Morphology and Gas-Sensing Properties of Tin OxideFoams with Dual Pore Structure

KYUNGJU NAM,1 HYEONG-GWAN KIM,2 HYELIM CHOI,1,6

HYEJI PARK,1 JIN SOO KANG,3,4 YUNG-EUN SUNG,3,4

HEE CHUL LEE,2,7 and HEEMAN CHOE1,5

1.—School of Materials Science and Engineering, Kookmin University, 77 Jeongneung-ro,Seongbuk-gu, Seoul 02707, Republic of Korea. 2.—Department of Advanced Materials Engineer-ing, Korea Polytechnic University, Siheung 15073, Republic of Korea. 3.—Center for NanoparticleResearch, Institute for Basic Science (IBS), Seoul 08826, Republic of Korea. 4.—School of Chemicaland Biological Engineering, Seoul National University, Seoul 08826, Republic of Korea.5.—Cellmotive Co. Ltd., #518, Engineering Building, Kookmin University, Seoul 02707, Republicof Korea. 6.—e-mail: choihl@kookmin.ac.kr. 7.—e-mail: eechul@kpu.ac.kr

Tin oxide is a commonly used gas-sensing material, which can be applied as ann- or p-type gas sensor. To improve the gas-sensing performance of tin oxide,we successfully synthesized tin oxide foam via an ice-templating or freeze-casting method. The tin oxide foam samples showed different morphologicalfeatures depending on the major processing parameters, which include sin-tering temperature, sintering time, and the amount of added powder. Basedon scanning electron microscopy images, we could identify dual pore structureof tin oxide foam containing ‘wall’ pores ranging from 5.3 lm to 10.7 lm, aswell as smaller secondary pores (a few micrometers in size) on the wall sur-faces. Gas-sensing performance tests for the synthesized tin oxide foams re-veal a sensitivity of 13.1, a response time of 192 s, and a recovery time of 160 sat an ethanol gas concentration of 60 ppm at 300�C. This is a remarkableresult given that it showed p-type semiconductor behavior and was usedwithout the addition of any catalyst.

Key words: Tin oxide, gas sensing, foam, porous ceramic, sensor

INTRODUCTION

With the dramatic advances being made in mod-ern technology, numerous toxic gases are beingproduced, and there is growing concern aboutaccidental leakages. These gases can cause variousforms of air pollution that can seriously affect theenvironment and the health of both humans andanimals.1–3 Technical difficulties arise because mostof these toxic gases are present at relatively lowconcentrations. Therefore, it has become increas-ingly important to develop sensitive detectiondevices in order to effectively monitor the varietyof toxic gases with appropriate gas sensors, whichcan detect the presence of a toxic gas or a high

concentration of gas particles in the surround-ings.4,5 There are various types of gas sensors, andthe ceramic gas sensor is one of the most commonlyused. The ceramic gas sensor is based on theprinciple of the change in electrical resistanceresulting from the gas adsorption reaction on thesurface of sensor material exposed to the targetgas.6,7 It is very important to select an appropriatesensing material that shows a sensitive response tothe target gas.6,7

A variety of oxide-based gas-sensing semiconduc-tor materials, e.g., SnO2,7–9 ZnO,10 In2O3,11 WO3,12

TiO2,13 NiO,14 CuO,15 Co3O4,16 and Cr2O3,17 havebeen extensively studied as chemical sensing mate-rials, which are categorized into n-type and p-typegas sensors. Over the past few decades, n-type gassensing materials have been receiving more atten-tion because they generally exhibit more satisfying

(Received September 30, 2016; accepted December 15, 2016)

Journal of ELECTRONIC MATERIALS

DOI: 10.1007/s11664-016-5242-6� 2017 The Minerals, Metals & Materials Society

sensitivity and selectivity than p-type gas sensingmaterials.18,19 Nevertheless, the scientific impor-tance of p-type oxide semiconductors as chemiresis-tive materials should not be underestimated, giventhat a number of p-type oxide semiconductors suchas NiO, CuO, Cr2O3, Co3O4, and Mn3O4 have beenextensively used as decent catalysts to promote theselective oxidation of various volatile organic com-pounds (VOCs).19–23 From this perspective, p-typeoxide semiconductors are also promising materialplatforms for developing new functionalities inchemiresistors.19–23

The most representative sensing materials areSnO2 and ZnO, both of which exhibit n-type oxidesemiconductor-sensing behavior. Other n-type oxidesemiconductors such as TiO2, WO3, In2O3, andFe2O3 are also being widely explored to discovernew chemiresistivity functionalities due to theirgood response and selectivity to ambient conditions,as well as the simplicity of their material synthesisand sensing device fabrication.19,24 The markedlydifferent gas-sensing characteristics of n- and p-type oxide semiconductors must be understood inthe context of the functions of the receptor, conduc-tion paths, and gas-sensing mechanisms of specificmaterials that have different majority charge carri-ers.19 It is noted that tin oxide can be used as eitheran n- or p-type gas-sensing material. In otherwords, SnO2 is a typical n-type semiconductor witha wide band gap, high conductivity, and goodstability, whereas SnO is a p-type semiconductor.25

It is important to understand the microstructure ofthe sensing material because the grain sizes, mor-phology, surface state, and accessible surface area allplay an important role in its gas-sensing perfor-mance.25,26 Moreover, the adsorption/desorption pro-cesses take place mainly on the surface of the sensinglayer and determine the sensor response.26 Therefore,controlling the (surface) morphology of a gas-sensingmaterial has been proven to be an effective method forenhancing the sensing properties.26 Recently, therehas been a dramatic increase in research on the designof gas-sensing materials, including nanowires,27

nanoparticles,28 nanorods,29 nanotubes,30 porousmaterials,31,32 and hollow spheres.33,34

In this paper, tin oxide foam is created for itspotential use as a gas-sensing material by a simple,scalable method known as freeze casting. This paperalso reports on the ethanol-sensing properties of thetin oxide foam with various ethanol gas concentra-tions, because ethanol is often known as a majorbreakdown product of foodstuffs where bacteria orfungi grow and also as the most common interferinggas.35–37 We carried out x-ray diffraction (XRD), x-ray absorption near edge structure (XANES) anal-ysis, x-ray photoelectron spectroscopy (XPS) analy-sis, and morphological analysis to characterize thephysical and chemical properties of freeze-cast tinoxide samples and relate the properties to the gas-sensing performance.

EXPERIMENTAL METHODS

Tin oxide (SnO2) powders (15.5 g, with a meanparticle size of<80 nm and a purity of 99.9%;Inframat Advanced Materials, Manchester, CT,USA) were suspended in a 20-ml solution of deionizedwater containing polyvinyl alcohol binder (0.47 g, 3.0wt.%, mean molecular weight of 89,000–98,000, anda purity of 99%; Sigma-Aldrich, MO, USA). Addi-tionally, to improve the densification of the SnO2

powders, we added copper oxide (CuO) powders(0.08 g, 0.5 wt.%, with a particle size of 40-80 nmand a purity of 99.9%; Inframat Advanced Materials)as a sintering agent. The slurry was then dispersedby stirring for 30 min and then by sonication for 1 h.To ensure sufficient particle dispersion, the processwas repeated twice. The slurry was then cooled to afew degrees above the freezing point of water andpoured into a cylindrical mold consisting of insulatedTeflon walls (with 43 mm and 56 mm inner and outerdiameters, respectively, and a height of 70 mm) on aCu rod, which was cooled using liquid nitrogen andcontrolled using a thermocouple and temperaturecontroller. Once the freezing process was complete,the frozen SnO2-ice sample was removed from themold and dried at 80�C for 48 h in a freeze-dryerunder 10�2 torr residual pressure. The resultinggreen-body foam sample was then heat-treated intwo separate steps. It was first subjected to a pre-heat treatment at 300�C for 3 h to remove the binderprior to sintering at elevated temperature in a boxfurnace in atmospheric air. To determine the effect ofthe sintering temperature on the foam, we usedsintering temperatures of 1050�C, 1100�C, 1150�C,and 1250�C for 10 h with heating and cooling rates of5�C min�1. We then controlled the sintering times to6 h, 10 h, and 24 h to determine the effect of thesintering time on the morphology of SnO2 foam byfixing the temperature at 1150�C, because the heattreatment at 1050�C or 1100�C resulted in somewhatincomplete sintering (with some degree of brittle-ness) and the heat treatment at 1250�C resulted inexcessive sintering (with less surface area available).

We observed the microstructure of the tin oxidefoams by scanning electron microscopy (SEM;JSM7401F; JEOL, Tokyo, Japan), and measuredthe mean pore sizes and strut widths from the SEMimages. We further characterized the selected foamsusing x-ray diffraction (XRD; DMAX2500; Rigaku,Tokyo, Japan) and x-ray absorption near edgestructure [XANES, in fluorescence mode at the 10C beamline of Pohang Accelerator Laboratory(PAL), Republic of Korea]. During the XANESmeasurement, the beam was monochromatized witha Si (111) monochromator, and the intensity wasdetuned by 20% to inhibit undesirable higherharmonics. Additionally, x-ray photoelectron spec-troscopy (XPS; Thermo SIGMA PROBE; ThermoFisher Scientific, MA, USA) analysis was also usedto investigate the chemical state of the tin oxidefoam’s surface. The pore distribution was measured

Nam, Kim, Choi, Park, Kang, Sung, Lee, and Choe

by mercury intrusion porosimetry (MIP; AutoPoreIV 9520; Micromeritics, GA, USA).

More than ten random struts were selected oncross-sectional SEM images to obtain the meanstrut size. Likewise, the mean pore size was alsomeasured by selecting more than ten random poreson top-view SEM images. Their porosity was mea-sured by comparing our image analysis results fromthe SEM images with the results obtained by theArchimedes method.

To evaluate the gas-sensing performance of theSnO2 foam sample, two Pt electrode pads werepatterned with a spacing gap of 200 lm by a lift-offprocess using a shadow mask. The gas sensitivitywas then measured at five different ethanol gasconcentrations from 20 ppm to 60 ppm at a fixedtemperature of 300�C. During the measurements,we balanced and controlled the gas variations by amass flow controller, and injected gas into thesensor chamber while continuously measuring theelectrical resistance of the sensor at a constantcurrent of 1 lA. The gas sensitivity was defined asthe resistance ratio of Rg/Rair, where Rair and Rg

are the electrical resistances in pure air and thetarget gas, respectively.

RESULTS AND DISCUSSION

Microstructural Observation

We used a two-step heat-treatment process toachieve a better sintering result than is possibleusing a single-step heat-treatment process asdescribed previously.38 It is well known that SnO2

powders are difficult to densify by sintering becausethe mass transport mechanism in SnO2 is rathercomplex. It is controlled by surface diffusion at lowtemperatures (T< 1000�C) but by evaporation con-densation at high temperatures (T> 1300�C).39

When either of these mechanisms predominates,densification in SnO2 is difficult to achieve.

To enhance the densification of the SnO2 powders,we added a small amount of CuO powder (1.0 mol.%or 0.5 wt.%) to the SnO2 slurry. CuO helps thedensification process by playing a significant role inpromoting bulk and/or grain-boundary diffusion orforming a liquid phase in the sintering of SnO2.40

Figure 1 shows digital photographs of a represen-

tative SnO2 foam before (left, 45 mm in diameterand 16 mm in height) and after (right, 31 mm indiameter and 11 mm in height) the sintering pro-cess in air. The volumetric shrinkage of this foam isDV/V �67%, and the radial and height shrinkagesare roughly Dr/r �31% and Dh/h �31%, respec-tively. Therefore, we consider the shrinkage to benearly isotropic, which suggests that the residualstress due to shrinkage is minimal.

Figure 2a compares the XRD patterns of the fourtin oxide foam samples heat-treated at 1050�C,1100�C, 1150�C, and 1250�C for 10 h. Based on theXRD pattern analysis results, we confirmed thatthere were no noticeable differences between theXRD peaks of the four foam samples, and they allshowed the typical peaks of SnO2. No CuO peakswere observed, most likely because the normaldetection limit of laboratory x-ray source is around2–3 wt.%.41 The mean crystallite size was calcu-lated using the Scherrer formula. As previouslyreported for the calcination process of ceramicnanoparticles,42,43 the crystallite size graduallyincreases from 51 nm to 71 nm with increasingheat-treatment temperature from 1050�C to1250�C.

The oxidation state of Sn in the tin oxide foamwas characterized by obtaining the Sn K-edgeXANES spectra (Fig. 2b), and comparing theabsorption edge position of the tin oxide foam andthose of the commercial SnO and SnO2 powders(purchased from Sigma-Aldrich, USA). XANES datawere processed by following a standard procedureusing the Athena interface to the IFFEFIT pro-gram.44 The XANES analysis results revealed thatthe average Sn valency is 3.36 in the tin oxide foamsample, suggesting that the phase of tin oxide in theentire tin oxide foam sample is predominantlySnO2. Furthermore, we also carried out XPS mea-surements to investigate the chemical state of thetin oxide foam on its surface. From the XPS surveyspectrum in Fig. 2c, the presence of both Sn and Owas clearly confirmed. On the other hand, we couldhardly observe any signal from Cu, suggesting thata negligible amount of CuO exists on the surface inaccordance with the XRD result. Additionally, thevalence state of Sn was characterized by the Sn 3dspectrum displayed in Fig. 2d. Based on the bindingenergy positions of the Sn 3d peaks (486.3 eV for3d5/2 and 494.7 eV for 3d3/2), we could verify that thesurface of the tin oxide foam is mainly composed ofSnO.45 According to the Sn-O phase diagram, SnOis in a stable phase in the temperature range of100–270�C under sufficient atomic percent ranges ofO.46 Thus, we assume that SnO may have beenformed either during the first heat-treatment stepapplied at around 300�C or in the slow furnacecooling process. The content of crystalline SnO inthe tin oxide foam sample is assumed to be less than3 wt.%, because the detection limit of XRD analysisusing laboratory x-ray source is known to be around2–3 wt.%.41

Fig. 1. An ice-templated SnO2 foam before and after sintering,which shows considerable volume shrinkage without cracking orwarping.

Morphology and Gas-Sensing Properties of Tin Oxide Foams with Dual Pore Structure

Morphological Analysis

Figure 3 shows SEM images of a natural top view(perpendicular to the freezing direction), whereasFig. 4 shows those of the fracture surface (parallelto the freezing direction) for the tin oxide foamscreated by varying the sintering temperature from1050�C to 1250�C while keeping the sintering timefixed at 10 h. In both the natural top-view (Fig. 3)and the fracture images (Fig. 4), we observednumerous round-shaped struts, as would be seenin individual powders. This is perhaps because it isgenerally more difficult to sinter ceramics thanmetals due to their high melting points. As seen inthe series of SEM images in Figs. 3 and 4, thevariation in sintering temperature certainly influ-enced the microstructure of the tin oxide foamsamples. For example, as also confirmed in Fig. 5b,when the sintering temperature was increased from1050�C to 1250�C, the mean pore size decreasedfrom 10.7 ± 2.2 lm to 5.3 ± 0.5 lm.

SEM images of the fracture surface parallel to thefreezing direction (Fig. 4) show the typical lamellarstructure wherein porous walls are formed that arealmost parallel due to the higher growth velocity inthe direction parallel to rather than perpendicularto the temperature gradient.47 Additionally, weobserved that secondary dendrites had also grownon the surface of the porous walls of the tin oxidefoam samples, essentially showing a ‘dual’ porestructure (macropores between the walls and micro-pores on the wall surfaces). On the other hand,these secondary dendrites are present predomi-nantly on only one side of the walls. This uniquetendency results from the competition between thethermal and interfacial energies. The interfacialenergy creates tilted struts on the lamellar icecrystals along the preferred orientation with respectto the macroscopic thermal gradient direction, sothe secondary dendrites form only on the upper sideof the walls.48–50

Fig. 2. (a) XRD patterns of the tin oxide foam samples heat-treated at 1050�C, 1100�C, 1150�C, and 1250�C for 10 h; (b) Sn K-edge XANESspectra of the tin oxide foam heat-treated at 1050�C for 10 h and comparison of its absorption edge position with those of commercially obtainedSnO and SnO2 powders; and XPS spectra of (c) survey and (d) Sn 3d of the tin oxide foam after sintering in air at 1050�C for 10 h.

Nam, Kim, Choi, Park, Kang, Sung, Lee, and Choe

MIP was used for the analysis of pore sizedistribution, pore area, and porosity of the tin oxidefoam sample heat-treated at 1150�C for 10 h. Fig-ure 5a shows a pore diameter distribution curve ofthe tin oxide foam, as measured by MIP. A largesharp peak of pore size distribution lies near thepore diameter of 3.4 lm, corresponding primarily tothe major lamellar wall pores, whereas a small dullpeak of pore size distribution lies near the porediameter of �2 lm, corresponding primarily to theminor dendrite pores on the surface of the walls(Fig. 4). Based on the MIP measurement result, thetotal porosity and total pore area of the Fe foams are62.7% and 0.3 m2 g�1, respectively. The pore area of0.3 m2 g�1 for this tin oxide foam is roughly anorder of magnitude higher than the pore area of0.037 m2 g�1 reported for a commercially availableCu foam (electroless Cu plated on polymer tem-plate).51 Figure 5b, c also shows the variations of

mean pore size, wall width, and porosity as afunction of sintering temperature. As expected, themean pore size decreases with increasing sinteringtemperature (Fig. 5b). Similarly, Fig. 5b shows thevariation of the tin oxide foam wall width withrespect to sintering temperature. With increasingsintering temperature, more complete sintering isachieved, resulting in a corresponding increase inthe wall width. The wall width of the tin oxide foamincreased from 2.2 ± 0.4 lm to 3.9 ± 1.0 lm with anincrease in the sintering temperature from 1050�Cto 1250�C. The increasing mean wall width tendedto reach a plateau at higher temperatures asexpected, based on the wall growth mechanism ofdiffusion (Fig. 5b). We measured and compared theporosity by two methods: image analysis and theArchimedes method. The measurements obtainedfrom both show the same tendency of decreasedporosity with increasing temperature (Fig. 5c),

Fig. 3. SEM images of radial sections of SnO2 foams sintered at various temperatures: (a) 1050�C, (b) 1100�C, (c) 1150�C, and (d) 1250�C for10 h. The arrows indicate junctions typical of the feature of ceramic powder sintering process.

Fig. 4. SEM images of fracture surfaces of SnO2 foams sintered at various temperatures: (a) 1050�C, (b) 1100�C, (c) 1150�C, and (d) 1250�C for10 h. The arrows indicate secondary dendrites on the wall struts.

Fig. 5. (a) Pore distribution of the tin oxide foam sample heat-treated at 1050�C for 10 h determined by MIP, (b) mean pore size and mean wallwidth, and (c) porosity with respect to sintering temperature at a constant slurry volume fraction of 10 vol.%.

Morphology and Gas-Sensing Properties of Tin Oxide Foams with Dual Pore Structure

which is in good agreement with the observationsmade from the SEM images in Figs. 3 and 4. Wenote that the porosity measurements by imageanalysis are roughly 24% smaller than those bythe Archimedes method, which is in agreement witha previous report,52 perhaps because the imageanalysis was carried out on 2D SEM images andbecause this method analyzed only a localizedportion of the sample unlike the Archimedesmethod.

The strut size (wall thickness) graduallyincreased with increasing sintering time, as morestruts in the tin oxide foam tended to sinter togetherto form larger struts. We estimated the mean strutsize to be around 2.3 ± 0.4 lm, 2.7 ± 0.4 lm, and3.8 ± 0.5 lm for the SnO2 foam samples for sinter-ing times of 6 h, 10 h, and 24 h at 1150�C,respectively.

In the theory of the grain growth mechanism, thegrowth of the mean grain size can be described interms of the sintering time and temperature asexpressed in Eq. 1.53 In other words, the cube of themean grain size increases linearly with sinteringtime at a constant temperature. Because weexpected the relationship between the mean strutsize and sintering time for the tin oxide foam used inthis study to follow the same general relationshipbetween the mean grain size and sintering time, weattempted to describe the relationship between themean strut size and sintering time of the tin oxidefoams using Eq. 1:

R3 � R30 ¼ Kt ð1Þ

where R is the mean strut size at a certain time tfrom the initial mean strut size, R0. K is a growth-rate constant, which is affected by the sinteringtemperature. Figure 6 shows the dependence of themean strut size on sintering time at a constanttemperature of 1150�C. Based on the behavior of the

curve, the mean strut size tended to follow the samegrowth kinetics dependence of t1/3 in good agree-ment with the relationship between the mean grainsize and sintering time.53 Additionally, we esti-mated the growth-rate constant K, which is calcu-lated as the slope of the linear fit, to be 2.5.

Gas-Sensing Properties

To assess the gas-sensing response of our tinoxide foam, we performed a gas-sensing test byexposing it to ethanol (C2H5OH) gas with concen-trations ranging from 20 ppm to 60 ppm at 300�C indry air. The working temperature of 300�C isassumed to be ideal, given that the highest sensi-tivity was achieved for zinc oxide sensor exposed toethanol gas at 300�C.54 Figure 7a shows the varia-tion in resistance (sensitivity) of the tin oxide foamsample (heat-treated at 1150�C for 10 h) withrespect to the measurement time during exposureto ethanol gas with concentrations ranging from20 ppm to 60 ppm, suggesting that the tin oxidefoam sample synthesized in this study can be usedas a suitable gas-sensing material. The graycolumns depicted in Fig. 7a show that the tin oxidefoam sample responded with satisfactory sensitivityto ethanol gas exposure with a substantial increasein electrical resistance; additionally, the resistancereturned to stable values when the ethanol gas wasturned off and pure air was injected, as illustratedby the white columns in Fig. 7a.

Interestingly, the increase in electrical resistanceduring exposure to ethanol gas implies that themajor carrier type in the tin oxide foam samplesynthesized in this study is the p-type semiconduc-tor with the majority charge carriers being posi-tively charged holes,55 whereas most metal oxidegas sensors show the n-type semiconductor-sensingbehavior by responding with negative resistancechange during exposure to gas.19 Furthermore, thedetection of dominant p-type semiconductor behav-ior indicates that the major tin oxide composition onthe surface of our tin oxide foam is SnO, becauseSnO and SnO2 are known to exhibit p-type and n-type semiconductor behaviors, respectively.56 Thisfinding is in good agreement with the XPS surfaceanalysis results in Fig. 2c and d, as the gas-sensingreaction mainly occurs on the surface of the tinoxide foam, and the dominant p-type behavior canthus be explained by SnO being the major tin oxidedetected by XPS. On the other hand, the overalloxidation state of Sn in the entire tin oxide foamsample appears to be considerably closer to theoxidation state in SnO2, based on the XANESanalysis result in Fig. 2b; the average Sn valencyis 3.36 in the whole tin oxide foam sample, indicat-ing that the oxidation state of Sn is predominantlySnO2 in the entire tin oxide foam sample. Theoperational difference and underlying mechanism isalso schematically shown for the n-type and p-typebehaviors of the tin oxide foam sensor in Fig. 7b.

Fig. 6. Cube of the mean strut size measured by varying the sin-tering time from 6 h to 24 h at a constant temperature of 1150�C.Scale bars 5 lm.

Nam, Kim, Choi, Park, Kang, Sung, Lee, and Choe

Gas sensitivity is commonly defined as the resis-tance ratio of Rg/Rair, where Rg and Rair stand forthe electrical resistances in the exposed gas and inair, respectively.57 We measured the gas sensitivityof our tin oxide foam to be 13.1 at an ethanol gasconcentration of 60 ppm at a temperature of 300�C,which is quite remarkable given that its behavior isa p-type semiconductor with no added catalyst. Asshown in Fig. 7a, the sensitivity (variation in resis-tance) of the tin oxide foam increases almostlinearly with increasing ethanol concentration.When we define the response time as the timerequired for the resistance to reach 90% of themaximum resistance value after the injection ofethanol gas and the recovery time as the timerequired to reach 90% of the initial resistance valueafter the ethanol gas is turned off, the measuredaverage response and recovery times for our tinoxide foam were 122 s and 113 s, respectively.

In Fig. 8, we compare the gas-sensing response (Ra/Rg) of our tin oxide foam with those from theliterature with similar operating conditions (ethanolgas concentrations from 20 ppm to 60 ppm at300�C).55,58,59 The tin oxide foam in the present studyexhibits a significantly higher response (13.1 at anethanol gas concentration of 60 ppm at a tempera-ture of 300�C) than the reported n-type SnO2

nanowires58 (3.5 under the similar testing condition).This is remarkable because our tin oxide foam has amicroscale pore structure and shows normally lessefficient p-type semiconductor behavior. Further-more, the sensing performance of the p-type tin oxidefoam sensor synthesized in this study is far superiorto that of the reported p-type NiO nanoparticles.55

For example, the response of the present tin oxidefoam sensor is 13.1 at an ethanol gas concentration of60 ppm at a temperature of 300�C, whereas that ofthe p-type NiO nanoparticles is �2 under the similarconditions as seen in Fig. 8. Indeed, the gas sensingperformance of 13.1 obtained from our p-type tinoxide foam sensor is even greater than that of the 1.0wt.% ruthenium-loaded NiO nanoparticles (�10 at

an ethanol gas concentration of 3000 ppm at300�C).55 The much higher response of our tin oxidefoam sensor is mainly attributed to its unique well-connected 3D struts with a combination of small andlarge pores, which results in a larger surface area. Onthe other hand, the addition of a Fe2O3-nanoparticlecoating on the n-type SnO2 nanowires58 achieved asignificantly improved response (better sensitivity),as shown in Fig. 8. This is because an effectivechemical doping of Sn4+ sites can stabilize the SnO2

surface and reduce its grain size to enhance both thecatalytic activity and sensor response as comparedwith pure SnO2.58,60 Iron (Fe3+; 0.80 A) doping intothe SnO2 matrix (Sn4+; 0.74 A) could improve theelectrocatalytic activity of the SnO2 nanowires.58,60

Figure 8 also compares the performance of the SnO2

nanorods created on n-type SnO2 nanoparticles,59

Fig. 7. (a) Variation in electrical resistance (gas sensitivity) of tin oxide foam sensor (heat-treated at 1150�C for 10 h), depending on themeasurement time during exposure to various ethanol gas concentrations, and (b) schematic diagram describing the general operating sensingmechanisms of the n- and p-type semiconductor sensors and the sensing mechanism in the tin oxide foam sensor used in this study.

Fig. 8. Comparison of sensing responses of the present tin oxidefoam with NiO nanoparticles,55 as-synthesized SnO2 nanowires,58

Fe2O3 nanoparticle-coated SnO2 nanowires,58 and SnO2 nanorodson nanoparticles59 as a function of the ethanol concentration at300�C. All four dotted lines are obtained from the Refs. 55,58, and 59for comparison with our tin oxide foam result.

Morphology and Gas-Sensing Properties of Tin Oxide Foams with Dual Pore Structure

which shows the best performance of the four tinoxide materials, including our p-type tin oxide foam.This excellent performance is most likelyattributable to the exceptionally large surface areacreated by the morphological combination of nanor-ods on nanoparticles.

We believe there to be plenty of room for perfor-mance improvement and that it is worthwhile toidentify performance enhancement strategies. First,we can enhance the gas-sensing properties of our tinoxide foam by decreasing its pore size and increas-ing the surface area by adjusting major processingparameters such as sintering condition and powdercontent. Second, we can also improve its gas-sensingperformance by adding a small amount of a chem-ical catalyst such as Pd. The fine-tuned pore struc-ture in the tin oxide foam can facilitate gas diffusionby shortening the reaction path, while the additionof a catalyst can reduce unnecessary side reactions.Finally, we can also alter the heat-treatmentprocess to completely transform a p-type SnO intoan n-type SnO2, which is known to exhibit bettergas-sensing performance.

SUMMARY AND CONCLUSIONS

We successfully synthesized tin oxide foam by afreeze-casting method. We added a small amount ofCuO powder (1.0 mol.%) to the SnO2 slurry toenhance the densification of the SnO2 powders. Weobserved morphological changes in the pore size,wall width, and porosity by varying the sinteringtemperature from 1050�C to 1250�C for a constantsintering time of 10 h. We then demonstrated thefeasibility of using freeze-cast tin oxide foam in gas-sensing applications by exposing it to ethanol gas atvaried concentrations from 20 ppm to 60 ppm. Thegas sensitivity of the freeze-cast tin oxide foam wasaround 13% at an ethanol gas concentration of60 ppm at a temperature of 300�C, which is remark-able considering that it shows p-type semiconductorbehavior with no catalyst added.

ACKNOWLEDGEMENTS

This work was supported by the NRF-2016-Fos-tering Core Leaders of the Future Basic ScienceProgram/Global Ph.D. Fellowship Program from theNational Research Foundation (NRF, 2016H1A2A1909161) of Korea. Choi and Choe also would like toacknowledge additional support from the NRF(2015R1D1A1A01060773, 2009-0093814, 2014R1A2A1A11052513).

REFERENCES

1. F. Schedin, A.K. Geim, S.V. Morozov, E.W. Hill, P. Blake,M.I. Katsnelson, and K.S. Novoselov, Nat. Mater. 6, 652(2007).

2. B.F. Bessac and S.-E. Jordt, Proc. Am. Thorac. Soc. 7, 269(2010).

3. C. Hou, J. Li, D. Huo, X. Luo, J. Dong, M. Yang, and X. Shi,Sens. Actuators B. 161, 244 (2012).

4. A. Hempel-Jorgensen, S.K. Kjaergaard, L. Molhave, andK.H. Hudnell, Arch. Environ. Health 54, 416 (1999).

5. M.L. Boeglin, D. Wessels, and D. Henshel, Environ. Res.100, 242 (2006).

6. G.F. Fine, L.M. Cavanagh, A. Afonja, and R. Binions,Sensors 10, 5469 (2010).

7. N. Yamazoe, Sens. Actuators B. 5, 7 (1991).8. B. Zhang, W. Fu, H. Li, X. Fu, Y. Wang, H. Bala, X. Wang,

G. Sung, J. Cao, and Z. Zhang, Appl. Surf. Sci. 363, 560(2016).

9. J.F. Boyle and K.A. Jones, J. Electron. Mater. 6, 717 (1977).10. J. Xu, Q. Pan, Y. Shun, and Z. Tian, Sens. Actuators B. 66,

277 (2000).11. T. Waitz, T. Wagner, T. Sauerwald, C.-D. Kohl, and M.

Tiemann, Adv. Funct. Mater. 19, 653 (2009).12. X.-L. Li, T.-J. Lou, X.-M. Sun, and Y.-D. Li, Inorg. Chem.

43, 5442 (2004).13. B. Karunagaran, P. Uthirakumar, S.J. Chung, S. Velu-

mani, and E.-K. Suh, Mater. Charact. 58, 680 (2007).14. I. Hotovy, V. Rehacek, P. Siciliano, S. Capone, and L.

Spiess, Thin Solid Films 418, 9 (2002).15. J. Zhang, J. Liu, Q. Peng, X. Wang, and Y. Li, Chem. Mater.

18, 867 (2006).16. H. NguyenandS.A.El-Safty,J.Phys.Chem.C115,8466 (2011).17. S. Pokhrel, C.E. Simion, V. Quemener, N. Barsan, and U.

Weimar, Sens. Actuators B . 133, 78 (2008).18. N. Barsan, C. Simion, T. Heine, S. Pokhrel, and U. Weimar,

J. Electroceram. 25, 11 (2010).19. H.-J. Kim and J.-H. Lee, Sens. Actuators B. 192, 607 (2014).20. K. Pirkanniemi and M. Sillanpaa, Chemosphere 48, 1047

(2002).21. M.M. Bettahar, G. Costentin, L. Savary, and J.C. Lavalley,

Appl. Catal. A. 145, 1 (1996).22. Y. Motooka and A. Ozaki, J. Catal. 5, 116 (1966).23. S.S. Kaye and J.R. Long, J. Am. Chem. Soc. 127, 6506 (2005).24. M. Kodu, T. Avarmaa, A. Floren, and R. Jaaniso, J. Eur.

Ceram. Soc. 33, 2335 (2013).25. J. Huang, L. Wang, C. Gu, and J.-J. Shim, Mater. Lett. 136,

371 (2014).26. N. Yamazoe, G. Sakai, and K. Shimanoe, Catal. Surv. Asia

7, 63 (2003).27. W.Y. Yin, B.Q. Wei, and C.W. Hu, Chem. Phys. Lett. 471,

11 (2009).28. T. Sahm, L. Madler, A. Gurlo, N. Barsan, S.E. Pratsinis,

and U. Weimar, Sens. Actuators B. 98, 148 (2004).29. L.C. Tien, S.J. Pearton, D.P. Norton, and F. Ren, Appl.

Phys. A 91, 29 (2008).30. G.X. Wang, J.S. Park, M.S. Park, and X.L. Gou, Sens.

Actuators B. 131, 313 (2008).31. D.J. Wales, J. Grand, V.P. Ting, R.D. Burke, K.J. Edler,

C.R. Bown, S. Mintova, and A.D. Burrows, Chem. Soc. Rev.44, 4290 (2015).

32. N.T.P. Nhung, P.V. Tong, C.M. Hung, N.V. Duy, N.V.Chien, N.V. Vinh, N.T. Tuyen, and N.D. Hoa, RSC Adv. 6,64215 (2016).

33. Y.H. Cho, X. Liang, Y.C. Kang, and J.-H. Lee, Sens. Actu-ators B. 207, 330 (2015).

34. Y.C. Goswami, V. Kumar, and P. Rajaram, Mater. Lett.128, 425 (2014).

35. P. Ivanov, E. Llobet, X. Vilanova, J. Brezmes, J. Hubalek,and X. Correig, Sens. Actuators, B 99, 201 (2004).

36. R. Zhang, T. Zhou, L. Wang, Z. Lou, J. Deng, and T. Zhang,New J. Chem. 40, 6796 (2016).

37. C.M. Hung, N.D. Hoa, N.V. Duy, N.V. Toan, D.T.T. Le, andN.V. Hieu, J. Sci. Adv. Mater. Devices 1, 45 (2016).

38. H. Jo, M.J. Kim, H. Choi, Y.-E. Sung, H. Choe, and D.C.Dunand, Metall. Mater. Trans. E 3, 46 (2016).

39. E.R. Leite, J.A. Cerri, E. Longo, J.A. Varela, and C.A.Paskocima, J. Eur. Ceram. Soc. 21, 669 (2001).

40. T. Kimura, S. Inada, and T. Yamaguchi, J. Mater. Sci. 24,220 (1989).

41. C.I. Drowley, 1-Application of Materials CharacterizationTechniques to Silicon Epitaxial Growth (Woburn: Butter-worth Heinemann, 1993).

Nam, Kim, Choi, Park, Kang, Sung, Lee, and Choe

42. M.K. Shobana, V. Rajendran, K. Jeyasubramanian, andN.S. Kumar, Mater. Lett. 61, 2616 (2007).

43. M.K. Shobana, W. Nam, and H. Choe, J. Nanosci. Nan-otechnol. 13, 3535 (2013).

44. B. Ravel and M. Newville, J. Synchrotron Radiat. 12, 537(2005).

45. X.-T. Chen, K.-X. Wang, Y.-B. Zhai, H.-J. Zhang, X.-Y. Wu,X. Wei, and J.-S. Chen, Dalton Trans. 43, 3137 (2014).

46. G. Leveque, S. Abanades, J.-C. Jumas, and J. Olivier-Fourcade, Ind. Eng. Chem. Res. 53, 5668 (2014).

47. H. Park, J.H. Um, H. Choi, W.-S. Yoon, Y.-E. Sung, and H.Choe, Appl. Surf. Sci. 399, 132 (2017).

48. K. Nagashima and Y. Furukawa, J. Cryst. Growth 171, 577(1997).

49. G.W. Young, S.H. Davis, and K. Brattkus, J. Cryst. Growth83, 560 (1987).

50. S. Deville, E. Saiz, and A.P. Tomsia, Acta Mater. 55, 1965(2007).

51. J.H. Um, M. Choi, H. Park, Y.-H. Cho, D.C. Dunand, H.Choe, and Y.-E. Sung, Sci. Rep. 6, 18626 (2016).

52. Y.-H. Cho, J. Wolfenstine, E. Rangasamy, H. Kim, H. Choe,and J. Sakamoto, J. Mater. Sci. 47, 5970 (2012).

53. S.I. Bae and S. Baik, J. Mater. Sci. 28, 4197 (1993).54. K. Mirabbaszadeh and M. Mehrabian, Phys. Scr. 85,

0357001 (2012).55. V. Kruefua, A. Wisitsoraatb, D. Phokharatkulb, A.

Tuantranontb, and S. Phanichphant, Sens. Actuators, B236, 466 (2016).

56. W. Guo, L. Fu, Y. Zhang, K. Zhang, L.Y. Liang, Z.M. Liu,H.T. Cao, and X.Q. Pan, Appl. Phys. Lett. 96, 042113(2010).

57. I. Jimenez, J. Arbiol, G. Dezanneau, A. Cornet, and J.R.Morante, Sens. Actuators B. 93, 475 (2003).

58. K.S. Choi, S. Park, and S.P. Chang, Sens. Actuators B. 238,871 (2017).

59. Y.J. Chen, L. Nie, X.Y. Xur, Y.G. Wang, and T.H. Wang,Appl. Phys. Lett. 88, 083105 (2006).

60. N. Lavanya, C. Sekar, S. Ficarra, E. Tellone, A. Bonav-ita, S.G. Leonrdi, and G. Neri, Mater. Sci. Eng., C 62, 53(2016).

Morphology and Gas-Sensing Properties of Tin Oxide Foams with Dual Pore Structure