Structure and thermal stability of nanocrystalline Ce1−xRhxO2−y in reducing and oxidizing atmosphere

Materials Research Bulletin 48 (2013) 852–862

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Structure and thermal stability of nanocrystalline Ce1xRhxO2y in reducing and oxidizing atmosphere M. Kurnatowska, L. Kepinski * Institute of Low Temperature and Structure Research, Polish Academy of Sciences, Okolna 2, 52-422 Wroclaw, Poland

A R T I C L E I N F O

A B S T R A C T

Article history: Received 14 August 2012 Received in revised form 13 November 2012 Accepted 17 November 2012 Available online 24 November 2012

Nanocrystalline (4–5 nm) Ce1xRhxO2y mixed oxides (x = 0.03–0.21) were synthesized using water in oil microemulsion method. Morphology, microstructure and phase composition of the samples subjected to heat treatment in oxidizing and reducing atmosphere were investigated by TEM, SEM-EDS, XRD, Raman and IR spectroscopy. Samples with x  0.16 were structurally stable in oxidizing atmosphere up to 850 8C. Above this temperature the samples decomposed into Rh deficient, nanosized Ce1xRhxO2y and large (few mm) Rh2O3 crystals. In hydrogen atmosphere segregation of small (1–2 nm) Rh crystallites in special epitaxial orientation Rh (1 1 1)jjCeO2 (1 1 1) started at 500 8C. The epitaxial orientation of small Rh crystallites was preserved up to 1000 8C indicating their strong bonding to the ceria surface. Partial substitution of Rh for Ce strongly enhanced reducibility of ceria at low temperatures (below 200 8C) and hindered the sintering of ceria at high temperatures. Ce0.89Rh0.11O2y shows an interesting property of reversible extraction–dissolution of Rh upon reduction–oxidation cycles at 500 8C which is important for potential catalytic applications. ß 2012 Elsevier Ltd. All rights reserved.

Keywords: A. Oxides A. Nanostructures B. Chemical synthesis C. Electron microscopy C. X-ray diffraction D. Microstructure D. Catalytic properties

1. Introduction Nanocrystalline cerium oxide has attracted a wide interest due to the unique redox properties which can be improved by its doping with noble metals (NM). Rhodium is a noble metal catalyzing many catalytic reactions [1,2]. Thanks to a synergy effect Rh loaded on CeO2 is a very active catalyst for: TWC [3] and reactions important for fuel cell industry such as reforming of elcohols for hydrogen production [4,5] and preferential CO oxidation in the presence f hydrogen (PROX) [6,7]. Addition of CeO2 to Rh/Al2O3 or Rh/SiO2 catalysts improves their activity [3,5,8]. It has been shown that the presence of mixed sites consisting of an oxidized Rhx+ species stabilized with two oxygen vacancies (and associated to Ce3+ cations) of the partly reduced support accounts for high catalytic activity of ceria containing Rh catalysts [2,3,5,8,9]. Creation of the Ce–O–NM surface species enhances also the activity of other noble metals supported on ceria [2,10]. One of the ways of producing such active sites is dispersion of NM ions in ceria lattice. It has been shown that Ce1xNMxO2y (x = 0.01–0.02) mixed oxides exhibit better activity and stability than impregnated NM/CeO2 catalysts [11]. Till now Ce1xRhxO2y mixed oxide was obtained by solution combustion method, which

* Corresponding author at: Institute of Low Temperature and Structure Research, Polish Academy of Sciences, P.O. Box, 1410, 50-950 Wroclaw, Poland. Tel.: +48 71 343 50 21; fax: +48 71 344 10 29. E-mail address: [email protected] (L. Kepinski). 0025-5408/$ – see front matter ß 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.materresbull.2012.11.076

has important drawbacks. Despite of using low doping levels (x = 0.005–0.02) the samples produced were inhomogeneous: Rh enrichment at the surface or even phase separation (for x = 0.02) was observed [12]. Moreover, mean size of the obtained crystallites was rather large (around 50 nm). Similar problems were reported for other Ce1xNMxO2y mixed oxides (NM = Pd and Pt) obtained with this method [13,14]. In this work we investigated for the first time structure and thermal stability of Ce1xRhxO2y mixed oxide over a wide range of compositions x = 0–0.21. The oxides were obtained using microemulsion method, which, as it has been reported recently for Ce1xPdxO2y, allows to obtain small, homogeneous nanocrystals with narrow size distribution [15]. The results of this study on structure evolution during heating in reducing and oxidizing atmosphere are important because of potential applications of Ce1xRhxO2y as active catalysts. 2. Experimental Nanosized Ce1xRhxO2y mixed oxides were synthesized by water in oil microemulsion method described in detail in Refs. [15,16]. Triton X-100 was used as a non-ionic surfactant and cyclohexane and 1-pentanol as an organic phase. Water phase containing a solution of cerium and rhodium nitrates was poured into organic phase and stirred until transparent emulsion was obtained. Then the second reactant tetramethylammonium hydroxide (25% water solution) was added and the microemulsion

M. Kurnatowska, L. Kepinski / Materials Research Bulletin 48 (2013) 852–862

was stirred for 30 min. The organic phase was removed by decanting and the precipitated oxide was washed with acetone (two times) and then with methanol (four times). Finally the solid was separated by centrifugation. All powder samples were dried and then heated in oxygen at 500 8C for 2 h. Morphology, microstructure and composition were investigated by TEM (Philips CM20 SuperTwin at 200 kV), SEM-EDS (EDAX Pegasus XM4 spectrometer installed on FEI NovaNanoSEM 230) and XRD (Powder Diffractometer PANanalytical using Cu Ka radiation in the 2u range 10–1108). Analysis of HRTEM images was made with Digital Micrograph program (Gatan) and XRD patterns were refined using X’Pert HighScore Plus program by Rietveld method. The Raman spectra were acquired with Bruker FT Raman spectrometer with resolution 2 cm1 using Nd:YAG laser with emission line at 1064 nm. Infrared spectra were measured with Bio-Rad 575 C FT-IR spectrometer in KBr pellets for region 4000– 400 cm1 with a spectral resolution of 2 cm1. H2-TPR (temperature programmed reduction) was performed by heating the sample up to 900 8C in H2 (5 vol.%)/Ar flow (30 ml/min) with the heating rate of 108/min. The hydrogen consumption was monitored by thermo-conductivity (TCD) detector. Nanocrystalline Rh2O3 was synthesized by a hydrothermal method as follows. To 0.5 M solution of Rh nitrate a 25% NH4OH solution was added with vigorously mixing. Then the mixture was poured into a Teflon container and placed into a stainless steel autoclave equipped with a microwave heating. Reaction time and temperature was set on 15 min and 195 8C, respectively. The pressure during the synthesis raised to 40 bar. The sample was thoroughly washed with water, dried at 110 8C and calcined at 500 8C. As prepared sample was amorphous but after additional calcination at 650 8C nanocrystalline a-Rh2O3 was obtained in accordance with [17]. 3. Results 3.1. Fresh samples Composition of the as prepared Ce1xRhxO2y samples was confirmed by energy dispersive spectroscopy (EDS). The measured Rh contents (x = 0.06, 0.11, 0.16 and 0.21) were close to nominal ones (0.05, 0.1, 0.15 and 0.2). Multi-point EDS analysis was also performed to check the homogeneity of the samples. Deviation of composition was around 0.02 for all samples and no noticeable macroscopic phase separation was observed. XRD patterns of all the ‘‘as prepared’’ samples are shown in Fig. S1 (supporting information). All reflections belong to a fluorite structure of CeO2 with no evidence of any Rh2O3 reflections up to x = 0.21. Parameters characterizing structure of the samples (lattice parameter, mean crystallite size, micro strain and oxygen occupancy) were calculated using the Rietveld refinement method and are shown in Fig. 1. Mean crystallite size was also calculated from TEM images and the results are consistent with the XRD data

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(Fig. 1a). Observed decrease of the mean crystallite size of ceria with increasing amount of doping was reported before for similar systems [15,18]. The lattice parameter of Ce1xRhxO2y initially slightly decreases and then increases with increasing x (inset to Fig. 1a). The initial small but noticeable decrease of the lattice parameter may be connected with the fact that the size of Rh3+ ion (66.5 pm) is much smaller than that of Ce4+ (97 pm). At higher Rh contents this tendency is overwhelmed by the more profound effect of the lattice expansion of ceria due to decreasing crystallite size. This effect has been discussed in literature and in our previous paper [15]. In accordance with our data Gayen et al. [12] reported that the lattice parameter of Ce0.99Rh0.01O2y prepared by solution combustion method (0.54109 nm) was smaller than that of pure CeO2 (0.54113 nm). It should be mentioned however, that the mean size of ceria crystallites obtained using solution combustion method (50 nm) was much larger so that the effect of the size on the lattice parameter was much weaker. Morphology and size distribution of the crystallites was investigated by TEM. Up to x = 0.16 small, well define crystallites with sharp edges were observed (Fig. 2a) while for x = 0.21 an amorphous phase was also present as thin coating and small round forms at the surface of crystallites (Fig. 2b). This amorphous phase may consist of an excess Rh oxide that has not been incorporated into the ceria lattice. Phase composition of the samples was also checked by a vibration spectroscopy. Raman spectroscopy was not useful because of very strong heating of the dark samples, but infrared spectroscopy has proved to be helpful (Fig. 3). The spectra of the Rh doped samples are generally similar to those of pure ceria and are dominated by strong absorption below 750 cm1 due to the CeO2 stretching vibration d Ce–O–Ce [19]. With increasing Rh content the absorption edge shifted to higher wavenumbers, the band around 850 cm1 moved to lower wavenumebers and the band at 720 cm1 slowly disappeared. These observations indicate changes in the fluorite structure of cerium oxide caused by substitution of cerium by rhodium in the lattice (creation of oxygen vacancies and change ot the bond lengths). For x = 0.21 an additional broad band occurs at around 650 cm1, which corresponds to the strongest band of a-Rh2O3. This result corresponds to TEM and XRD data indicating that the maximum amount of Rh which may be introduced into the ceria lattice is around x = 0.16. For Pd doped ceria prepared using the same procedure slightly higher doping limit x = 0.21 was reported [15]. 3.2. Heating in oxidizing atmosphere Fig. 4 presents XRD patterns of three samples (x = 0.06, 0.11, 0.16) heated in oxygen at two temperatures 950 8C and 1000 8C. In Table 1 lattice parameters, mean crystallite size and phase composition of the samples are shown. It is seen that the presence of rhodium in the ceria lattice strongly inhibits the growth of crystallites at elevated temperatures (Table 1). After heating at

Fig. 1. Variation of the microstructure parameters of Ce1xRhxO2y with Rh content: mean crystallite size and lattice parameter (inset) (left), micro strain and oxygen occupancy (right).

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Fig. 2. HRTEM images and crystallite size distributions (insets) of Ce1xRhxO2y heated in O2 at 500 8C for 2 h: (a) x = 0.11 and (b) x = 0.21. In Fig 2b amorphous forms of Rh2O3 are marked.

800 8C in oxygen mean crystallite size of CeO2 was 56 nm while that of Ce0.89Rh0.11O2y only 8.8 nm. We observed similar, but less pronounced effect also for Ce0.89Pd0.11O2y (mean crystallite size 14 nm) [15] and for Ce0.9Ln0.1O2y (Ln–Yb, Lu) (mean crystallite size 21 nm) [20]. These findings may be explained by results of DFT calculations [21], which showed that the lattice distortions induced by doping (expressed as deviation from ideal 8 coordination of cations in CeO2) are high for transition metals (Ni, Mn, Cu, Ru), less for noble metals (Pt, Pd) and the smallest for rare earth metals (La, Y). First evidence of a phase separation (segregation of a-Rh2O3) occurred in the Ce0.84Rh0.16O2y sample heated at 950 8C,

Fig. 3. IR spectra of Ce1xRhxO2y, CeO2 and a-Rh2O3.

but after treatment at 1000 8C also for other samples (Fig. 4). The presence of a-Rh2O3 in the samples after heating in oxygen at 1000 8C was confirmed by Raman (Fig. 5a) and IR spectroscopy (Fig. 5b). It is interesting that despite relatively small amount of Rh2O3 segregated in the Ce0.89Rh0.11O2y (4.5 wt.% according to XRD – Table 1) intensities of the Raman bands are higher than those of ceria. It implies that Rh2O3 phase consists of crystallites bigger and better ordered than those of the mixed oxide. In fact this agrees with XRD patterns, showing small, but very sharp reflections of a-Rh2O3 (Fig. 4). Analysis of TEM micrographs revealed that heating in oxygen caused growth of the mean size of the Ce1xRhxO2y crystallites with simultaneous strong broadening of the crystallite size distribution. Unexpectedly however, no Rh2O3 phase could be detected by high resolution imaging and electron diffraction in the samples heated at 1000 8C. This result clearly disagrees with both XRD and Raman data. SEM observations clarified the mystery revealing the presence of scarce, large and well developed crystals different from ordinary doped ceria (Fig. 6). The EDS confirmed that the crystals had Rh2O3 composition. Such big crystals could be easily overlooked by TEM due to preparation method used (making suspension in methanol). Numerous point EDS analyses of the areas of the sample free from large Rh2O3 crystallites always revealed the presence of Rh in concentration corresponding to 3 wt.% Rh2O3. Analysis of the results presented above rises a few issues. The first concerns the amount of Rh2O3 determined from XRD patterns of the samples after oxidation at 1000 8C which is significantly less than that expected is the whole amount of Rh would be extracted from the lattice and transformed into crystalline Rh2O3 (Table 1).

Table 1 XRD characterization of the structure of Ce1xRhxO2y heated in oxygen.

0.06 0.11 0.16

950oC 1000oC

x 0.06 0.11 0.16

Fig. 4. XRD patterns of Ce1xRhxO2y (x = 0.06, 0.11 and 0.16) heated in oxygen at 950 8C and 1000 8C. Inset presents the change of the mean crystallite size for Ce0.89Rh0.11O2y and CeO2.

x

Temperature (8C)

a (nm)

Mean size (nm)

Rh oxide (wt.%)

0 0.11 0 0.11 0.06 0.11 0.16 0.06 0.11 0.16

500 500 800 800 950 950 950 1000 1000 1000

0.5411 0.5412 0.5410 0.5409 0.5411 0.5412 0.5412 0.5411 0.5412 0.5410

8.6 4.2 56 8.8 20.5 14 13.5 31.5 37 33

– – – – – – Rh2O3 Rh2O3 Rh2O3 Rh2O3

2.7% traces 4.5% 6.3%

Note: Total amount of Rh in the samples corresponds to 4.5, 8.4 and 12.3 wt.% Rh2O3 for x = 0.06, 0.11 and 0.16, respectively.

M. Kurnatowska, L. Kepinski / Materials Research Bulletin 48 (2013) 852–862

a

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b

Fig. 5. (a) Raman spectra of Ce1xRhxO2y (x = 0.06, 0.11) heated at 1000 8C in O2. (b) IR spectra of Ce1xRhxO2y (x = 0.03, 0.06, 0.11 and x = 0.16) heated in O2 at 950 8C and 1000 8C.

This result combined with EDS data presented above imply that significant part of Rh remained in the ceria lattice or is dispersed as an amorphous phase at the surface. Since no such phase was found in HRTEM images we assume that solubility of Rh in ceria lattice is quite high even at 1000 8C. Another confirmation of the presence of Rh in the ceria lattice after high temperature oxidation was obtained by observation of the changes induced in the sample during its reduction in hydrogen. Fig. S2 shows HRTEM image of the sample reduced at 500 8C for 1 h, where very small Rh particles distributed at the surface of ceria crystallites are seen. It will be shown in the next section that such images are obtained when as prepared Ce1xRhxO2y samples are reduced in hydrogen at or above 500 8C. Persistence of a depleted Ce1xRhxO2y even after oxidation at 1000 8C observed in this work is consistent with results of Chen et al. [9] who reported formation of Rh–O–Ce surface phase in Rh/Ce–Zr–O oxidized at 1000 8C for 5 h. In line with our data this surface phase produced no visible contrast in HRTEM images or additional diffraction features. Another issue connected with microstructure of the samples heated in oxidizing atmosphere is why after oxidation at 1000 8C we observe a-Rh2O3, which is in fact a low temperature polymorph, and what mechanism of the phase separation can account for the growth of such big, well-developed a-Rh2O3 crystals. According to literature a-Rh2O3 transforms irreversibly into high temperature b-Rh2O3 at temperatures of around 850– 950 8C [22] so the latter phase should be observed in our case. To address the first question we checked the kinetics of a-Rh2O3 to bRh2O3 transformation. Fig. S3 shows XRD diffractograms of pure, nanocrystalline a-Rh2O3 heated in oxygen at 1000 8C for 2 and 17 h. Refinement of the patterns revealed the presence of both aRh2O3 and b-Rh2O3 with percentage composition 70/30 and 30/

70 wt.%, respectively. This experiment shows that the a–b-Rh2O3 phase transition is slow and in the Ce1xRhxO2y samples treated at 1000 8C for 2 h a-Rh2O3 exists because of too short heating time. An interesting point is that sintering of bare a-Rh2O3 during heating in oxygen is slow and mean size of the crystallites was 13, 30 and 53 nm after heating at 650 8C for 2 h, 1000 8C for 2 h and 1000 8C for 17 h, respectively. Clearly, the crystal size of the aRh2O3 is much smaller than in Ce1xRhxO2y samples treated at similar conditions. Another experiment was performed to find out a possible role of ceria in the sintering and in phase transformations in Rh2O3. The nanocrystalline a-Rh2O3 (annealed at 650 8C) was impregnated with an aqueous solution of cerium nitrate, dried and calcined at 400 8C. Then the sample was heated in oxygen for 15 h at 1000 8C. XRD revealed the presence of crystalline CeO2 and Rh2O3 in both a and b forms. Interestingly, the crystallite size of Rh2O3 was still in nanometre size indicating that the presence of ceria is not enough to induce the abnormal crystal growth of Rh2O3 observed for Ce1xRhxO2y. The growth of large Rh2O3 crystals with well developed faces in the Ce1xRhxO2y indicates that some kind of chemical transport may be involved. One possibility could be a formation of a mobile (liquid) cerium–rhodium oxide phase at the surface of the Ce1xRhxO2y and its transport to the growing aRh2O3 crystals. A unique atomic mixing of Ce and Rh in the Ce1xRhxO2y seems to be necessary to form the mobile phase because no abnormal growth was observed in Rh2O3 impregnated with ceria. Mechanism of the phase separation in oxidizing atmosphere was also investigated by ‘‘in situ’’ XRD. The sample with x = 0.16 was heated with a rate 58/min to a desired temperature and then XRD pattern was acquired during 30 min (Fig. 7). Up to 850 8C only reflections corresponding to the fluorite structure were observed.

Fig. 6. SEM images of Ce1xRhxO2y (a) x = 0.06 (b) x = 0.11 heated in O2 at 1000 8C. Examples of EDS spectra of big crystals of Rh2O3 and Ce1xRhxO2y are shown as insets.

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Fig. 7. XRD patterns for Ce0.89Rh0.16O2y recorded ‘‘in situ’’ during heating in air.

The only effect of increasing temperature being narrowing and small shift of the reflections. Very small and broad new peaks occurred at 875 8C at around 33.5–358 and then at 950 8C characteristic sharp reflections of a-Rh2O3 appeared (Fig. 7). Upon further heating to 1000 8C some growth of intensity of the a-Rh2O3 reflections was observed. No noticeable changes of the XRD patterns of the sample were observed upon cooling the sample (Fig. S4). In the second ‘‘in situ’’ experiment the kinetics of the structure change of the same sample at elevated temperatures was studied (Fig. 8). When temperature of the sample achieved 800 8C the first XRD pattern was recorded during 30 min in 30–408 2Q range, where the strongest reflections of Rh oxides occur. Next, the temperature was kept constant for another 30 min and second pattern was recorded. The same procedure was repeated at higher temperatures up to 1000 8C. No reflections except those of fluorite type structure were observed till 875 8C, when a small but sharp feature occurred at 358 in the second XRD pattern (Fig. 8a). The position of the reflection coincides with the strongest (1 2 0) reflection of a-Rh2O3, pointing out that the decomposition of the Ce0.84Rh0.16O2y begins in air at 875 8C. The intensity of the reflection achieved the maximum at 950 8C but simultaneously broad feature at 34.28 appeared. The position of this feature corresponds to the strongest (1 1 4) reflection of b-Rh2O3 and its intensity increases with increasing temperature at the expense of reflection at 358 (a-Rh2O3). This data confirms that above 900 8C aRh2O3 transforms slowly to b-Rh2O3. Careful analysis revealed also the presence of very weak additional peaks at around 35.4–35.78 in the patterns of the sample heated at or above 925 8C (Fig. 8b). In appears that intensity of these features rather weakly depends on the temperature. The position of the reflections roughly corresponds to those of Ce0.55Rh2O4 [23] or RhO2 [24], but the fit is not very good (Fig. 8b). It is worth to stress that behavior of Rh/CeO2 in oxidizing atmosphere differs substantially from that of Rh supported on other oxides (La2O3, Y2O3, MgO, Ta2O5), where RhMOx compounds

a

form at high temperatures 800 8C preventing Rh sintering [25,26]. Among these compounds there are perovskites MRhO3, which easily form for Ln3+ lanthanides [27,28] but only under special, reducing conditions for Ce4+ [29]. Formation of such compounds hinders the reduction of Rh in the Rh/Ln2O3 catalysts (presence of the reduction maxima at 500–600 8C in H2-TPR profiles) [25]. Such effect does not appear for Rh/SiO2 [25], but has been reported for Rh/Al2O3, where calcination at 700 8C caused formation of hardly reducible Rh(AlO2)y [26]. Recently, successful synthesis of CeRh2O5 and Ce0.55Rh2O4 compounds containing Ce4+ ions by the flux method (PbO, V2O5, NaCl) at 1000–1100 8C was reported [23]. The authors were unable however, to obtain the same products by simple reaction of the oxides, which may indicate that small amount of additives may be necessary to stabilize the structures. One of these complex oxides Ce0.55Rh2O4 could possibly be observed in our samples (Fig. 8b). There is a discussion in the literature on the presence and stability of the Rh–O–Ce phase at the surface of CeO2 in Rh/CeO2 catalysts oxidized at high temperature. Li et al. [30] showed by TEM and supporting DFT calculations that a monolayer of Rh deposited on CeO2 (1 1 1) surface transforms during oxidation at 600 8C into a stable O-Rh-O sandwiched structure, which inhibits Rh agglomeration. The authors stated also, that contrary to PdO/ CeO2, the Rh–O–Ce surface layer is not reduced during TEM observation. The authors did not study the behavior of the system at higher oxidation temperatures, but showed that 20 min treatment in 1% H2/N2 at 200 8C reduced the surface Rh–O–Ce phase. In agreement with our data (see below) small Rh particles formed upon reduction were strongly bound with CeO2 (mean crystallite size after 5 h treatment at 700 8C in nitrogen was 2.2 nm). Interestingly, for Pd monolayer on CeO2 (1 1 1) similar treatment caused severe sintering of the metal phase (mean crystallite size was 11 nm). Chen et al. [9] reported formation of Rh–O–Ce surface phase in Rh/Ce–Zr–O oxidized at 1000 8C for 5 h, but could not observe any distinct surface structure using sophisticated TEM techniques. As similar conclusion was drawn by other groups [30,31] it is suggested that Rh–O–Ce layer at the surface does not form a crystal structure distinct from that already present in the sample and consequently does not result in diffraction or mass contrast. An example of such structure could be a solid solution where Rh3+ is substituted for Ce4+ in the surface layer of CeO2 to form Rh–O–Ce. The surface phase appears to be stable enough to withstand severe conditions during HRTEM observation, because no beam induced reduction was observed [9]. Bernal et al. [32] reported that oxidation of Rh particles on CeO2 at 900 8C caused spreading of the Rh oxide phase over the surface creating an amorphous layer a few angstroms thick. Structure of this phase could not be however, observed by TEM because contrary to [9,30] it underwent fast reduction under the electron beam into small, crystalline Rh particles. The authors suggest that

b

Fig. 8. (a) Kinetic ‘‘in situ’’ XRD study of Ce0.84Rh0.16O2y in air. First measurements (10 ) were performed during 30 min immediately after reaching the desired temperature. Second measurements (20 ) were performed after 30 min heating at the temperature of the first measurement. (b) Possible assignment of minor reflections in the XRD pattern of the Ce0.89Rh0.16O2y sample heated in air at 1000 8C.

M. Kurnatowska, L. Kepinski / Materials Research Bulletin 48 (2013) 852–862

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high temperature oxidation at 900 8C followed by a low temperature reduction can be used to redisperse a strongly sintered Rh/ CeO2 catalysts. Results of our study as well as data reported in literature suggest that there may be maximum amount of Rh that may be stabilized at the surface of CeO2 at high temperatures in oxidizing atmosphere. It is difficult to evaluate the exact value but rough estimation can be made basing on EDS data. A 3 wt.% Rh2O3 content and mean crystallite size of ceria 37 nm, (corresponding to 23 m2/g surface area) would give a value 10 times higher than a value of 0.25 g Pt for 100 g of CeO2 (surface area 30 m2/g) reported as maximum, monolayer surface coverage for thermally stable Pt–O–Ce surface phase [33]. It should be remembered however that in our case Rh is distributed not only at the surface but also within the bulk of ceria.

Table 2 Structure characterization of the Ce0.89Rh0.11O2y and CeO2 heated in hydrogen.

3.3. Heating in reducing atmosphere

Note: Total amount of Rh in the Ce0.89Rh0.11O2y corresponds to 6.8 wt.% Rh.

XRD patterns of the Ce0.89Rh0.11O2y sample heated in hydrogen at 300–1000 8C for 1 h are shown in Fig. 9a and b. After reduction up to 500 8C only reflections of fluorite type structure were observed. Reduction at 600 8C caused appearance of small, broad feature at 40–418, which at higher reduction temperatures developed into (1 1 1) reflection of metallic rhodium. Reduction treatment at temperatures higher than 300 8C caused also growth of the mean crystallite size of Ce0.89Rh0.11O2y (Fig. 9c) and small shift of the reflections towards higher angles (decrease of the lattice parameter). It is interesting that contrary to what occurred in oxidizing atmosphere the presence of Rh in the ceria lattice only moderately hindered the crystal growth of ceria (Fig. 9c). Refinement of the XRD patterns using X’Pert HighScore Plus program provided the structure parameters and enabled an estimation of the phase composition. Results presented in Table 2 show that the amount of segregated Rh increased from 2.7 wt.% at 800 8C to 5.5 wt.% at 900 8C and 7.4 wt.% at 1000 8C. The latter value is higher than the total content of Rh in the sample (6.8 wt.%) indicating that accuracy of the calculation is rather low (due to low intensity of the Rh reflections). Anyway it suggests that complete phase separation occurred during heating at 1000 8C. These results show that the Ce0.89Rh0.11O2y mixed oxide exhibits higher thermal stability in reducing atmosphere than similar Ce0.89Pd0.11O2y, oxide, where complete phase separation occurred after heating the sample at 800 8C [15]. Microstructure changes of the Ce0.89Rh0.11O2y crystallites after treatment in reducing atmosphere was studied by TEM. In accordance with XRD no noticeable changes appeared after reduction at 400 8C. At 450 8C small dark spots became visible at

some places on ceria crystallites, which could represent rhodium particles (Fig. 10a). Distinct changes occurred after reduction 500 8C such as growth of the size of ceria crystallites and appearance of numerous small (average size 1.2 nm) rhodium metal particles at the surface of ceria crystallites (Fig. 10b). Rh particles could be identified due to the presence of Moire fringes and weak reflections in SEAD patterns (inset to Fig. 10b). The distances between the Moire fringes 0.75 nm agree well with the value of 0.74 nm calculated assuming that (1 1 1) fringes of Ce0.89Rh0.11O2y (0.312 nm) are parallel to (1 1 1) fringes of Rh (0.219 nm). The epitaxial orientation of Rh particles on ceria Rh (1 1 1)jjCeO2 (1 1 1) remained after reduction at 800 8C and only small growth of the Rh particles occurred (Fig. 11a). The same kind and stability of the epitaxial orientation was reported before in Rh/ CeO2 catalysts, where Rh was deposited by impregnation [32,34,35]. The size of rhodium particles was however, bigger and particle size distribution was wider than in our samples, despite much lower total Rh content (2.4 or 0.78 wt.% versus 6.8 wt.%). After heating at 700 8C in H2, mean size of Rh crystallites was 3.7 nm [35] and 3.2 nm [34], while in our sample, the mean size was 2 nm after heating at 800 8C. Severe microstructure changes appeared after reduction of the sample in H2 at 1000 8C. There was a significant growth of the Rh particles connected with strong broadening of the particle size distribution (Fig. 11b). Interestingly, most of the particles were still in epitaxial orientation, though some of them forms twins. Examples of another characteristic microstructure features at the rhodium–ceria interface observed in the sample reduced at

Temperature

a (nm)

Crystallite size (nm)

Ce0.89Rh0.11O2y As prep. 300 8C 500 8C 600 8C 800 8C 900 8C 1000 8C

0.5412 0.5423 0.5418 0.5417 0.5415 0.5416 0.5414

4.2 4.8 12.4 20 30 41 43

CeO2 As prep. 500 8C 800 8C

0.5411 0.5410 0.5411

8.6 9.9 44

a

Rh (wt.%)

Rh crystallite size (nm) XRD

TEM

– – – – 2.7 5.5 7.4

– – – 2.5 5 6

– 1.2 (0.5–2.2) 1.6 (0.5–2.6) 2 (1–3.5) n.m. 4.5 (1–14)

– – –

– – –

– – –

b

c

Fig. 9. XRD patterns of the Ce0.89Rh0.11O2y heated in hydrogen (a) and zoomed region in the vicinity of Rh (1 1 1) reflection (b). Variation of the mean crystallite size with the reduction temperature (c).

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M. Kurnatowska, L. Kepinski / Materials Research Bulletin 48 (2013) 852–862

Fig. 10. HRTEM images of Ce0.89Rh0.11O2y heated in hydrogen at 450 8C (a) and 500 8C (b). SAED pattern with weak Rh reflections is shown as inset.

1000 8C are shown in Fig. 12. Fig. 12a depicts an example of a decoration of Rh particle by ceria crystallites (marked with arrows). This effect was reported already for Rh/CeO2 catalysts prepared by an impregnation method and reduced at high temperature [32,35]. Fig. 12b shows an example of small Rh particle, which has a twinned structure. The FFT from the particle (inset to Fig. 12b) indicates that all three sets of lattice fringes correspond to (1 1 1) lattice planes of Rh. There are two twin boundaries with an angle between normals to the two sets of (1 1 1) faces equal to 408. As it is seen in FFT pattern the twin planes are (1 1 1). Similar situation was observed by STM for gold twinned particles [36]. The authors presented and discussed a hypothesis that the twinned structure may indicate that a transition from an icosahedral structure to a cubic form had occurred as the particle grew in size [36]. Tao et al. [37] proposed that nanoparticles of noble metals often raise in polyhedral structures because of low twinning energies, which can accommodate the strain induced by a completely (1 1 1) bound particle. Images in Fig. 12 c and d present a Rh particle in slightly different orientations. It is seen that the crystal structure of ceria in the vicinity of the Rh particle is strongly disturbed: amorphization or deformation (bending of the lattice fringes shown by arrow) was observed. When comparing the structure evolution of the Ce0.89Rh0.11O2y in oxidizing and reducing atmosphere an important difference can be noticed. In the former the oxide is structurally and chemically stable up to 875 8C with no evidence of any phase separation and rather small growth of the mean crystallite size. At higher temperatures (especially above 950 8C)

there is however, sudden decomposition of the mixed oxide with dramatic sintering of Rh oxide into micron size Rh2O3 crystals. In hydrogen, on the contrary, some decomposition of the mixed oxide with formation very small (1–2 nm) Rh particles begins already at 450 8C but then even after reduction at 1000 8C Rh stays as highly dispersed phase (mean crystallite size 6 nm). It appears also that the Ce0.89Rh0.11O2y shows higher stability in reducing atmosphere than similar Ce0.89Pd0.11O2y [15]. The mean crystallite size of Pd in the latter sample subjected to heating in hydrogen at 800 8C (14 nm) was bigger than the size of Rh particles in the sample heated at 1000 8C. Interestingly, even after such severe treatment most of Rh particles preserved the epitaxial orientation on the support. Calculations of electronic properties of various NM/CeO2 systems (NM-noble metal) revealed that Rh/CeO2 (1 1 1) interface is energetically more stable than the corresponding Pd or Pt/CeO2 (1 1 1) interfaces [38]. This could explain higher stability of rhodium particles on the ceria surface. 3.4. Alternate heating in reducing/oxidizing atmosphere Nanocrystalline Ce0.89Rh0.11O2y and CeO2 were subjected to successive heat treatment at 500 8C in oxygen (for 2 h) and in hydrogen (for 1 h). XRD showed that the first reduction caused noticeable sharpening of the reflections of the mixed oxide but then only minor changes were detected (Fig. S5). No other crystal phase except the fluorite type Ce0.89Rh0.11O2y was seen. Refinement of the XRD profiles revealed however, changes of lattice parameters and of micro strain after successive treatments (Fig. 13,

Fig. 11. HRTEM images of Ce0.89Rh0.11O2y heated in hydrogen at 800 8C (a) and 1000 8C (b). Distributions of rhodium crystallite sizes are shown as insets.

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Fig. 12. HRTEM images of Rh particles in the Ce0.89Rh0.11O2y after heating in hydrogen at 1000 8C: (a) decoration effect, (b) twinned particle, (c, d) amorphization of distortion of ceria close to a Rh particle shown in slightly different orientations.

Table 3). In Fig. 13 step zero refers to the fresh samples, step 1 to reduction, step 2 oxidation and so on. The behavior of the mixed oxide is very interesting. After initial, significant modifications during the first reduction step, expressed by the increase of the mean crystallite size from 4.5 nm to 12 nm, the changes observed in subsequent steps are reversible. Oxidation increases the lattice parameter and the microstrain, whereas reduction decreases these parameters. Such periodic changes were not observed for undoped CeO2, which proves their connection with the presence of Rh in the ceria lattice. Reversible structure changes of the Ce0.89Rh0.11O2y were also observed by TEM. As it has been described above reduction at 500 8C causes occurrence of small Rh particles (1– 2 nm, mean size 1.2 nm) in epitaxial orientation relative to ceria crystallites (Fig. 14a). After oxidation step no Rh crystallites could be seen and also there was not any amorphous or crystalline phase that could be assigned to rhodium oxide (Fig. 14b). Next reduction

step caused reappearance of Rh crystallites with morphology very similar to that observed after step 1. Mean crystallite size was almost the same (1.4 nm) (Fig. 14c) and the crystallites exhibited preferential orientation on the support. The shape of the Rh particles (height/diameter < 1) imply strong interaction with the support. Next oxidation step caused disappearance of the Rh crystallites and also significant disorder at the surface of the ceria crystallites. The surface of the ceria crystallites is very rough after the oxidation steps and contains numerous defects such as bended lattice fringes or kinks (Fig. 14b and d). Above results suggest reversible diffusion of rhodium from the ceria lattice to the surface during reduction and its incorporation back into the lattice during oxidation.

Table 3 Structure characterization of the Ce0.89Rh0.11O2y and CeO2 subjected to reduction–oxidation treatment. Treatment

a (nm)

Crystallite size (nm)

Strain (%)

Ce0.89Rh0.11O2y 0 ‘‘As prepared’’ 1 Reduced 500 8C H2 2 Re-oxidized 500 8C O2 3 Reduced 500 8C H2 4 Re-oxidized 500 8C O2 5 Reduced 500 8C H2

0.5411 0.5418 0.5421 0.5415 0.5420 0.5415

4.5 12.4 12.3 13.3 13.2 13.9

0.87 0 0.18 0 0.14 0

CeO2 0 1 2 3 4

0.54115 0.54104 0.54102 0.54104 0.54113

8.6 9.9 9.8 11.7 10.3

0.26 0.11 0 0 0

t

Step

Fig. 13. Changes of the microstructure parameters of the Ce0.89Rh0.11O2y and CeO2 heated alternately in H2 and O2 at 500 8C: 0 – fresh sample, 1 – first reduction, 2 – reoxidation and so on. Micro strain (Ce0.89Rh0.11O2y – circle, CeO2 – triangle) and lattice parameter (Ce0.89Rh0.11O2y – square, CeO2 hexagon).

‘‘As prepared’’ Reduced 500 8C H2 Re-oxidized 500 8C O2 Reduced 500 8C H2 Re-oxidized 500 8C O2

860

M. Kurnatowska, L. Kepinski / Materials Research Bulletin 48 (2013) 852–862

Fig. 14. HRTEM images of the Ce0.89Rh0.11O2y heated alternately in H2 and O2 at 500 8C. First reduction (a), re-oxidation (b), third reduction (c) and re-oxidation (d). Rh particles on the oxide surface are marked with arrows.

790 825

835

460 510

c

450

b

335

113

318 376

163

a

148

In our previous work [15] we reported similar effect of reversible extraction-dissolution of Pd in Ce0.89Pd0.11O2y. The changes occurring after successive reduction–oxidation cycles were however more pronounced (visible in XRD patterns) due to lower stability of Ce0.89Pd0.11O2y. Ciston et al. [39] studied ‘‘in situ’’ the behavior Ce0.8Cu0.2O2y solid solution upon change of gas atmosphere from reducing to oxidizing at 300 8C using TEM and XRD. As in our case, precipitation of metallic Cu was observed on reduction. Under oxidizing conditions the authors observed however, that only 50% of Cu returned to a Ce1xCuxO2 solid solution, while the remainder was observed by in situ TEM to form an amorphous CuOx phase. The differences may be due to higher initial Cu doping and weaker bonding of Cu in ceria lattice. Basing on ‘‘in situ’’ XRD the authors of Ref. [39] reported that in the initial sample 25% of Cu was present at the surface as Cu oxide and already after reduction at 160 8C the Cu dopant is rapidly and entirely expelled from the solid solution. The size of Cu particles after heating in H2 at 400 8C (10 nm) was also much bigger than in our case. The concept of reversible extraction – dissolution of metal phase in complex oxides is a basis of so called ‘‘self-regenerative property’’ discovered for Pd-doped perovskite catalyst for automotive emission control [40]. In a recent TEM study Katz et al. [41] showed that the process of extraction–dissolution of Pd in LaFe0.95Pd0.05O3d under high-temperature oxidizing and reducing conditions, was quite limited and restricted to a topmost (few nm) layer of the mixed oxide. 3.5. H2-TPR Fig. 15 depicts two successive H2 TPR profiles obtained for Ce0.89Rh0.11O2y in comparison with similar profiles reported for

Fig. 15. H2-TPR profiles (I and II run) of the Ce0.89Rh0.11O2y (a), Ce0.89Pd0.11O2y (b) and CeO2 (c). Temperatures of the maxima in the first runs are given.

M. Kurnatowska, L. Kepinski / Materials Research Bulletin 48 (2013) 852–862

Ce0.89Pd0.11O2y [15] and CeO2 [42] synthesized in a similar way. After the first run up to 900 8C the samples were outgassed and then oxidized at 500 8C for 1 h. Three regions can be distinguished in the profiles of Ce0.89Rh0.11O2y (Fig. 15a): low temperature (below 200 8C), intermediate (200–500 8C) and high temperature (above 500 8C). The first run contains two bands in the low temperature region: weak band at 113 8C and very strong one at 163 8C. According to literature [4,12] pure Rh2O3 reduces in a single band at 150–180 8C, and the difference may be due to different dispersion of the oxide and/or experimental details. For Rh/CeO2 catalysts prepared by impregnation generally two bands are reported in the low temperature region. Mizuno et al. [43] reported for 5.4 wt.% Rh/CeO2 catalyst the presence of bands at 77 and 137 8C. The authors assign the latter band to reduction of Rh2O3 not interacting with the support (the band is similar to that in Rh supported on Al2O3 and SiO2) while the former band at 77 8C is due to reduction of Rh2O3 in close contact with the support (CeO2 promotes the Rh2O3 reduction). Cai et al. observed for 1 wt.% Rh/ CeO2 catalyst a strong, single band at 140 8C. Though its position is similar to that for Rh2O3, the intensity is much to high to be assigned to bare Rh2O3. The authors suggest that this band must involve also surface reduction of CeO2 support in vicinity of Rh, facilitated by the presence of metallic Rh (hydrogen spill-over). Bueno-Lopez et al. [44] reported H2-TPR profiles of 0.5 wt.% Rh supported on CeO2 of various surface area. In each case a single band occurred at 100 8C, but its intensity increased strongly for the catalysts supported on high surface CeO2. The authors estimated, that up to 11% of Ce4+ is reduced to Ce3+ together with Rh in the low temperature band in the catalyst supported on high surface ceria. In the medium temperature region (200–500 8C) two weak reduction bands are seen at 318 and 376 8C. Finally, the broad band at around 790 8C is due to bulk reduction of ceria. Completely different profile appeared in the second TPR run. Instead of a very strong band at 160 8C there are much weaker multiple bands in the temperature range up to 250 8C. There is also some reduction of intensity of the middle temperature bands. All these changes are caused primary by sintering of ceria during the first TPR run accompanied with extraction of Rh from the ceria lattice. It is worth to stress that the first TPR profile of our Ce0.89Rh0.11O2y closely resembles that of Ce0.98Rh0.02O2y prepared by a solution combustion method [12] but differs from that of Ce0.89Pd0.11O2y (Fig. 15b) [15]. In particular for Pd doped ceria the intensity ratio of the low temperature and middle temperature bands is much lower. Moreover, middle and high temperature reduction bands occur in Ce0.89Rh0.11O2y at lower temperatures. These observations indicate that rhodium is more effective than Pd as a dopant promoting the ceria reduction. High intensity of the low temperature reduction bands for the Rh and Pd doped ceria as well as the fact that the ratio of middle/high temperature reduction bands for the doped ceria is smaller than for bare ceria strongly suggest that in the former some reduction of Ce4+ occurs simultaneously with Rh3+ or Pd2+ below 200 8C. Dramatic change of the TPR profiles of the doped samples recorded during the second run confirms the observations presented above that treatment in hydrogen at high temperatures (900 8C) causes extensive modification of the structure of the samples (phase decomposition). 4. Conclusions Nanocrystalline Ce1xRhxO2y solid solutions (x = 0.03–0.16) with mean crystallite size 4–5 nm and narrow size distribution was prepared using microemulsion method. All the samples were stable in oxidizing atmosphere up to about 850 8C, showing no phase separation and only limited growth of the mean crystallite size. At higher temperatures (950 8C) phase separation into Rh

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deficient mixed oxide and large (few mm) Rh2O3 crystals was observed indicating that Rh solution limit in ceria lattice decreaes with temperature (crude estimation gives about 3 at% at 1000 8C). In hydrogen segregation of Rh as small crystallites (1–2 nm) in epitaxial orientation relative to ceria lattice Rh (1 1 1)jjCeO2 (1 1 1) started at 500 8C. Only limited growth of the Rh particles size occurred up to 800 8C what confirms that rhodium is strongly bonded with the ceria support. After treatment at 1000 8C strong, chemical interaction with ceria support caused decoration of Rh particles by the support. Reversible diffusion of rhodium from the ceria lattice to the surface during reduction and its incorporation back into the lattice during oxidation was observed during repeated reduction/oxidation cycling at 500 8C. The effect was observed also for Pd doped ceria [15] and is an example of so called ‘‘self-regenerative property’’ discovered for Pd-doped perovskite [40]. Nanocrystalline Ce0.89Rh0.11O2y exhibits strongly enhanced reducibility (below 200 8C) as compared to pure CeO2 what in addition to high structural and chemical stability makes it an interesting catalytic material for reactions were fluctuating oxidation-reducing conditions are expected. Acknowledgments This work was financially supported by the National Science Center in Poland (grant 2011/01/N/ST5/05609). The authors thank Dr A. Ga˛gor and Mrs. E. Bukowska for XRD work, Mr. M. Ptak for recording IR and Raman spectra and Dr. W. Mista for carrying out TPR experiments. Appendix A. Supplementary data Supplementary data associated with this article can be found, in the online version, at http://dx.doi.org/10.1016/j.materresbull.2012. 11.076. References [1] M.A. Newton, Chem. Soc. Rev. 37 (2008) 2644–2657. [2] S. Hosokawa, M. Taniguchi, K. Utani, H. Kanai, S. Imamura, Appl. Catal. A 289 (2005) 115–120. [3] F. Fajardie, J.F. Tempere, J.M. Manoli, O. Touret, G. Blanchard, G. Djega-Mariadassou, J. Catal. 179 (1998) 469–476. [4] W. Cai, F. Wang, A.C. Van Veen, H. Provendier, C. Mirodatos, W. Shen, Catal. Today 138 (2008) 152–156. [5] M. Boaro, V. Modafferi, A. Pappacena, J. Llorca, V. Baglio, F. Frusteri, P. Frontera, A. Trovarelli, P.L. Antonucci, J. Power Sources 195 (2010) 649–661. [6] N. Bion, F. Epron, M. Moreno, F. Marino, D. Duprez, Top. Catal. 51 (2008) 76–88. [7] C. Galletti, S. Fiorot, S. Specchia, G. Saracco, V. Specchia, Top. Catal. 45 (2007) 15– 19; A. Yee, S.J. Morrison, H. Idriss, Catal. Today 63 (2000) 327–335. [8] T. Miyazawa, K. Okumura, K. Kunimori, K. Tomishige, J. Phys. Chem. C 112 (2008) 2574–2583. [9] J. Chen, H. Jiang, H. Qian, M. Malac, J. Phys. Chem. C 115 (2011) 14173–14179. [10] X. Tang, B. Zhang, Y. Li, Y. Xu, Q. Xin, W. Shen, Catal. Lett. 97 (2004) 163–169. [11] M.S. Hedge, G. Madras, K.C. Patil, Acc. Chem. Res. 42 (2009) 704–712. [12] A. Gayen, K.R. Priolkar, P.R. Sarode, V. Jayaram, M.S. Hegde, G.N. Subbanna, S. Emura, Chem. Mater. 16 (2004) 2317–2328. [13] P. Bera, K.R. Priolkar, A. Gayen, P.R. Sarode, M.S. Hegde, S. Emura, R. Kumashiro, V. Jayaram, G.N. Subbanna, Chem. Mater. 15 (2003) 2049–2060. [14] K.R. Priolkar, P. Bera, P.R. Sarode, M.S. Hedge, S. Emura, R. Kumashiro, N.P. Lalla, Chem. Mater. 14 (2002) 2120–2128. [15] M. Kurnatowska, L. Kepinski, W. Mista, Appl. Catal. B: Environ. 117–118 (2012) 135–147. [16] M.A. Małecka, L. Kepinski, W. Mis´ta, J. Alloys Compd. 451 (2008) 567–570. [17] S. Music, J. Mol. Struct. 924–926 (2009) 221–224. [18] Y. Zhou, M.N. Rahaman, J. Mater. Sci. Technol. 11 (1995) 429–434. [19] J.-S. Lee, S.-C. Choi, J.-S. Lee, S.-C. Choi, Mater. Lett. 58 (2004) 390–393. [20] M.A. Małecka, U. Burkhardt, D. Kaczorowski, M.P. Schmidt, D. Goran, L. Kepinski, J. Nanopart. Res. 11 (2009) 2113–2124. [21] A. Gupta, U.V. Waghmare, M.S. Hegde, Chem. Mater. 22 (2010) 5184–5198. [22] L.A. Carol, G.S. Mann, Oxid. Met. 34 (1990) 1–12. [23] H. Mizoguchi, L.N. Zakharov, N.S.P. Bhuvanesh, A.W. Sleight, M.A. Subramanian, J. Solid State Chem. 184 (2011) 1381–1386. [24] O. Muller, R. Roy, J. Less-Common Met. 16 (1968) 129.

862 [25] [26] [27] [28] [29] [30] [31] [32] [33] [34]

M. Kurnatowska, L. Kepinski / Materials Research Bulletin 48 (2013) 852–862 E. Ruckenstein, H.Y. Wang, J. Catal. 190 (1999) 32–38. C.P. Hwang, C.T. Yeh, Q.M. Zhu, Catal. Today 51 (1999) 93–101. R.B. Macquart, M.D. Smith, H.C. zur Loye, Cryst. Growth Des. 6 (2006) 1361–1365. T. Taniguchi, W. Iizuka, Y. Nagata, T. Uchida, H. Samata, J. Alloys Compd. 350 (2003) 24–29. V.B. Lazarev, I.S. Shaplygin, Russ. J. Inorg. Chem. 23 (1978) 1449–1452. B. Li, O.K. Ezekoye, Q. Zhang, L. Chen, P. Cui, G. Graham, X. Pan, Phys. Rev. B 82 (2010) 125422. S. Imamura, T. Yamashita, R. Hamada, Y. Saito, Y. Nakao, N. Tsuda, C. Kaito, J. Mol. Catal. A: Chem. 129 (1998) 249–256. S. Bernal, F.J. Botana, J.J. Calvino, G.A. Cifredo, J.A. Perez-Omil, J.M. Pintado, Catal. Today 23 (1995) 219–250. M. Hatanaka, N. Takahashi, T. Tanabe, Y. Nagai, K. Dohmae, Y. Aoki, T. Yoshida, H. Shinjoh, Appl. Catal. B 99 (2010) 336–342. J.M. Gatica, R.T. Baker, P. Fornasiero, S. Bernal, G. Blanco, J. Kaspar, J. Phys. Chem. B 104 (2000) 4667–4672.

[35] S. Bernal, J.J. Calvino, M.A. Cauqui, J.A.P. Omil, J.M. Pintado, J.M. RodriguezIzquierdo, Appl. Catal. B: Environ. 16 (1998) 127–138. [36] C.H. Pontifex, P. Zbang, Z. Wing, T.L. Haslett, D.A.I. Mawlawi, M. Moskovits, J. Phys. Chem. 95 (1991) 9989–9993. [37] A.R. Tao, S. Habas, P. Yang, Small 4 (2008) 310–325. [38] Z. Lu, Z. Yang, J. Phys. Cond. Matter 22 (2010) 475003. [39] J. Ciston, R. Si, J.A. Rodriguez, J.C. Hanson, A. Martinez-Arias, M. Fernandez-Garcia, Y.M. Zhu, J. Phys. Chem. C 115 (2011) 13851–13859. [40] Y. Ishihata, J. Mizuki, T. Akao, H. Tanaka, M. Uenishi, M. Kimura, T. Okamoto, N. Hamada, Nature 418 (2002) 164–167. [41] M.B. Katz, G.W. Graham, Y. Duan, H. Liu, C. Adamo, D.G. Schlom, X. Pan, J. Am. Chem. Soc. 133 (2011) 18090–18093. [42] M. Małecka, L. Kepinski, W. Mis´ta, Appl. Catal. B: Environ. 74 (2007) 290–298. [43] T. Mizuno, Y. Matsumura, T. Nakajima, S. Mishima, Int. J. Hydrogen Energy 28 (2003) 1339–1399. [44] A. Bueno-Lopez, I. Such-Basanez, C.S.M. de Lecea, J. Catal. 244 (2006) 102–112.