM image of Pt-decorated SnO2 NWs; Reprinted withpermission from Refs. [127]. Copyright (2011), Royal Society of Chemistry. (d) TEM images of In2O3 NWs with different Au loading; Reprinted withpermission from Refs. [114]. Copyright (2011), American Chemical Society. (e) TEM image of a single Nb2O5 NW and Fourier-filtered HRTEM image ofthe marked area with the corresponding crystal structure (inset); Reprinted with permission from Refs. [95]. Copyright (2012), Elsevier. (f) SEM image offree-standing CuO NW arrays; Reprinted with permission from Refs. [128]. Copyright (2010), IOP Publishing. (g) SEM image of locally synthesized ZnONWs by local heating; (h) magnified SEM image of (g); Reprinted with permission from Refs. [129]. Copyright (2011), American Chemical Society. (i)SEM image of uniaxially grown ZnO NW clusters. Reproduced with permission from Ref. [130]. Copyright (2007), Philosophical Magazine.


Discrete Co3O4 clusters, with lengths ranging from 80 to110 nm, were grown on the ZnO NWs by thermal evapora-tion of CoCl2 powders at 500 �C in an Ar–O2 gas mixture.Co3O4–ZnO NW sensors showed much enhanced responseto C2H5OH (21.9 at 400 �C) while that to 5 ppm NO2

became negligible (1.36). The Co3O4 nanoparticles providea very effective catalytic effect with respect to C2H5OH gassensitivity. The coating of a discrete configuration of p-typesemiconductors onto n-type semiconductor NWs providesa simple, versatile and powerful tool for controlling bothgas selectivity and sensitivity.

Noble metals such as Pt and Pd can also be decoratedonto metal oxide NWs. Lin et al. reported atomic layerdeposition (ALD)-assisted Pt decoration on SnO2 NWsexhibiting a very uniform coating of discrete 5 nm Ptnanoparticles (Fig. 13c) and ultrahigh gas response of

8400–500 ppm ethanol gas at 200 �C [127]. Consideringthe uniform conformability of the ALD method, catalyticdecoration by ALD is a promising approach for the fabri-cation of high-sensitivity gas sensors. Singh et al. reportedon Au-nanoparticle functionalized In2O3 NWs preparedvia a self-assembled monolayer of p-minophenyltrimethox-ysilane (APhS-SAM) (Fig. 13d) [114]. Pre-synthesized Aunanoparticles were immobilized on In2O3 NWs and servedto enhance CO oxidation due to increased oxygen chemi-sorption on the conductive Au-nanoparticle surfaces[114]. In addition, metal oxide NWs can also be directlysynthesized by thermal oxidation of metallic foils such asNb (Fig. 13e) [95] or Cu [128] within a vacuum tube fur-nace, leading to NWs of Nb2O5 and CuO of approximately30 and 30–100 nm diameter as shown in Fig. 13e and f,respectively.

Fig. 14. Schematics of adsorbed O� adions on surfaces of different geometrical structures such as (a) thin-walled films and (b) tubes with large surface-to-volume ratios, (c) illustration of procedures for fabricating random and aligned hollow ZnO tubes on interdigitated electrodes using an electrospunnanofiber templating method. Reproduced with permission from Ref. [134]. Copyright (2009), American Chemical Society.

1 For interpretation of color in Fig. 14, the reader is referred to the web


Hydrothermal growth has been a viable strategy forobtaining well-controlled metal oxide NWs. Jin et al.reported on localized growth of ZnO NWs based on localheating [129]. As shown in Fig. 13g and h, Zn precursorsare precipitated in a selected nanoscale region via localizedapplication of an array of individually addressable NWJoule heaters [129]. Uniaxial fuzzy ZnO NW clusters weresynthesized through a simple vapor-transport depositionprocess in a single-zone horizontal tube furnace [130].These nanostructures show increased surface activity(Fig. 13i) [130].

4.1.2. Nanotubes

1-D hollow metal oxide nanotubes have been consideredas ideal structures for application in chemical sensors dueto their desirable geometrical characteristics leading to highreactivity with the gas phase owing to their large surface-to-volume ratio. As illustrated in Fig. 14a and b, the num-bers of adsorbed O� adions on tubes are approximately 2ptimes higher than those on planar thin films. An increase inthe density of electron-depleted regions significantlyimproves gas response. By controlling the wall thicknessesand diameters of tubes, one can manipulate the resistanceof sensor layers and their effective gas response. Hollowfiber sensors have high-resistance surface-depleted regionsand relatively conductive cores as depicted schematicallyin Fig. 14b. The thickness and resistance of the

surface-depleted regions (see green1 color zone) vary as afunction of gas composition due to adsorption of electro-negative gases, while the remaining relatively conductivecores are in blue.

Thus far, several promising routes have been suggestedfor the synthesis of hollow metal oxide tubes. These includeelectrochemical anodization of metallic plate [131–133],anodized aluminum oxide (AAO) nanoporous membranetemplating [117], sacrificial templating routes using poly-meric nanofibers as a sacrificial layer [134–137], and selec-tive removal of the core in core/shell nanofibers [138]. Veryrecently, a sacrificial templating approach, using electro-spun nanofibers, was introduced to synthesize hollow metaloxide tubes. Polymer nanofibers are one of most promisingtemplates, which can be massively synthesized and entirelydecomposed after calcination at elevated temperature, asillustrated in the preparation of hollow ZnO tubes(Fig. 14c). These nanofiber template-assisted routes couldbe combined with physical vapor deposition (sputtering)[134–137] and ALD [98,139,140] for deposition of variousmetal oxide films on nanofibers. High-temperature heattreatment then leads to the removal of the core polymertemplate and crystallization of the metal oxide overlayer,resulting in hollow metal oxide fibers that replicate the

version of this article.

Fig. 15. The resistance response (R/R0) during cyclic exposure to increasing NO2 concentrations at 350 �C of sensors comprising a network of non-aligned(red) or quasi-aligned hollow ZnO fibers (black) and a reference ZnO thin film sensor (blue). Reproduced with permission from Ref. [134]. Copyright (2009),American Chemical Society. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

Fig. 16. (a) Cross-sectional view of InGaZnO4 tubes with thin wall thickness;(b) back-side view of InGaZnO4 tubes showing inner and outer surface;Reproduced with permission from Refs. [137]. Copyright (2011), Elsevier;(c) thin-walled Al2O3 tubes with uniform thickness using ALD; Reproducedwith permission from Refs. [140]. Copyright (2007), American ChemicalSociety. (d) schematic illustration of Pt-functionalized hollow NiO tubes withthin-walled structure using co-sputtering technique. Reproduced withpermission from Ref. [136]. Copyright (2011), Royal Society of Chemistry.


polymer fiber template (Fig. 14c). For example, polymer(polyvinyl acetate, PVAc) nanofiber templates were pro-duced by electrospinning and subsequently coated by a thinlayer of ZnO by means of reactive sputtering. FollowingZnO deposition, the samples were calcined (at 500 �C) toremove the polymer template and crystallize the ZnO over-layer, resulting in hollow ZnO fibers that replicated thepolymer fiber template. Calcined ZnO tubes showed asym-metric wall thickness due to the line of site deposition char-acteristic of sputter deposition. Alignment of ZnO tubesbecomes possible by appropriate control of the electrodeground design (Fig. 14c).

A key advantage of the nanofiber-assisted route is its ver-satile synthesis of ultra-thin-walled nanotubes with a num-ber of materials such as ZnO, SnO2, TiO2, etc. achieved bychanging sputtering target materials or metal oxide precur-sor coating materials. Fig. 15 shows the gas responses ofsensors using ZnO thin films, randomly oriented hollowfibers and quasi-aligned fibers for detection of NO2 gasesat 400 �C. The inset SEM image shows ZnO hollow fibersthat were found to exhibit higher sensitivity to 2 ppmNO2 than the ZnO thin films. Both the inner and outer sur-faces of the hollow fibers are easily accessible to NO2, whileonly the outer surface is accessible to the gas phase in thethin film case. Surface depletion regions evolve on boththe inner and outer surfaces of the ZnO fibers, resulting inenhanced sensitivity. The sensitivity enhancement ofquasi-aligned compared to random fibers may result fromimproved gas transport properties and the reduced effectof inter-fiber charge transport at the junctions betweencrossing fibers, which appear to be less accessible to gasmolecules compared to the fiber surface [134].

Amorphous InGaZnO4 metal oxide tubes were synthe-sized via the nanofiber templating route combined with

Fig. 17. (a) Hollow Pt–NiO–Pt tube-coated sensor; (b) magnified image ofthe sensing layer of (a) showing random networks; (c) gas response (Rgas/Rair) of Pt-functionalized hollow NiO tubes, pure hollow NiO tubes andNiO thin films to C2H5OH with gas concentration in the range of 2.5–100 ppm at 400 �C. Reproduced with permission from Ref. [136].Copyright (2011), Royal Society of Chemistry.


radiofrequency (RF) sputtering of the InGaZnO4 thin layer(Fig. 16a and b) [137]. Open tubular hemitubes of InG-aZnO4 were successfully synthesized. Fig. 16a shows thatthe InGaZnO4 hemitubes had a top wall thickness of35 nm and a side wall thickness of 20 nm. Due to ashadowing effect, a slight asymmetry in wall thickness

Fig. 18. (a) SEM images of TiO2 tube arrays fabricated by anodization; Reprodimages of Ag nanoparticle embedded SnO2 tubes. Reproduced with permissio

was observed [134,137]. Shorter sputtering times inducehemitubes with thinner walls. The step coverage by sputter-ing is not as good as that achieved by chemical vapor depo-sition. Recently, several groups have attempted to combineelectrospun polymeric nanofiber templates with ALDmetal oxide growth [98]. Very thin-walled Al2O3 tubes withhomogeneous thickness distribution were obtained asshown in Fig. 16c [140]. More recently, Cho et al. demon-strated sandwiched Pt–NiO–Pt open tubular structures forapplication in ethanol sensors. Fig. 16d illustrates hollowthin-walled NiO tubes functionalized by catalytic Pt, syn-thesized via nanofiber templating and multilayered sput-ter-coating of Pt and NiO thin overlayers followed byheat treatment at 600 �C [136]. Fig. 17a and b exhibitSEM images of sensors using Pt–NiO–Pt tube networks withwall thicknesses of 10–20 nm. Fig. 17c shows the gasresponse of NiO thin film, NiO hollow tube and Pt–NiO–Pt tube sensors to 2.5–100 ppm C2H5OH at 400 �C.Sandwich Pt–NiO–Pt tube networks exhibited a superiorC2H5OH sensing response. Hollow thin-walled NiO tubesfunctionalized by Pt, i.e. sandwich Pt–NiO–Pt tubenetworks, were synthesized via nanofiber templating andmultilayered sputter-coating of Pt and NiO thin overlayersfollowed by heat treatment at 600 �C. Fig. 18 shows a sche-matic illustration of a sandwich Pt–NiO–Pt tube. This struc-ture provides promising platforms for ultra-selective andsensitive detection of C2H5OH with minimal interferencefrom small-molecule gases. Approximately 6 times highersensitivity was obtained in Pt–NiO–Pt tubes compared withpure NiO tubes and NiO thin films [136]. The advantage ofthe polymeric nanofiber templating route is the multilayeredgrowth of various films. Catalyst-decorated n-type or p-typetubes and p–n–p or n–p–n junctions are possible.

Metal oxide tubes were also prepared by electrochemicalanodization of metallic foil such as Ti in an electrolyte solu-tion containing NaF (0.5 wt.%) + Na2SO4 (0.2 mol L–1) anda mixture of glycerol and water (50:50 vol.%). The lengthand thickness of the nanotubes can be controlled by vary-ing the anodization time and applied bias. Fig. 18a shows atypical TiO2 tube array [141]. Joo et al. reported on H2 gas

uced with permission from Ref. [141]. Copyright (2010), Elsevier. (b) TEMn from Ref. [142]. Copyright (2011), Wiley.


sensors of Pt- and Pd-added anodic TiO2 nanotube filmssynthesized by anodization of Pt–Ti and Pd–Ti alloy thinfilms [99]. These dispersed Pt or Pd particles improvedthe performance of the hydrogen gas sensor presumablydue to accelerated hydrogen chemisorption on the nano-tube wall [99]. Al2O3 tubes prepared by the transformationof NH4Al(OH)2CO3 nanotubes (synthesized via homoge-neous precipitation with assistance of surfactant) annealedat 600 �C were used for humidity sensors [100]. The tube-like nanostructures not only increase efficient sites for gasadsorption, but also promote the dissociation of waterabsorbed onto the surfaces of the nanotube walls. In addi-tion, sacrificial 1-D structures, which can be removed by

Fig. 19. (a) SEM image of ZnO nanofibers prepared by electrospinning; (b) mainvestigation of ZnO nanofibers using TEM; (d) magnified image of (c) for la(2007), American Institute of Physics. (e) TEM image of pure SnO2 nanofiber[110]. Copyright (2010), Wiley.

subsequent etching or high-temperature calcination, havealso been utilized for the synthesis of hollow tubes. 1-Dcoaxial nanocables were used as sacrificial templates forthe fabrication of SnO2 nanotubes loaded with Ag2O nano-particles [142]. Fig. 18b shows Ag nanoparticle-embeddedSnO2 tubes [142].

4.1.3. Nanofibers

Among the different strategies for producing sensingdevices based on nanosized 1-D building blocks, electros-pinning provides several unique advantages including easeof fabrication and robustness due to fiber network redun-dancy. Several recent reports emphasize the importance

gnified image of (a) for investigation of surface morphology; (c) structuralttice inspection. Reproduced with permission from Ref. [147]. Copyright

and (f) Pd-loaded SnO2 nanofiber. Reproduced with permission from Ref.

Fig. 20. The electrical response (I/I0) of unloaded and 30 mol.% Pd-loaded SnO2 sensors to H2 with concentrations in the range of 50–1000 ppb. From Ref. [110]. Copyright (2010), Wiley.


of electrospinning routes for chemical sensor fabrication[143–146]. In particular, nanofibrous metal oxide layersdisplay unique polycrystalline morphologies, i.e. nanofi-bers are composed of tiny nanoparticles that give rise tohigh surface areas and large surface-to-volume ratios, lead-ing to effective gas modulation of the electrical resistance.Sol–gel solutions, containing inorganic precursors andhigh-viscosity polymers, can be electrospun on sensor sub-strates, i.e. interdigitated electrode, to collect the inorganiccomposite fibers. During the electrospinning process, reac-tions such as hydrolysis, condensation and gelation of theprecursors are involved, leading to the morphological andmicrostructural evolution of the fibers. Subsequent calcina-tion results in the decomposition of the organic compo-nents while the inorganic precursors oxidize andcrystallize to form metal oxide nanoparticles aligned alongwhat used to be the fibers in the as-spun state. This processresults in highly porous structures that often display abimodal pore size distribution with relatively large pores(submicron to a few microns) and small nanosized pores.The large pores serve to facilitate fast gas diffusion intoand out of the nanofiber mats, while the nanopores mark-edly enhance the active surface area for the interaction withthe gas species. Both features have important merits forgas-sensing applications, as the former ensures fast sensorresponse and recovery and the latter is essential for highgas sensitivity. A rich variety of metal oxide nanofiberssuch as simple binary oxides (i.e. ZnO, TiO2, SnO2, NiOand CuO), complex oxides (i.e. BaTiO3, LaNiO3, PbTiO3,BiFeO3, Bi3.15Nd0.85Ti3O12), and two-phase mixtures ofNiO/ZnO and other couples have been synthesized forapplication in chemical sensors [120,142].

In order to tailor the microstructure and functionalproperties of chemical sensors utilizing nanofiber mats,the main process parameters should be carefully and sys-tematically controlled. These include: (1) the chemistry ofthe electrospun solution; (2) the inorganic precursor chem-istry; (3) the rheology of the electrospun solution(determined by the polymer/solvent ratio); (4) the electros-pinning time (controlling the layer thickness); (5) the tem-perature and pressure during the hot-pressing step; and(6) the calcination temperature–time profile (controllingthe crystallization process). These parameters should bemanipulated to optimize the preparation process for differ-ent sensor materials [143,144].

As one example, Fig. 19a and b show field emission(FE)-SEM images of ZnO nanofiber mats after calcinationat 500 �C. A random distribution of nanofibers is obtainedwith no distinct alignment. The randomness is due to theinstability of the electrospinning jet. The nanofibers shrinkfollowing the calcination process due to the removal of thepolymer component (in this case, polyvinylpyrrolidone/poly(methyl methacrylate) bi-component). The diametersof the ZnO/polymer composite nanofibers and the calcinedZnO nanofibers were approximately 1000–1200 and 700–1000 nm, respectively. The network structure was wellmaintained after the calcination step with a high degree

of porosity. The surfaces of the as-spun composite nanofi-bers are smooth. In contrast, the surfaces of calcined ZnOnanofibers show a distinct polycrystalline character asshown in Fig. 19b. Fig. 19c and d show high-magnificationtransmission electron microscopy (TEM) images of theZnO nanofiber. The ZnO nanoparticles, with diameter of20 nm, were close-packed to form a nanofiber. Such struc-tural characteristics of the nanofiber results from the sol–gel reaction accompanied by subsequent thermal treatmentthat induces the creation and growth of the ZnO nuclei inthe nanofiber [147].

Various SMO nanofiber mats, functionalized with cata-lysts, exhibit unique morphologies, facilitating efficient gastransport into the layers, combined with high surface areaand enhanced reactivity, well suited for high sensitivity.Catalyst-loaded metal oxide nanofiber sensor prototypescan be fabricated by means of a modified electrospinningprocess that involves electrospinning of polymer fibers con-taining inorganic and catalyst precursors that subsequentlyform the catalyst-loaded metal oxide fibers. Fig. 19e and fshow TEM images of microstructures of pristine SnO2 andPd-loaded SnO2 nanofibers. The very fine nanosizeddimensions, achieved by grain boundary pinning via intro-duction of crystal growth inhibitors, are desirable forenhanced sensing performance. Pd additions reduced theaverage crystallite size comprising SnO2 fiber mats (cal-cined at 600 �C) from �15 to �7 nm [110]. This demon-strates that catalytic Pd works very effectively in reducinggrain growth in electrospun SnO2 fiber mats, probablydue to the suppression of grain boundary migration andelevation of the energy barrier for grain growth. Fig. 20illustrates the exceptional sensivity to hydrogen gasachieved by the addition of the Pd catalyst to the SnO2

fiber sensor.


Although SMO nanofibers are often considered as idealbuilding blocks for high-sensitivity gas sensors, poor sub-strate adhesion, as well as difficulties with achieving reliableelectrical contacts to the fibers, are key issues that must becarefully addressed to enable reproducible sensor fabrica-tion using the electrospinning method. Thus far, severalapproaches have been introduced to overcome these barri-ers, including (1) the introduction of a hot-pressing inter-mediate step following the electrospinning process andprior to the final calcination step [148]; or (2) drop-casting(dripping) and/or deposition of chopped metal oxidenanofibers, so-called “nanorods”, dissolved in solvents.As illustrated in Fig. 21a, the hot-pressing step markedlyimproves the adhesion of the electrospun fibers to the

Fig. 21. (a) Schematic illustrating hot-pressing method; (b) SEM image of calcof (b) for surface morphological inspection. Reproduced with permission fromcoating method; (e) SEM image of the drop-coated sensing layer on the sendistribution (unpublished work).

substrate and the underlying (embedded) electrical con-tacts, and at the same time, as recently discovered, breaksthe outer layers of the fibers, exposing their fine fibrillarinner structure following the calcination step. This resultsin a unique structure comprising nanosized fibrils alignedalong what used to be the polymer fibers following the elec-trospinning process and before the calcination step [148].Fig. 21b and c shows the microstructures of calcinedSnO2 nanofiber mats, followed by hot-pressing of theinorganic precursors/polymer composite nanofibers. Thehot-pressing step serves to drive the PVAc above its glasstransition temperature (28–30 �C), resulting in markedlyimproved adhesion to the substrate. Subsequent calcina-tion at 450 �C for 30 min results in porous structures with

ined SnO2 nanofiber mats following hot-pressing step; (c) magnified imageRef. [148]. Copyright (2010), Springer. (d) Schematic illustration of drop-sor substrate; (f) magnified image of (e) illustrating bi-modal pore size


high surface area (�73.5 m2 g–1, measured by BET) andsmall grain size (�5–15 nm) as shown in Fig. 21c [148].

In the case of drop-casting of nanofibers (or nanorods),ultrasonication or ball-milling can be used to disperse themetal oxide nanofibers in the solvents. During sonicationand ball-milling, the nanofibers are dispersed and/orchopped, resulting in the formation of “nanofiber ink” or“nanorod ink” (Fig. 21d). As shown in Fig. 21e and f, thesechopped nanorods are still composed of tiny nanoparticleswith a bimodal pore size distribution, with micrometricpores between what used to be the rods and nanometricpores between the nanofibrils [104].

4.2. Hollow nanostructures

Hollow nanobuilding blocks with thin-walled structuresare considered as ideal sensing layers for high-sensitivitychemical sensors [149–151]. The electron depletion layerthickness is known to be 10–25 nm. Hollow spheres andhemispheres with shell thickness of <25 nm can thus pro-vide significantly enhanced gas sensitivity as compared todense planar thin films with the potential for an insensitiveinterfacial layer to form between film and substrate [152].

Fig. 22. (a) Schematic fabrication procedure for fabrication of hollow SnO2 hfabricated hollow TiO2 hemispheres, reproduced with permission from Ref. [1arrays, reproduced with permission from Ref. [87]. Copyright (2009), Royalhemispheres, reproduced with permission from Ref. [155]. Copyright (2006), Afrom substrate and flipped over, reproduced with permission from Ref. [87]. Cmultilayered TiO2 hemispheres on Si substrate, (g) larger-diameter TiO2 hemi

In addition, small nanopores enable gas penetration intoinner layers, leading to active gas–oxide interactions inboth outer and inner layers.

Both polymeric beads, which are pyrolyzed during high-temperature calcination, and metallic particles, which aredissolved in acidic solution, have been used as hard tem-plates in the fabrication of hollow nanostructures. Flexibledimensional control can be achieved by manipulating thediameter of the polymer beads or metallic particles. Inparticular, repetitive adsorption of precursors containingcations on polymer beads, so called layer-by-layer (LBL)coating, and subsequent high-temperature heat treatment,leads to hollow metal oxide spheres with precisely con-trolled wall thickness [153,154].

Recently, these colloidal templating routes have beensuccessfully combined with physical or chemical vapordeposition methods. Kim et al. demonstrated the potentialuse of macroporous films prepared via polymeric bead tem-plating combined with sputter deposition of various metaloxides such as CaCu3Ti4O12 [155], SnO2 [87], TiO2

[156,157] and NiO [86]. Fig. 22a illustrates the synthesismethod for RF-sputtered SnO2 coating on polymethylmethacrylate (PMMA) sacrificial templates, followed by

emispheres using PMMA sacrificial templates; (b) cross-sectional view of57]. Copyright (2008), Wiley. (c) Tilted view of hollow SnO2 hemisphereSociety of Chemistry. (d) Cross-sectional view of hollow CaCu3Ti4O12

merican Chemical Society. (e) Inside view of SnO2 hemispheres removedopyright (2009), Royal Society of Chemistry. (f) Cross-sectional image of

spheres (2 lm).


high-temperature heat treatment, leading to hollow hemi-sphere structures. As shown in Fig. 22b and d, TiO2

[157], SnO2 [87] and CaCu3Ti4O12 [155] hemispheres couldbe obtained by this method. Fig. 22e provides a back-sideview of SnO2 hemispheres. A wall thickness of 18.2 nmwas achieved via precise control of film thickness. In addi-tion, multilayered TiO2 hemispheres could be built up byrepeating the overall coating process (Fig. 22f). Flexibledimensional control is obtained by changing the size ofthe polymeric microspheres. As shown in Fig. 22g, TiO2

hemitubes, 2 lm in diameter, were prepared by use ofapproximately 2 lm sized PMMA beads. Fig. 23 showsthe NO2 gas response of SnO2 planar thin films andSnO2 hemitube networks. SnO2 hollow hemispheres, syn-thesized with the aid of PMMA microsphere templates,exhibited an approximately 3-fold enhancement in NO2

gas response compared to their non-templated counter-parts due to the increased surface area and enhanced car-rier depletion of the thin hollow SnO2 shells.

In these approaches, catalyst particles can be decoratedonto the top of various metal oxide hemispheres by co-sputtering of the catalyst, e.g. Pt, and the metal oxide tar-gets [88]. The coating processes are conducted at room tem-perature to prevent thermal deformation of the polymericbeads. For example, co-sputtered layers (10 nm) of Ptand SnO2 were deposited following the deposition of thesupport SnO2 layer (90 nm). After calcination above500 �C, catalytic Pt particles were decorated onto the sup-port metal oxide layers (10 nm). Pt-decorated SnO2 hemi-spheres exhibited 4-fold higher NO2 gas responsecompared to catalyst-free SnO2 hemispheres [88].

Besides polymeric beads, nanosized metallic particlessuch as Ni (300 nm) can be used as a hard template forthe coating of thin hydrous Sn(OH)2 layers (30 nm) [82].After calcination at 400 �C, hydrous Sn(OH)2 layers are

Fig. 23. The response (I0/I) during cyclic exposure to increasing NO2

concentration in dry air at 250 �C. Inset: schematic illustration of hollowSnO2 hemispheres and response to NO2 adsorption (inset). Reproduced withpermission from Ref. [87]. Copyright (2009), Royal Society of Chemistry.

converted into crystalline SnO2 shells and the outer partof the Ni particles become oxidized to NiO. Due to thedilation induced during the oxidation step between metallicNi and NiO, nanopores are generated in the SnO2 shells(Fig. 24a). Through these open accessible pores, the innernon-oxidized Ni can be removed by etching in diluteHCl. Finally SnO2 hollow spheres with functionalizedinner NiO layers are obtained (Fig. 24b). Lee et al. reportedvery fast response (5 s) and recovery (2 s) speeds due to theenhanced catalytic activity combined with the open nano-porous hollow structures (Fig. 24) [82].

Soft templating routes have also attracted much atten-tion. Mixed solutions containing cation precursors and

Fig. 24. (a) SEM image of NiO-functionalized hollow SnO2 spheres usingNi templates and a schematic illustration of a NiO-functionalized hollowSnO2 sphere. (b) Dynamic C2H5OH resistance transient measured forsensor prepared from NiO-functionalized hollow SnO2 spheres at 450 �C.Reproduced with permission from Ref. [82]. Copyright (2010), RoyalSociety of Chemistry.

Fig. 25. (a) TEM image of a zigzag Zn2SnO4 NW; (b) IDS–VDS curvesmeasured at different gate voltages in a single Zn2SnO4 NW FET.Reproduced with permission from Ref. [165]. Copyright (2011), RoyalSociety of Chemistry.

Table 1FET-based chemical sensors.

Materials Synthesis method Wire diameter(nm)

Analytes Operatingtemperature

Detectionrange

Response time(s)

ZnO [167] Thermal evaporation of ZnOpowders

– CO 275 �C 50–400 ppm 500

CuO [168] Thermally grown on copperfoils

50–100 CO 200 �C 5–1200 ppm <10

In2O3 [166] Chemical vapor deposition 30–100 H2S Room temp. 1–160 ppm 48Pd/Si [169] Vapor–liquid–solid (VLS) 30 H2 Room temp. 1–1000 ppm <100Zn-doped In2O3

[162]Chemical vapor deposition 50–300 CO Room temp. 1–5 ppm 20

ZnO [170] Chemical vapor deposition 80 CO Room temp. 400–1600 ppm

–

In2O3 [171] Vapor–liquid–solid (VLS) 165 NO Room temp. 5 ppm 160SnO2 [163] Chemical vapor deposition 30–80 Acetone

C2H5OH192–373 �C 20–80 ppm –

In2O3 [172] Laser ablation 10 NO2 Room temp. 5–500 ppb 5–10In2O3 [173] Laser ablation 10 NO2, NH3 Room temp. 5–20 ppm

NO2

20

In2O3 [174] Chemical vapor deposition 10 NH3 Room temp. 0.002–1%NH3

30

In-SnO2 [175] Thermal evaporation 70–150 Ethanol 400 �C 10–1000 ppm <2SnO2 [176] Electrochemical method 60 O2, CO 300 �C C 0.05% CO –ZnO [177] Thermal evaporation 20 NH3 Room temp. 4–1000 ppm �5


surfactants are loaded into an autoclave and reacted (mildtemperature and high pressure), followed by high-tempera-ture calcination, leading to the formation of hollow nano-structures. The morphologies, microstructures anddiameters are largely governed by surfactants and the con-centration of precursors, as well as processing parameters(temperature and pressure). For soft templating synthesis,surfactants such as glucose [85,149,158,159] and lysine[84] have been utilized for hydrothermal and solvothermalapproaches.

5. Field effect transistor-based chemical sensors

SMOs have been used as the active channel layer in fieldeffect transistors (FETs). Field-induced current modulationcan lead to further improvements in chemical sensing char-acteristics. NW-based FETs using ZnO [160,161], In2O3

[114,162], SnO2 [163,164] and Zn2SnO4 [165] have beenreported with superior NO2, CO and VOC sensing behav-ior. Although FET-based sensors using single NW chan-nels have very good sensing performance, pick-and-placemethods for positioning of the channel layer, connectingsource and drain, induce large variations in channel length,leading to large signal variations.

Fig. 25a shows a TEM image of a zigzag Zn2SnO4 NWsynthesized by thermal evaporation. FETs, using single zig-zag Zn2SnO4 NWs, exhibited an on/off ratio of 104 and adevice mobility of 17.2 cm2 (V s)�1 (Fig. 25b). The currentincreased over 10 times with UV light irradiation, indicat-ing possible applications in photoswitches. As summarizedin Table 1, single NW-based FETs exhibit superior sensing

properties. Chemical sensors based on an individual In2O3

NW transistor were able to detect H2S down to the 1 ppmlevel at room temperature [166]. More recently multiplechannel networks of NWs have been intensively studied,offering more uniform properties based on average channelcharacteristics [163–165]. Dattoli et al. reported on the suc-cessful demonstration of a highly selective SnO2 NW net-work-based FET via combined temperature and gatevoltage modulation. Fig. 26 shows a schematic illustrationand SEM micrograph of a NW microhotplate sensor. A Ptmeander heater is embedded in the nitride membrane,while the NW sensor is located on the membrane top sur-face. As shown in Fig. 26e and f, FETs with SnO2 NW net-works exhibit effective modulation upon variation oftemperature and gate bias. Table 1 summarizes recent pro-gress on FET-based chemical sensors.

6. Future directions

The authors have attempted to demonstrate that thefield of solid-state gas-sensitive devices is a very dynamicand fast evolving field, and that it can be expected to con-tinue to grow rapidly with advances in materials and devicefabrication, characterization and modeling. Under thesecircumstances, it is very likely that the authors will misscertain important future trends. Nevertheless, it is usefulto consider which developments may deserve carefulscrutiny.

A key challenge for chemoresistive sensors, which relyon chemisorption, is their inherent lack of selectivity.Reducing gases contribute to decreased resistance of n-type

Fig. 26. (a) Schematic view of sensor fabricated on a nitride membrane-based microhotplate platform. (b) Optical image of the sensor device. (c)Electrical/thermal simulation of sensor. (d) SEM images of the NW sensor. (e) Traces, left axis, sensor conductance measurements taken at three differenttemperatures during the validation run (Vgs = 0 V). Bars, right axis, analyte flow schedule for a single exposure run. (f) Sensor conductance taken at fivedifferent Vgs conditions during validation run (Theater = 373 �C). Reproduced with permission from Ref. [163]. Copyright (2012), Royal Society ofChemistry.


semiconductors, while oxidizing gases lead to increasedresistance. It therefore becomes of interest to identifymeans for modulating the device in such a way that selec-tivity is somehow advanced. One promising approach isvia the field effect. CMOS-compatible field effect gas sensordevices were recently characterized and modeled by Velas-co-Velez et al. [178]. They calculated that for a gas-sensinglayer thickness in the range of the Debye length, the

controlling electrical field reaches the surface, enablingeffective control of gas sensitivity. Whitfield and Tullerexamined a thin film transistor with an amorphous InG-aZnO4 channel both experimentally and by numerical sim-ulation [179]. Gate bias was found to have a significantimpact on sensitivity, particularly as the channel thicknessdropped below 100 nm with positive gate bias acceleratingchemisorption of oxidizing gases and negative gate bias


accelerating desorption. This points to the possibility ofdesigning chemoresistive sensors with improved sensitivity,selectivity and response time.

Above band gap illumination has been shown tocontribute to the desorption of oxygen and NO2 fromthe surface of n-type semiconductors [180,181].This promises to increase sensor reversibility andreduce response time under reduced temperatureoperation. Sub-band gap illumination may hold outpromise for selective desorption and thus improvedselectivity.

Processing routes that offer great flexibility in control-ling the morphology of nanostructured SMOs promise toaccelerate our understanding of the precise role of crystal-lographic orientation, interconnectivity, grain boundaries,particle-to-particle contact, scale, etc. In a recent study, sin-gle-phase nanostructured p-type CuO was prepared bymicrowave-assisted synthesis leading to different morphol-ogies, including hierarchical urchin-, fiber- and nanorodstructures achieved by modification of the solvent utilizedin hydrothermal reactions [182]. These materials wereexamined simultaneously in the same sensor chamber andthe urchin-like structures were found to be most effectivefor hydrogen detection in the range of ppb, and consider-ably superior to previously reported hydrogen CuOsemiconducting oxide-based sensor response. Based on aconsideration of the model proposed by Barsan et al.[183], the reduced contact area between particles with theurchin-like morphology is the likely source of the enhancedresponse.

Power consumption and the resultant heating of nearbycomponents remains a major hindrance for integrating gassensors into microelectronics and portable devices. Micro-hotplate structures, as illustrated in Fig. 26, certainly makeintegration more feasible, but at the cost of more complexprocessing. Recent work by the authors demonstrates theability to self-heat nanocolumnar metal oxide thin filmsto hundreds of degrees Celsius while requiring only a fewtens of microwatts of power [184]. For NO2 and ethanol,the response of the 360 nm thick WO3 films was threeorders of magnitude higher than for that of dense planarfilms. The exceptional response of the nanocolumnar sen-sor is attributed to the combined effects of self-heating, aporous nanostructure with high surface-to-volume ratio,and the presence of narrow necks between the columns.These features, coupled with transparency, suggest thatsuch chemical sensor structures will lend themselves tomore effective integration into microelectronic devices.

Acknowledgements

H.L.T. thanks the following organizations for supportof his research on topics related to this work: National Sci-ence Foundation, Division of Materials Research, Materi-als World Network (DMR-0908627), Korea ResearchCouncil Industrial Science and Technology (B551179-10-01-00), The Brazil-MIT program, Department of Energy,

Basic Energy Sciences (DE SC0002633). I.D.K. acknowl-edges the support by the Engineering Research Center(ERC-N01120073) program from the Korean National Re-search Foundation. I.D.K. and A.R. acknowledge the sup-port of the Korean Ministry of Research (GrantN01120137) and the Israeli Ministry of Science and Tech-nology (Grant No. 3-8272). This work was supported bythe Center for Integrated Smart Sensors funded by theMinistry of Education, Science and Technology as GlobalFrontier Project (CISS-2012M3A6A6054188). A.R. andH.L.T. acknowledge the support of the US–Israel Bina-tional Science Foundation (BSF Grant No. 2006295).

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Advances and New Directions in Gas-sensing Devices