Interfacial Doping of Heteroatom in Porous SnO2 for Highly Sensitive Surface Properties.

The design and synthesis of heteroatom-doping porous materials with unique surface/interfaces are of great significance for enhancing the sensitive surface performance in the fields of catalytic energy, especially gas sensor, CO oxidation, and ammonium perchlorate decomposition. Usually, the template method followed by a high-temperature calcination process is considered as the routes of choice in preparing ion-doped porous materials, but it requires extra templates and will undergo complicated steps. Here, we present a simple fusion/diffusion-controlled intermetallics-transformation method to synthesize various heteroatom-doping porous SnO2 only by changing the species of intermetallics. By this new method, Ni-doped popcornlike SnO2 with plenty of ∼30 nm pores and two kinds of Cu-doped SnO2 nanocages was successfully constructed. Phase-evolution investigations demonstrated that growth kinetics, diffusion, and solubility of the intermediates are highly related to the architecture of final products. Moreover, low-solid-solution limit of MO x (M: Ni, Cu) in SnO2 made the ion dope close to the surface to form a special surface/interfaces structure, and selective removal of MO x produce abundant pores to increase the surface area. As a consequence, Ni-doped composite exhibits higher sensitivity in formaldehyde detection with a relative low-operating temperature in a short response time (i.e., 23.7-50 ppm formaldehyde, 170 °C, and 5 s) and Cu-doped composites show excellent activity in decreasing the catalytic temperature of CO oxidation and ammonium perchlorate decomposition. The fusion/diffusion-controlled intermetallics-transformation method reported in this work could be readily adopted for the synthesis of other active heteroatom-doping porous materials for multipurpose uses.

Introduction

Modified surface/interfaces with more internal and outside active sites would provide an excellent platform for catalytic reactions because most of the catalytic reactions are found to be influenced by the structure of surface or interfaces.[1−7] Heteroatom doping is a promising route to decorate the microstructures including the surface/interfaces structure of materials because the previous study demonstrated that doping metal ions into semiconductors could regulate the Fermi level and activate the lattice oxygen. Using such a chemistry, one is able to create excellent catalytic materials for environmental and energy applications.[8] However, simple ion doping could not meet the higher requirement with the development of materials science and technology. More efficient and multifunction materials are needed. Thus, pore-creating becomes a good way to further modify the architecture of material and increase its surface activity sites to improve their gas adsorption and sensing ability, energy storage capacity, and many other intriguing merits.[9−12] Consequently, porous semiconductors with heteroatom doping would become excellent multifunctional materials if the special surface/interfaces structure and open-pore channels could be derived from the doping ions and architecture of pores.

Many strategies have been reported for pore-creation, which include soft/hard template methods. With these literature methods, varied porous structures such as molecular sieve (e.g., dual-mesoporous organosilica nanoparticles),[13] hollow (e.g., hollow Co3O4 nanocages, SnO2–Fe2O3–C hollow sphere),[14,15] and yolk–shell (e.g., Au@NiO, Fe3O4@SiO2) structures have been synthesized successfully.[16,17] Despite these progresses, the conventional inserted template strategy often needs extra templates and multistep processes. Afterward, the development of self-template strategy simplified the procedures and expanded the diversity of porous materials. CoSn(OH)6 hollow nanoboxes,[18] CoMn2O4 hollow microcubes, NiCo2O4 hollow sphere, and NiCo2S4 ball in ball-hollow sphere could be synthesized.[19−22] However, it still suffers difficulties in preparing porous materials with desired architectures and barely accessible to synthesize semiconductors merited with pores and heteroatom doping as well. It appears that searching for simple methodologies of preparing porous materials with heteroatom doping would enrich the multifunctional material system, which would further affect the scientific and technological applications.

Herein, we took porous SnO2 as a model semiconductor to study because it has a plenty of rutile-phase counterparts (e.g., TiO2, VO2) with many uses. Considering that, for ion doping, high-temperature calcination and/or high-pressure hydrothermal reactions are frequently used to import metal ions into the matrix in virtue of a strong driving force.[23] We designed a simple route that transform Sn-containing intermetallics to porous SnO2 semiconductors with a heteroatom doping through calcination and subsequent acid-etching process. The first step affords heteroatom doping in some regions through ion diffusion between solids, and the second step will produce more accessible pores in between the SnO2 particles. Such a strategy is based on our considerations: (i) the doping type of metal in the porous semiconductors relies on the choice of metal compositions in intermetallic compounds, (ii) the doping content relies on the solid solubility limit of the dopants in the final oxides, and (iii) the architecture of the final porous semiconductor is relative to the growth kinetics of intermediates MO (M: Ni, Cu, and Co). If the excess of transition metal oxides were selectively removed, one may expect to get heteroatom doping SnO2 semiconductors with pores around. With this strategy, Cu-doped nanocages, SnO2 derived from Sn–Cu, and Ni-doped popcornlike porous SnO2 from Sn–Ni intermetallics were synthesized successfully. These materials are strikingly featured by special surface/interfaces structure and open-pore channels, which show promising performance in catalytic field, pollutant detection, and control. This method reported here provides a picture that intermetallics could be excellent precursors to fabricate multicavity heteroatom doping semiconductors with diverse nanostructures for multipurpose uses.

Results and Discussion

Transformation of Sn-Containing Intermetallics to Diverse Heteroatom-Doping Porous SnO2

Here, we initiated a new route to prepare diverse heteroatom-doped porous SnO2 through a transformation from Sn-containing intermetallics, differing from the conventional template methods ever used which need extra templates and will undergo complicated procedures. The key to this route is the transformation reaction. To make clear the transformation process for diverse porous SnO2, we prepared four kinds of Sn-containing intermetallic compounds (Cu3Sn, Cu6Sn5, Sn–Ni, and Sn–Co) following the procedure reported in our previous work.[24] These intermetallic compounds were taken as precursors for the transformation, which underwent an oxidization reaction in O2 atmosphere and a subsequent acid etching by hydrothermal condition. Prior to the oxidization reaction, the optimum oxidation temperature of the precursors was initially determined by thermogravimetric analysis (TGA)–differential scanning calorimeter (DSC) data analysis. As indicated in Figure S2, the onset oxidation temperature is about 200–300 °C, whereas the terminated transformation temperature is about 800–900 °C for all of these intermetallic precursors. Thus, we selected 800 °C as the calcination temperature for the following transformations under the considerations that at this temperature, there would be complete oxidation. As indicated by XRD in Figure S3, all Sn-containing intermetallic precursors transformed to the corresponding mixed oxides completely after calcinations at 800 °C. Except for SCo-800, acid-etching of the resultant mixed oxides (SN-800, SC1-800, and SC2-800) produces a pure tetragonal-phase SnO2 with a space group of P42/mnm (JCPDS, no. 41-1445) (Figures g and S2). This transformation relies on the selective dissolution of the component oxides. For instance, component oxides such as NiO, CuO, and Co3O4 are relatively easier to be dissolved in an acid solution, whereas SnO2 is on reverse, resulting in the formation of porous SnO2-based products. Products originated from intermetallic Sn–Co did not form porous SnO2-based materials, which is probably because of the relative low solubility of Co3O4 in acid condition, consistent with the results previously reported.[25] From the photograph in Figures c and 2, one could clearly see colored acid aqueous solutions which suggest partial or complete removal of NiO, Co3O4, or CuO from their mixed oxidations. Representative scanning electron microscopy (SEM) and transmission electron microscopy (TEM) images demonstrate that three kinds of porous SnO2-based products were successfully synthesized (Figures and 2). Meanwhile, SN-800H derived from intermetallic Sn–Ni shows a popcornlike geometry, whereas both SC1-800H derived from intermetallics Cu6Sn5 and SC2-800H from Cu3Sn showed the well-shaped nanocage architectures. In spite of the residual Co3O4 in SCo-800H, the skeleton of polyhedron remained with Co3O4 aggregated in between particles (Figures , S4, and 3).

Figure 1

SEM and TEM images that monitor the transformation processes to the porous SnO2: (a) As-prepared precursor (Sn–Ni), (b) mixed oxide (SN-800) after calcination of the precursor, and (c–f) porous SnO2 (SN-800H) after acid-etching and hydrothermal treatment. The picture in the inset of (c) is the photograph of the dispersion after hydrothermal treatment using HNO3 aqueous solution. (g) Rietveld refinement results for X-ray diffraction (XRD) pattern of SN-800H. Symbol (×) in black represents the experimental diffraction data, and the red solid line denotes the calculated data. Green line at the bottom is the deviation between the experimental and calculated value. Heart symbol represents the internal standard peaks of Al. (h) N2 adsorption/desorption isotherm and (i) X-ray photoelectron spectroscopy (XPS) spectrum for SN-800H.

Figure 2

SEM images of the as-prepared precursors: (a1) Cu6Sn5, (b1) Cu3Sn, and (c1) Sn–Co; the corresponding mixed oxides: (a2) SC1-800, (b2) SC2-800, and (c2) SCo-800 after calcinations of the precursors; and the final porous SnO2: (a3) SC1-800H, (b3) SC2-800H, and (c3) SCo-800H. The picture in the far right is the photograph for the products after hydrothermal treatment using HNO3 aqueous solution.

Figure 3

(a,d) Dark-field TEM spectrum, EDS elemental mapping of (b) Sn and (c) Ni, and (e) EDS line in SN-800. Scale bar in (a) and (d) is 500 nm. (f,i) Dark-field TEM spectrum, EDS elemental mapping of (g) Sn and (h) Ni, and (j) EDS line in SN-800H. Scale bar in (f,i) is 200 nm.

Different from SC1-800H and SC2-800H, SN-800H inherits the skeleton of the precursor Sn–Ni, forming a multicavity popcornlike porous SnO2. TEM and N2 adsorption/desorption were carried out to investigate the inner architecture of SN-800H (Figure h). It is seen that popcornlike porous SnO2 is constituted by plenty of nanoparticles with a size in the range of 17–34 nm, leaving some ∼30 nm pores in between the particles. In contrast, the other two porous SnO2 derived from the precursor Sn–Cu showed nanocage microstructures with a large cavity in the middle and multigaps outside. Measurements of size and Brunauer–Emmett–Teller (BET) surface area of diverse porous SnO2 were conducted with an aim to describe the morphology features digitally. As indicated by XRD patterns in Figure a, different half-width at half-maximum for final porous SnO2 products suggested distinct sizes. Thus, we calculated the particle size from XRD data using Scherrer formula. Results showed that porous SnO2 derived from Sn–Cu intermetallics (Cu6Sn5 and Cu3Sn) showed a relatively larger size (∼79 nm for SC1-800H and ∼46 nm for SC2-800H) than SnO2 (∼25 nm) when fabricated by Sn–Ni, similar to the tendency of statistics data from SEM images. The smaller crystal size of SN-800H might be because of the confined growth of SnO2 caused by NiO–SnO2 interactions and doping-effect-induced growth inhibition. Compared to the variations in particle size, BET surface area showed a reversed trend, which is logical (Table S3). The BET surface area of SN-800H is calculated to be 10.4 m2/g, which is larger than that of SC2-800H (8.8 m2/g) and SC1-800H (3.6 m2/g).

Figure 4

(a) XRD patterns for the samples SC1-800H, SC2-800H, and SN-800H; (b) particle size and BET surface area of varied porous SnO2. Triangular symbol and pillar represent crystal size calculated from XRD using Scherrer formula and SEM statistics, respectively. (c) UV–vis spectra and (d) EPR for as-synthesized porous SnO2 after an excitation at 450 K.

From calcination to acid etching, multistep treatments under high-temperature and high-pressure hydrothermal conditions might import foreign ions (i.e., Ni, Cu, and Co) into SnO2 because of ion diffusion. To make sure of this possibility, EDS elemental mapping images were recorded to determine the chemical composition of porous SnO2. As shown in Figure S6, weak but uniform Cu signals are seen throughout the SnO2 nanoparticles, which signified a homogenous doping of Cu. Similarly, from EDS line and EDS mapping, one can see that Ni and Sn signals rise and fall in turns and tend to become uniform after acid etching, and Ni was well-distributed in SN-800H particles. Cu and Ni signals were captured in XPS spectra and quantified by EDS spectra (Figures i and S7, Table S2), which gives a common specification that the multistep methodology reported in this work would acquire metal-doped porous SnO2. Doping contents for these porous SnO2 detected by EDS were found within the scope of 1–2 at. %, typically a low-level doping ratio suitable for many catalytic reactions.[26,27]

To verify successful doping and grasp more information on the unit cell structure of diverse porous SnO2 prepared by the current method, lattice parameters were calculated using least-squares method from XRD patterns. Rietveld refinement results in Table showed that the cell volume V and lattice parameter (c or a) of samples SC1-800H and SC2-800H are obviously increased when compared to bulk SnO2. This observation indicates that Cu2+ ions might incorporate into the lattice site of SnO2 crystal and elevate the lattice constant because Cu2+ (rCu = 0.073 nm) is larger than Sn4+ (rSn = 0.069 nm). However, there is no obvious change for SN-800H, which might be because of the fact that the ion radius of Ni2+ (rNi = 0.069 nm) is equivalent to Sn4+. From the viewpoint of electronic structure, ionic doping would generate impurity levels and decrease the band gap energy. Hence, we investigated the variation in band gaps (Eg) through measuring UV–vis diffuse reflection spectra of the samples. Pure SnO2 is a direct-type semiconductor with a gap of Eg = 3.6 eV, and its diffuse reflection spectrum is represented by a single-peak absorbance under ultraviolet region only.[28] Out of what one expects, from Figure c, one can see that the current-doped porous SnO2 showed two peaks in the ultraviolet region and expanded the absorbance edge to visible light. It might be because of the doping of foreign ions. For catalytic application, surface species play important roles, especially oxygen vacancies and surface-absorbed oxygen. High-temperature calcination under O2 will reduce the oxygen vacancies, however, ionic doping and subsequent acid etching might generate defects. Electron paramagnetic resonance (EPR) is an effective means to detect single electron trapped by oxygen vacancies or surface oxygen with a g factor at around 2.[29] As shown in Figure d, a valley appeared at about g = 2, which is associated with a signal of O– or electron trapped by surface defects. All-porous SnO2 showed a signal around g = 2 under a thermal excitation at 450 K, and this signal became more obvious for SN-800H and SC1800H. O 1s core-level XPS spectra for all of these samples are shown in Figure S8, which exhibit three peaks: the first one at the lowest binding energy of about 530.2 eV is attributed to the lattice oxygen of Cu–O or Ni–O; the second peak in the middle at a binding energy of about 530.6 eV can be identified to the lattice oxygen of Sn–O; and the third peak located in the highest energy of about 532.5 eV is deemed as O species adsorbed in oxygen vacancy. The peak area of SN-800H at 532.5 eV is greater than that in other areas, indicating the highest oxygen vacancy concentration of SN-800H, consistent with the EPR results.

Table 1

Lattice Parameters Obtained through Data Refinements for Porous SnO2 Prepared Using Different Intermetallic Precursors

samples	a or b (Å)	c (Å)	V (Å3)	wRp	Rp
SC1-800H	4.7500	3.1952	72.092	0.1127	0.0834
SC2-800H	4.7440	3.1905	71.804	0.1231	0.0865
SN-800H	4.7385	3.1871	71.559	0.1313	0.0985
SnO2-reference	4.7374	3.1860	71.51

Phase/Shape Evolution and the Relevant Mechanism during Transformation

It should be pointed out that apart from the large hollow in SC1-800H and SC2-800H, SN-800H showed multicavity property with popcorn morphology, which may be because of the distinct thermodynamic behavior of CuO and NiO. In the former case, CuO exhibits a faster growth rate, higher oxidation, and diffusion ability when compared to those of NiO. As a result, there occurs a serious structure destruction of cavernous Sn–Cu after high-temperature calcination. On the contrary, the popcorn morphology of Sn–Ni and the polyhedral shape of Sn–Co were maintained after calcination. This phenomenon is consistent with the observation from Figure b2, where CuO showed large agglomerates and complete phase separation from SnO2, exactly opposite to that of the uniform NiO–SnO2 and Co3O4–SnO2. However, it seems difficult to attribute the complete phase separation of CuO from SnO2 in SC1-800 only to the rapid growth of CuO; after all, crushed fragments coming from fast expansion of CuO covered on SnO2 particles did not exist. From Figure b2, one can see that a phase movement might occur during the calcination process, and we guess that there is a fusion behavior of intermetallics. Thus, we checked the phase diagram of binary metal alloy to find the melting point, as listed in Table . We found that Sn–Cu series own a relatively low melting point around 500–600 °C, whereas Sn–Ni or Sn–Co possesses a melting point higher than 1000 °C. For veracity, we tested the melting point (Tm) under N2 and the complete oxidization temperature (To) under O2 atmosphere using TGA–DSC measurement (Figure ). There is a sharp endothermic peak around 700 °C without mass changes, which presents a phase transition, that is, fusion. The higher melting point in TGA-DSC results than phase diagram might be because of the insufficient heat under small crucible in the thermogravimetric measurement. As expected, no endothermic peak appeared in Sn–Ni below 1000 °C. Thus, we suspect that there is a relation between the melting point and final architecture. Samples with higher Tm restrained phase separation, whereas those with lower Tm promoted phase moving. When dealing precursors with calcination temperature (T): To < T < Tm, the skeleton of the precursor remained. Considering that To is not a fixed value because oxidation is a kinetic process under control not only by temperature (T) but also by time (t), we put an unclear variable value (T) on To. As shown in Figure , when selecting a calcination temperature between Tm and To (±T), inherited-shape metal-doped porous SnO2 would be obtained, and when the calcination temperature is lower than Tm, the final products show an unclear shape.

Table 2

Melting Points and Oxidation Temperatures for the Precursors

precursor	melting point (°C)	oxidation temperature range (°C)
Cu3Sn	∼500	∼200–∼790
Cu6Sn5	∼600	∼300–∼890
Sn–Ni	∼1050	∼180–∼780
Sn–Co	∼1150	∼150–∼900

Figure 5

TGA–DSC curves for (a) Cu3Sn and (b) Sn–Ni measured at given atmospheres.

Figure 6

Schematic diagram for the strategy of synthesizing heteroatom-doping porous SnO2 through Sn-containing intermetallics. Tm is the melting point and To is the complete oxidation temperature. Considering that To is not a determined value, we put a range (T) to To.

To investigate the phase and shape-evolution process along with the sintering temperature, we took Cu3Sn and Sn–Ni as examples. Low-temperature oxidation environment (400 °C) means a slow oxidation rate, and only the surface composition could be oxidized to MO–SnO2 (Figure a,e). After the acid-etching, apart from the nanosized SnO2 grain attached on the outside, the major skeleton and composition of precursors maintained (Figure b,f). The oxidation degree enhanced as the calcination temperature increased, and the precursors were almost oxidized to MO–SnO2 by treating at 600 °C for 2 h (Figure S9). Porous SnO2 was obtained after acid etching. The morphology of SN-600H is similar to SN-800H as well as the precursor Sn–Ni, indicating an inheritance of the initial shape. However, SC2-600H showed a broken morphology, which verified our inference. It is just similar to a hollow architecture, not the same as the nanocages of SC2-800H, which may be because of its relatively lower diffusion rate than 800 °C. The formation of a hollowlike shape can be explained by Kirkendall effect.[30] When elevating the sintering temperature to 800 °C, as described before, SN-800 inherited the shape well, whereas SC2-800 was separated into two portions. This viewpoint would provide a speculation that Sn-containing intermetallics, especially Sn–Ni, are excellent precursors to fabricate multicavity porous SnO2 with a desired nanostructure using heteroatom doping.

Figure 7

SEM images for Sn–Ni series on the left: (a) SN-400, (b) SN-400H, (c) SN-600, and (d) SN-600H and for Cu3Sn series on the right: (e) SC2-400, (f) SC2-400H, (g) SC2-600, and (h) SC2-600H. The scale bars are 200 nm for all images.

Multicatalytic Applications of the Heteroatom-Doping Porous SnO2 Derived from the Transformations

Gas Sensoring

Because heteroatom-doping porous SnO2 has unique composition and morphology that endow it specific active sites as well as gas diffusion channels, doping porous SnO2 may have excellent performance in gas sensoring. To obtain superior gas-sensing properties, one has to examine the static response–recovery sensing, as reported elsewhere.[11] Here, we investigated the statics response–recovery sensing of Cu-doped nanocages (SC1-800H, SC2-800H) and Ni-doped popcornlike SnO2 (SN-800H) toward 50 ppm formaldehyde in a temperature range from 170 to 300 °C. It should be mentioned that SC1-800H and SC2-800H devices have resistance higher than the detection line (500 MΩ) at a relatively low temperature, thus no response signal could be found, whereas SN-800H showed a response in the whole detection temperature range. Besides, from the temperature-dependent curves of sensitivity represented in Figure S10, one can see that SN-800H showed a superior performance to the other two materials with the maximum response of 23.7 at 170 °C, whereas SC1-800H and SC2-800H showed maximum responses of 12.8 at 200 °C and 6.6 at 280 °C, respectively. In particular, SN-800H reduces the best operation temperature of most SnO2 materials from >300[12,26,31] to 170 °C, much lower than that of SC1-800H (200 °C) and SC2-800H (280 °C). The higher response of SN-800H might be because of the following facts: (i) Ni doping induced subsize leads to a relatively large surface area, (ii) the surface-active O– verified by EPR, and (iii) doping element of Ni is active in the gas-sensing performance.[26,27,32,33]Figure a compares the response–recovery curve of porous SnO2 to different concentrations of formaldehyde (5, 10, 20, 50, 100, and 200 ppm). The response amplitude sensors were increased with increasing the concentration, indicating superior abilities to distinguish varied concentration of toxic gas. In accordance with temperature-dependent experiments, SN-800H sensor showed a higher response compared with SC1-800H and SC2-800H. The Ni-doped popcornlike porous SnO2 sensor has a sensitivity of 52.8 to 200 ppm formaldehyde, which is nearly 5 times higher than that of SC1-800H (10.5) and 6 times than that of SC2-800H (8.6). Besides, the response and recovery curves in Figure c also indicate a very short response time of 5 s and a recovery time of 3 s to 50 ppm formaldehyde, faster than most SnO2 materials reported previously.[12,34] The superior performance of SN-800H might be attributed to the unique popcorn shape and mesoporous structure, which enabled the response gas to go through the pore easily. The sensor sensitivity of SnO2 toward several other organic gases (50 ppm) is also investigated at their optimum operation temperature. As shown in Figure b, the responses of SN-800H sensor to formaldehyde are about 3, 8, and 17 times than that to acetone, toluene, and benzene, respectively. This indicates that SN-800H has a superior ability to distinguish formaldehyde from other VOCs.

Figure 8

(a) Response–recovery sensing curves of SN-800H; (b) sensor responses of given porous SnO2 to different formaldehyde concentrations (5–200 ppm) at their optimum operation temperature; and (c) sensor sensitivities of SN-800H, SC1-800H, and SC2-800H to various organic vapors (50 ppm) at their optimum operation temperature (170 °C for SN-800H, 200 °C for SC1-800H, and 280 °C for SC2-800H). Catalytic effects of the as-synthesized metal-doped porous SnO2 in (d) CO oxidation performance and (e) thermal decomposition of AP (ammonium perchlorate).

Catalytic Applications in CO Oxidation and AP Thermal Decomposition

Metal-doping porous materials could also be used in other catalysis such as CO oxidation and AP thermal decomposition benefiting from their increased active sites and special surface/interfaces. First, we tested the catalysis activity toward CO oxidation. As can be seen in Figure d, Cu-doped porous SnO2 showed an activity superior to Ni doping. SC2-600H showed the optimal performance with an onset conversion temperature of 180 °C and a total conversion temperature of 300 °C, lower than those for Ni-doped SnO2 and most SnO2-based materials.[6,35−37] Considering the higher surface area but lower CO oxidation activity of Ni-doped porous SnO2, the influence of surface area on CO catalysis could be excluded. In CO oxidation reactions process, SnO2 usually acts as supporters to active oxygen species, and the doping ions of Cu or Ni are active sites for CO adsorption and activation. Thus, the superior activity of Cu-doped porous SnO2 can be attributed to (i) the high activity of copper species, as in previously reported studies[38−40] and (ii) a strong interaction between doped copper ions and SnO2. Then, we used it in catalytic decomposition of AP, an oxidizing agent in propellants. Previous investigations showed that metal elements have played key roles in reducing the decomposition temperature.[40] Metal-doped porous SnO2 is expected to promote the catalytic reaction. TGA and DSC tests were employed to show the heat-release process in AP decomposition, as presented in Figure e. There are three peaks in the DSC curves. The first endothermic peak without mass loss at about 250 °C represents phase transformation from orthorhombic to cubic AP. The other two exothermic peaks are referred as the low-temperature decomposition and high-temperature decomposition process of AP, respectively. It showed that all of these metal-doped porous SnO2 possess the ability to reduce the decomposition temperature (Figure e), especially Cu-doped compounds with about 80 °C decrease. Lower surface area with higher performance indicates the existence of abundant activity species. Cu species showed superior activity in AP decomposition than Ni species in our previous study,[24] and the accelerated electron transfer of Cu species may be the major factor for enhanced performance. It appears that metal-doped porous SnO2 is a multifunction material for gas catalysis as well as solid catalysis.

Conclusions

We report on a new route to heteroatom-doping of porous SnO2: Cu-doped nanocages and Ni-doped popcornlike porous SnO2 through transformation of Sn-containing intermetallic compounds. The transformation processes involve calcination and a subsequent acid-etching process. These composites with low-content metal-ion doping (1–2% atomic ratio of metal to Sn) have diverse architectures and are featured by lattice as well as surface-state regulation. Such features enable metal-doped porous SnO2 to possess open-pore channels with more in/out surface and catalytic active species. Investigation on phase/shape-evolution process suggested that the fusion behavior of SnM, growth kinetics, and diffusion ability of MO have crucial effects on the final architecture. It is demonstrated that Sn-containing intermetallics could act as excellent precursors to fabricate heteroatom-doping porous SnO2 with desired nanostructures, necessary for multifunction performance such as gas sensing, CO oxidation, and AP thermal decomposition. Because SnO2 is a prototype semiconductor with many counterparts, we expected that these findings achieved here will put forward one’s ability of synthesizing more ion-doping porous functional materials.

Experimental Section

Sample Syntheses

All chemicals were used as received without further purification. The sample syntheses undergo two-step processes: (1) intermetallic precursor formation and (2) transformation of the precursor to metal-doped porous SnO2. First, Sn-containing intermetallics such as Cu6Sn5, Cu3Sn, Sn–Ni and Sn–Co were taken as the precursors. These precursors were prepared by a hydrothermal method, just following the procedure we recently reported.[24] Second, the precursors of Cu6Sn5, Cu3Sn, Sn–Ni, and Sn–Co were transformed into a series of oxides named as SC1-800, SC2-800, SN-800, and SCo-800, respectively, through an oxidization reaction under O2 atmosphere at 800 °C for 2 h. Afterward, 30 mL of 0.3 M HNO3 aqueous solution was added to form a mixed solution, which was transferred to an autoclave with a Teflon liner and heated at 140 °C for 2 h. The sample after hydrothermal crystallization was centrifuged and washed with abundant distilled water. Then, variations of porous SnO2 (named as SC1-800H, SC2-800H, SN-800H, and SCo-800H) were obtained after drying at 70 °C for precursors of Cu6Sn5, Cu3Sn, Sn–Ni, and Sn–Co, respectively. Similar nomenclature was used for other samples obtained at different calcination temperatures, as presented in Table S1.

Sample Characterization

Sample structures were identified by powder X-ray diffraction on a Rigaku D/Max 2550 diffractometer with a graphite monochromator, operating at 50 kV and 200 mA at room temperature, and using Cu Kα radiation (λ = 1.5418 Å). Commercial Al powder serves as an internal standard for peak positions determination mixed with samples evenly before the measurement to exclude artificial factors. Morphologies of the samples were further examined by field-emission SEM (Hitachi SU8020 electron microscope) and TEM (Philips/FEI Tecnai G2S TWIN microscope, acceleration voltage of 200 kV). Samples for TEM were prepared by dispersing a powder in ethanol and leaving a droplet of the suspension on a copper (double-sided for magnetic sample) or molybdenum microscope grid covered with perforated carbon. The selection of the microscope grid depends on the elements in sample to eliminate the outside interference in EDS and mapping measurement. Chemical compositions and valence states of the samples were detected by XPS, a Thermo ESCALAB 250Xi electron energy spectrometer using Al-Kα (1486.6 eV) as the X-ray excitation source. The Brunauer–Emmett–Teller (BET) surface areas of the samples were obtained by N2 adsorption/desorption isotherm performed in an apparatus of Micromeritics ASAP 2020. The UV–vis diffuse reflectance spectra of the samples were recorded on a PerkinElmer Lambda 20 UV–vis spectrometer.

The thermal behaviors of the samples were performed on TGA and DSC in O2 or N2 atmosphere over a temperature range of 30–1000 °C (NETZSCH STA 449F3 thermal analyzer). The use of atmosphere depends on the purposes. EPR measurements were performed using an A300-10/12 X-band spectrometer (Bruker) operating at 9.84 GHz.

Catalytic Activity Test

Gas-Sensing Test

Gas sensor is fabricated by coating viscous slurry of the sample onto a ceramic tube (diameter 1 mm and length 4 mm) positioned with a pair of Au electrodes and four Pt wires on both ends of the tube. A Ni–Cr alloy coil in the tube was employed as a heater to control the operation temperature. Gas-sensing tests were performed on a commercial CGS-8 Gas Sensing Measurement System (Beijing Elite Tech Company Limited). Environmental air is used as both the reference gas and the diluting gas. After the calculated target gas was injected into the test chamber for about 30 min by a microsyringe, the sensor was put into the test chamber. After the response completed, the sensor was removed to fresh air for recovery. The sensor working temperature was adjusted by varying the electric current. The sensor response is defined as S (the ratio of Ra to Rg), where Ra and Rg are the electrical resistance of the sensor in atmospheric air and in target gas, respectively.

CO Oxidation Test

Catalytic activity of the as-prepared catalysts toward CO oxidation was evaluated in a fixed-bed quartz microreactor using 50 mg of the catalyst. The feed gas was composed of 1% CO and 20% O2 balanced by He at a flow rate of 35 mL min–1 (roughly 80 000 h–1 GHSV). The reaction temperature was monitored by a thermal couple in the middle of the quartz microreactor. The effluent gas was detected by a gas chromatograph equipped with a TC-detector. No products other than those resulted from CO (i.e., CO2) were observed under the applied reaction conditions. The conversion of CO in the oxidation process was calculated from different concentrations between inlet or outlet gases of CO.

AP Thermal Decomposition Test

The catalytic roles of the samples in the thermal decomposition of AP were detected by TGA and DSC in N2 atmosphere over a temperature range of 50–600 °C (NETZSCH STA 449F3 thermal analyzer).