Hydrogen Production through Bi-Reforming of Methane: Improving Ni Catalyst Performance via an Exsolution Approach

: Hydrogen production through the bi-reforming of methane over exsolution-derived Ni catalysts has been studied. Nickel-based catalysts were prepared through the activation of (CeM) 1 − x Ni x O y (M = Al, La, Mg) solid solutions in a reducing gaseous medium. Their performance and resistance to coking under the reaction conditions were controlled by regulating their textural, structural, morphological, and redox properties through adjustments to the composition of the oxide matrix (M/Ce = 0–4; x = 0.2–0.8; y = 1.0–2.0). The role of the M-dopant type in the genesis and properties of the catalysts was established. The efﬁciency of the catalysts in the bi-reforming of methane increased in the following series of M: M-free < La < Al < Mg, correlating with the structural behavior of the nickel active component and the anti-coking properties of the support matrix. The preferred M-type and M/Ce ratio determined the best performance of (CeM) 1 − x Ni x O y catalysts. At 800 ◦ C the optimum Ce 0.6 Mg 0.2 Ni 0.2 O 1.6 catalyst provided a stable H 2 yield of 90% at a high level of CO 2 and CH 4 conversions (>85%).


Introduction
Hydrogen is a versatile energy carrier and an important industrial raw material. About 75 million tons (Mt) of pure H 2 and 45 Mt of H 2 mixed with other gases are consumed annually [1]. The main consumer of pure hydrogen is the chemical sector (39 Mt), where H 2 is used in the production of ammonia, and oil refining (33 Mt), where it is needed for the hydrotreating and hydrocracking of oil [2]. Methanol synthesis and steel production are important sources of demand for hydrogen mixed with other gases. The large-scale use of hydrogen as an energy carrier requires the solution of various issues, including its efficient production, economical purification, safe storage, and transportation [3][4][5].
For H 2 production, conventional and renewable technologies are applied based on the use of fossil fuels (natural gas, heavy oils, naphtha, coal, etc.) and renewable resources (water, biomass, etc.), respectively [6][7][8]. At present, hydrogen is generally produced from natural gas ("gray hydrogen") and coal ("brown hydrogen"), with small contributions from oil and electricity. "Gray hydrogen" is mainly manufactured through the steam reforming of methane (SRM, Reactions (1) and (2)), which is accomplished by CO 2 emission [9]: Today, H 2 production from fossil fuels emits about 900 Mt of CO 2 [10]. It is predicted that by 2050, the production of hydrogen from natural gas will be almost entirely based on low-carbon technologies: water electrolysis accounts for more than 60% of world production, and natural gas in combination with CCUS (carbon capture, utilization, and storage) accounts for almost 40% ("blue hydrogen") [10]. However, the future of the latter technology is doubtful [6,9,11], since greenhouse gas footprints from "blue hydrogen" are only 18-20% less than those from "gray hydrogen" due to the use of fossil C-derived energy to implement the CCUS technology. Thus, the improvement of cost-effective processes of converting methane into a hydrogen-containing gas is still relevant.
The combined steam and carbon dioxide reforming of methane (bi-reforming of methane, Reaction (3)) is an attractive way to catalytically generate hydrogen-containing gas from natural gas with CO 2 utilization [12][13][14]: This is a unique environmentally and climate-friendly process that allows one to simultaneously utilize three greenhouse gases (carbon dioxide, methane, and water vapor). The obtained syngas can be used directly in certain processes, such as methanol production or the direct reduction of iron, or they can be used to produce pure hydrogen. In this case, concentrated CO streams can be captured and used in the production of formic acid, acetic acid, or functional copolymers [15][16][17][18][19]. The bi-reforming of methane is a complex process [20,21]. It involves the reactions of steam (1) and (2) and carbon dioxide (4) and the conversion of methane, leading to the production of the target products (synthesis gas or hydrogen), a number of additional reactions that increase the yield of hydrogen (5) and (6), and side reactions of carbon formation (6)- (9 (9) Extensive application in the reforming process was given to Ni-based catalysts due to their low cost, wide availability, and sufficiently high activity [22][23][24][25][26][27][28]. An obstacle that arises when using nickel catalysts is their deactivation under harsh reaction conditions and the impossibility of self-activation [29]. There are various strategies for improving the activity and stability of nickel catalysts: optimization of preparation mode [30,31] and support composition [32,33], application of different promoters (Mo, Pd, Pt, Rh, Re, or Mo) [34,35], variation of calcination and activation procedure [36][37][38][39]. For example, using La 2 O 3 or MgO as basic support modifiers can prevent Ni agglomeration and promote CO 2 activation, which increases the activity and stability of the catalyst in the dry reforming of methane (DRM) [40,41].
Critical factors controlling catalyst performance and the rate of carbon deposits formation include the strength of metal-support interaction (MSI) and the nickel particle size. The Ni particle size significantly affects the carbon formation and removal reaction paths in DRM and, consequently, the amount and the structure/morphology of carbon deposits [42]. The amount of carbon deposition is insignificant for Ni crystallites sizes less than 10 nm and more than 100 nm and is greatest for Ni of about 20-30 nm in size [36]. On Ni particles of medium size (20-45 nm), the formation of carbon in the form of nanotubes prevails, and on small particles (<10 nm), amorphous carbon is typically found [36,42]. Carbon species (C-S ) formed during DRM can be removed through reaction with labile lattice oxygen (O-L ) of support and creating oxygen vacancies (V o ) that are replenished by the dissociative adsorption of CO 2 on the support (10) and (11) [42]:

Characterization of Fresh Samples
The chemical composition and textural properties of fresh (CeM) 1−x Ni x O y samples are given in Table 1. The chemical composition of the samples is in good agreement with the theoretically defined values.
Three series of (CeM) 1−x Ni x O y samples were obtained using the polymerizable complex method, differing in the type (M = Al, La, Mg) and content (M/Ce = 0-4; x = 0.2; y = 1.0-2.0) of the dopant cation M, and one series based on an unmodified ceria, where the Ni/Ce molar ratio was varied over a wide range (0.01-4) and x was changed from 0.01 to 0.8.
When obtaining multicomponent solid solutions through the PC method (Figure 1), all metal cations are mixed at the molecular level in the water solution and then CA-complexes are homogeneously distributed in the gel and polymer resin. This preserves the initial stoichiometry of the molar ratio of cations in the final material, but does not exclude the coexistence of several amorphous or crystalline phases after the procedure of its thermal treatment [72].  The TA method was used to study the genesis of materials from a gel into a soli phase during thermal treatment. As typical example, Figure 2a illustrates the TG, DTG and DTA curves for a Ce0.8Ni0.2O1.8 sample. There are three main peaks on the DTG curv The first peak at T < 200 °C is accompanied by an endothermic effect and can be associate with the desorption of water and volatile organic compounds. The second peak at T ≈ 30 °C and the third peak at 400 °C < T < 600 °C match well with the exothermic events ob  The TA method was used to study the genesis of materials from a gel into a solid phase during thermal treatment. As typical example, Figure 2a illustrates the TG, DTG, and DTA curves for a Ce 0.8 Ni 0.2 O 1.8 sample. There are three main peaks on the DTG curve. The first peak at T < 200 • C is accompanied by an endothermic effect and can be associated with the desorption of water and volatile organic compounds. The second peak at T ≈ 300 • C and the third peak at 400 • C < T < 600 • C match well with the exothermic events observed in the DTA curve and can be assigned to multistage oxidation of the polymer matrix. The porous structure of the catalysts is formed under the condition of burnout of the organic matrix, accompanied by the release of gaseous combustion products. Data on lowtemperature nitrogen adsorption indicate that (CeM)1−xNixOy are mesoporous materials. As evidenced by the type IV adsorption isotherm with an H3 hysteresis loop at P/P0 > 0.6 ( Figure 3), the samples have predominantly textural mesoporosity (pores between primary particles) [75,76]. The specific surface area of the Ce1−xNixOy samples practically does not change when the mole fraction of Ni varies within 0.05-0.8 and amounts to 110 ± 10 m 2 /g (Table 1), which represents a typical value for the Ce-Ni-O system [12,77]. It is note- For (CeM) 1−x Ni x O y samples, compared to Ce 1−x Ni x O y , the profiles of the TG, DTG, and DTA curves have the similar behavior, but the temperature of organics burn-out shifts to the high-temperature region (Figure 2b). Thus, the introduction of M does not lead to a change in the sequence of decomposition of the organic matrix but rather slows it down. This effect increases in the series La < Al ≈ Mg and can be explained by differences in (i) the resistance of citrate-metal complexes to degradation and (ii) the catalytic activity of metals in the oxidation of organic compounds. The stability of citrate metal complexes increases in the series Mg (lgk 1 = 4.0) < Ni (lgk 1 = 5.4) < Ce (lgk 1 = 7.4) < Al (lgk 1 = 8.0) < La (lgk 1 = 8.4) [73], which has no correlation with the temperature of organic matrix degradation (Figure 2b). Apparently, the presence of a redox Ce 4+ /Ce 3+ pair accelerates the oxidation reactions of organic residues, which reduces their burnout temperature with increasing content of Ce in the gel. Indeed, cerium-free samples have the highest organic matrix burnout temperature (Figure 2b). Note that, according to [74], hindered organic decomposition has a positive effect on the dispersion of the crystalline phase due to the capping around the particles and their protection from sintering. Accordingly, it can be expected that differences in the parameters of the thermal genesis of materials will affect their textural and/or structural properties.
The porous structure of the catalysts is formed under the condition of burnout of the organic matrix, accompanied by the release of gaseous combustion products. Data on lowtemperature nitrogen adsorption indicate that (CeM) 1−x Ni x O y are mesoporous materials. As evidenced by the type IV adsorption isotherm with an H3 hysteresis loop at P/P 0 > 0.6 ( Figure 3), the samples have predominantly textural mesoporosity (pores between primary particles) [75,76]. The specific surface area of the Ce 1−x Ni x O y samples practically does not change when the mole fraction of Ni varies within 0.05-0.8 and amounts to 110 ± 10 m 2 /g (Table 1), which represents a typical value for the Ce-Ni-O system [12,77]. It is noteworthy that, unlike samples obtained by impregnation, even at a high nickel content, the S BET and V p of the samples are higher than for unmodified ceria with S BET = 83 m 2 /g ( Figure 4a). In particular, at x = 0.3 (15 wt.% Ni) S BET is equal to 117 and 65 for samples obtained by the polymerizable complex method and impregnation, respectively. The specific surface area and total pore volume of (CeM) 1−x Ni x O y samples increased with decreasing M/Ce molar ratio (Table 1, Figure 4b). This trend intensifies in the series of M: Mg < La < Al. The average pore diameter depends on the composition of the material and varies from 7.3 to 22 nm. It increases with an increase in the M/Ce molar ratio. There are inhibition as well as promotion effects from dopants on the textural characteristics. In our case, the incorporation of Ni cations improves the textural properties of CeO 2 , while for (CeM) 1−x Ni x O y samples only for samples with M/Ce equal to 0.3, the textural characteristics are higher than for unmodified ceria. The resultant dependence of the textural properties of the samples upon their chemical composition may reflect the effects of dopant cations (Ni, Al, La, and Mg) of ceria on the sintering mode of the material and, consequently, phase dispersion and material porosity [77][78][79]. surface area and total pore volume of (CeM)1−xNixOy samples increased with decreasing M/Ce molar ratio (Table 1, Figure 4b). This trend intensifies in the series of M: Mg < La < Al. The average pore diameter depends on the composition of the material and varies from 7.3 to 22 nm. It increases with an increase in the M/Ce molar ratio. There are inhibition as well as promotion effects from dopants on the textural characteristics. In our case, the incorporation of Ni cations improves the textural properties of CeO2, while for (CeM)1−xNixOy samples only for samples with M/Ce equal to 0.3, the textural characteristics are higher than for unmodified ceria. The resultant dependence of the textural properties of the samples upon their chemical composition may reflect the effects of dopant cations (Ni, Al, La, and Mg) of ceria on the sintering mode of the material and, consequently, phase dispersion and material porosity [77][78][79].    The morphological properties of the samples are close. They are aggregates of spongy particles ( Figure 5). The formation of their porous macrostructure occurs under conditions of intense gas evolution due to the spontaneous burnout of the organic matrix. The morphological properties of the samples are close. They are aggregates of spongy particles ( Figure 5). The formation of their porous macrostructure occurs under conditions of intense gas evolution due to the spontaneous burnout of the organic matrix. Table 2 demonstrates the phase and structure data for (CeM) 1−x Ni x O y samples. X-ray diffraction analysis indicates the presence of a CeO 2 -based cubic phase in all Ce-containing samples ( Figure 6). The only exception is the Ce 0.2 Al 0.7 Ni 0.1 O 1.5 sample, for which the XRD pattern of an amorphous phase is observed. The Ni-containing phase is observed at a high nickel content (x = 0.7, 0.8) and also at M = Mg. In the first case, highly dispersed particles of NiO with an average size of~5 nm are formed, and in the second case, a solid solution of NiO-MgO with a particle size of 4-8 nm is observed (Table 2, Figure 6). In the cerium-free samples-La 0. 8   The morphological properties of the samples are close. They are aggregates of spongy particles ( Figure 5). The formation of their porous macrostructure occurs under conditions of intense gas evolution due to the spontaneous burnout of the organic matrix.  Table 2 demonstrates the phase and structure data for (CeM)1−xNixOy samples. X-ray diffraction analysis indicates the presence of a CeO2-based cubic phase in all Ce-containing samples ( Figure 6). The only exception is the Ce0.2Al0.7Ni0.1O1.5 sample, for which the XRD pattern of an amorphous phase is observed. The Ni-containing phase is observed at a high nickel content (x = 0.7, 0.8) and also at M = Mg. In the first case, highly dispersed particles of NiO with an average size of ~5 nm are formed, and in the second case, a solid solution of NiO-MgO with a particle size of 4-8 nm is observed (Table 2, Figure 6). In the     The structural characteristics of a CeO 2 -based phase depend on the content (x = 0.001-0.8, M/Ce = 0-4), ionic radius (r Ni2+ = 0.69 Å, r Al3+ = 0.68 Å, r La2+ = 1.16 Å, r Mg2+ = 0.72 Å), and degree of oxidation (2+ for Ni and Mg cations, 3+ for Al and La cations) of the dopant. It can be seen that in case of doping cations (Ni, Al, Mg) with a radius smaller than that of Ce 4+ (r Ce4+ = 0.97 Å), the unit cell parameter a of cubic CeO 2 -based phase is 0.5412 ± 0.0004 nm and differs little from that for unmodified CeO 2 . On the contrary, in the case of La, which has a larger radius, a noticeable increase in the unit cell parameter is observed (Table 2), which can indicate the incorporation of La 3+ cations into the ceria lattice. For all samples, a decrease in the size of crystallites is observed in comparison with unmodified ceria. This effect is enhanced with increasing x, increasing M/Ce ratio and in the series La→Mg→Al (Figure 7), which correlates with the difficulty of burnout of the organic matrix ( Figure 2b) and indicates the inhibitory effect of the cations introduced into the ceria structure on the growth of crystallites and an increase in the thermal stability to sintering of (CeM) 1−x Ni x O y samples.  In the case of (CeLa)1−xNixOy samples, the absence of a nickel-containing phase, a decrease in the particle size, and an increase in the lattice parameter of the CeO2-based phase clearly points to the formation of a solid solution. For the Ce1−xNixOy and (CeAl)1−xNixOyseries, no Ni-containing phases were observed either, and the CSR size for CeO2 decreases with increasing x or M/Ce ratio. However, for these samples, the cell parameters remain virtually unchanged (Table 2). Based on the XRD data, it can be assumed that a solid solution is also formed in this case. But since the Ce 3+ /Ce 4+ ratio increases with decreasing particle size [80], an increase in the concentration of Ce 3+ with bigger ionic radius (rCe3+ = 1.14 Å) compensates for the expected decrease in the cell parameter of a CeO2-based solid solution. According to [80], the cell parameter a increases from 0.541 to 0.554 Å as the CeO2 particle size decreases from 70 to 3 nm. It was also shown [70] that differences in the formal charge of Ni 2+ and Ce 4+ promote the formation of oxygen vacancies and, as a consequence, lattice expansion. Less obvious from the XRD data is the formation of a solid solution for the (CeMg)1−xNixOy-series. In this case, although there is a decrease in the size of crystallites of the CeO2-based phase, a NiO-MgO solid solution is present in the sample. Its unit cell parameter has values in the range of 0.4182-0.4212 nm, which lies between the values characteristic of individual oxides (0.4177 for NiO and 0.4213 for MgO) and In the case of (CeLa) 1−x Ni x O y samples, the absence of a nickel-containing phase, a decrease in the particle size, and an increase in the lattice parameter of the CeO 2 -based phase clearly points to the formation of a solid solution. For the Ce 1−x Ni x O y and (CeAl) 1−x Ni x O yseries, no Ni-containing phases were observed either, and the CSR size for CeO 2 decreases with increasing x or M/Ce ratio. However, for these samples, the cell parameters remain virtually unchanged ( Table 2). Based on the XRD data, it can be assumed that a solid solution is also formed in this case. But since the Ce 3+ /Ce 4+ ratio increases with decreasing particle size [80], an increase in the concentration of Ce 3+ with bigger ionic radius (r Ce3+ = 1.14 Å) compensates for the expected decrease in the cell parameter of a CeO 2 -based solid solution. According to [80], the cell parameter a increases from 0.541 to 0.554 Å as the CeO 2 particle size decreases from 70 to 3 nm. It was also shown [70] that differences in the formal charge of Ni 2+ and Ce 4+ promote the formation of oxygen vacancies and, as a consequence, lattice expansion. Less obvious from the XRD data is the formation of a solid solution for the (CeMg) 1−x Ni x O y -series. In this case, although there is a decrease in the size of crystallites of the CeO 2 -based phase, a NiO-MgO solid solution is present in the sample. Its unit cell parameter has values in the range of 0.4182-0.4212 nm, which lies between the values characteristic of individual oxides (0.4177 for NiO and 0.4213 for MgO) and increases with increasing Mg content in the sample ( Table 2). This may indicate only a partial incorporation of Ni 2+ and Mg 2+ doping cations into the ceria lattice. The solubility limit of a dopant in the ceria lattice with the formation of a solid solution depends on the type of dopant and the method of its introduction. For Ce 1−x M x O y , a single-phase fluorite-like system is observed up to x = 0.1-0.5 for M = Ni [69,81,82], x = 0.6 for M = La [83][84][85] and x = 0.5-0.9 for M = Mg [86,87]. In addition, among the studied dopants (M = Al, La, and Mg), the Mg cations are characterized by the lowest stability of citrate complexes, which probably contributes to their less-successful entry into the ceria lattice.
Raman spectroscopy data provide additional information about the structural properties of the obtained samples (Table 2, Figure 8).
type of dopant and the method of its introduction. For Ce1−xMxOy, a single-phase fluoritelike system is observed up to x = 0.1-0.5 for M = Ni [69,81,82], x = 0.6 for M = La [83][84][85] and x = 0.5-0.9 for M = Mg [86,87]. In addition, among the studied dopants (M = Al, La, and Mg), the Mg cations are characterized by the lowest stability of citrate complexes, which probably contributes to their less-successful entry into the ceria lattice.
Raman spectroscopy data provide additional information about the structural properties of the obtained samples (Table 2, Figure 8). The spectrum of unmodified CeO2 is characterized by one strong band at 465 cm −1 , which is due to the F2g vibration mode of the cubic fluorite structure [88,89]. In addition, it also exhibits two weak bands at about 250 and 600 cm −1 , which can be attributed to the non-degenerate transverse and longitudinal optical phonon modes of CeO2, respectively [90]. The Raman spectral features of the (CeM)1−xNixOy samples were greatly affected by the content and type of dopant ( Figure 8). For these samples, Raman spectra also include main bands at 452-464 cm −1 and two weak bands at 230 cm −1 and 550-650 cm −1 , which can be attributed to the displacement of oxygen atoms from their ideal fluorite lattice positions [91] and the presence of defects such as oxygen vacancies [92]. This shows [93] that the emergence of the band at about 560 cm −1 (oxygen vacancies) results from the different oxidation state of the dopant compared to that of Ce 4+ , while the band at about 600 cm −1 is caused by the different ionic radius of the dopant compared to that of Ce 4+ . The I570/I465 ratio, used to estimate the concentration of oxygen vacancies in a material [94], increases with increasing x, increasing M/Ce ratio, and in the series Mg → Al →La ( Table 2). The higher the I570/I465, the more oxygen vacancies in the material are formed. In addition, a shift (465 → 452 cm −1 ) of the position of the main band in Raman spectra for (CeM)1−xNixOy samples compared to the position of this band for pure CeO2 was noted. This indicates a change in the binding energies due to the incorporation of Ni 2+ and M n+ cations into the lattice of CeO2 and confirms the formation of a solid solution. The considerable broadening of the bands observed for (CeM)1−xNixOy ( Figure 8) could be attributed to the reduction in phonon lifetime due to the decrease in grain size [88], which correlates well with XRD data ( Figure 6). Significant deformation of the spectrum shape for La-containing samples at La/Ce ≥ 1 points to a substantial rearrangement of the oxide structure compared to that of unmodified ceria. It was demonstrated [95] that with the increase of La concentration in Ce1−xLaxOy, the Raman spectra are progressively broadened and shifted to lower energy The spectrum of unmodified CeO 2 is characterized by one strong band at 465 cm −1 , which is due to the F2g vibration mode of the cubic fluorite structure [88,89]. In addition, it also exhibits two weak bands at about 250 and 600 cm −1 , which can be attributed to the nondegenerate transverse and longitudinal optical phonon modes of CeO 2 , respectively [90]. The Raman spectral features of the (CeM) 1−x Ni x O y samples were greatly affected by the content and type of dopant ( Figure 8). For these samples, Raman spectra also include main bands at 452-464 cm −1 and two weak bands at 230 cm −1 and 550-650 cm −1 , which can be attributed to the displacement of oxygen atoms from their ideal fluorite lattice positions [91] and the presence of defects such as oxygen vacancies [92]. This shows [93] that the emergence of the band at about 560 cm −1 (oxygen vacancies) results from the different oxidation state of the dopant compared to that of Ce 4+ , while the band at about 600 cm −1 is caused by the different ionic radius of the dopant compared to that of Ce 4+ . The I 570 /I 465 ratio, used to estimate the concentration of oxygen vacancies in a material [94], increases with increasing x, increasing M/Ce ratio, and in the series Mg→Al →La ( Table 2). The higher the I 570 /I 465 , the more oxygen vacancies in the material are formed. In addition, a shift (465 → 452 cm −1 ) of the position of the main band in Raman spectra for (CeM) 1−x Ni x O y samples compared to the position of this band for pure CeO 2 was noted. This indicates a change in the binding energies due to the incorporation of Ni 2+ and M n+ cations into the lattice of CeO 2 and confirms the formation of a solid solution. The considerable broadening of the bands observed for (CeM) 1−x Ni x O y ( Figure 8) could be attributed to the reduction in phonon lifetime due to the decrease in grain size [88], which correlates well with XRD data ( Figure 6). Significant deformation of the spectrum shape for La-containing samples at La/Ce ≥ 1 points to a substantial rearrangement of the oxide structure compared to that of unmodified ceria. It was demonstrated [95] that with the increase of La concentration in Ce 1−x La x O y , the Raman spectra are progressively broadened and shifted to lower energy (from 464 cm −1 (La = 0%) to 458 cm −1 (La = 20%)). These effects agree with the formation of homogeneous La-doped CeO 2 solid solutions.
The study of the nanostructure of the (CeM) 1−x Ni x O y by TEM shows that the samples are represented by particles of two types: (1) thin polycrystalline plates, from 100 nm to several microns in length and 10-20 nm in width; (2) small, irregularly shaped polycrystalline agglomerates up to 100 nm in size. According to the EDX data, the composition of all particles is close to the declared one. The size of primary crystallites is very small (Figure 9a The study of the nanostructure of the (CeM)1−xNixOy by TEM shows that the samples are represented by particles of two types: (1) thin polycrystalline plates, from 100 nm to several microns in length and 10-20 nm in width; (2) small, irregularly shaped polycrystalline agglomerates up to 100 nm in size. According to the EDX data, the composition of all particles is close to the declared one. The size of primary crystallites is very small (Figure 9a-d). For example, at x = 0.2 and M/Ce = 1 it is equal 1.5-2.5 nm for M = Al, 3-5 nm for M = La or Mg, and 5-8 nm for the sample without M. Analysis of the interplanar distances measured from the high-resolution images and microdiffraction patterns reveal the presence of a fluorite phase reflections corresponding to d111CeO2 (0.312 nm), d200CeO2 (0.271 nm) and d220CeO2 (1.913 nm). Also, in the case of the (CeMg)1−xNixOy sample, the reflection at 0.211 nm corresponding to d200 of NiO-MgO particles is clearly visible in the electron diffraction pattern. In the Ce1−xNixOy and (CeLa)1−xNixOy samples, only solitary highly dispersed NiO particles can be found on the surface of the samples with a typical interplanar distance of 0.209 nm corresponding to (d200NiO). In the case of M = Al, no reflections that would correspond to Ni-containing phases are visible. EDX mapping shows a uniform distribution of elements in the studied samples ( Figure 9, e-h) implying that even in the (CeMg)1−xNixOy sample, there is a very intense mixing of crystallites of the Ce-containing phase and Ni(Mg)O nanocrystallites, which, under imaging conditions at selected magnification, does not allow one to distinguish inhomogeneities of the local composition. Thus, solid solutions of various chemical compositions based on the cubic structure of ceria with a very-fine crystallite size were actually obtained. Both nickel and M cations were inserted into the structure of ceria through the synthesis of (CeM)1−xNixOy materials using the polymerizable complex method. This led to a significant decrease in the size of crystallites, an increase in the defectiveness of the material, and the formation of oxygen vacancies. Traces of highly dispersed Ni-containing species were also formed. This trend is intensified with an increase in the Ni content and with the use of Mg as a dopant. In the latter case, a MgO-NiO solid solution is formed. For the formation of catalytically active Ni 0 nanoparticles, the appropriate activation of (CeM)1−xNixOy materials is needed. Thus, solid solutions of various chemical compositions based on the cubic structure of ceria with a very-fine crystallite size were actually obtained. Both nickel and M cations were inserted into the structure of ceria through the synthesis of (CeM) 1−x Ni x O y materials using the polymerizable complex method. This led to a significant decrease in the size of crystallites, an increase in the defectiveness of the material, and the formation of oxygen vacancies. Traces of highly dispersed Ni-containing species were also formed. This trend is intensified with an increase in the Ni content and with the use of Mg as a dopant. In the latter case, a MgO-NiO solid solution is formed. For the formation of catalytically active Ni 0 nanoparticles, the appropriate activation of (CeM) 1−x Ni x O y materials is needed.

Activation of Samples
As a rule, to activate nickel reforming catalysts, samples were reduced at a temperature of 600-800 • C for 1-3 h [37,96,97]. In order to study the features of the reduction of (CeM) 1−x Ni x O y samples, methods of temperature-programed hydrogen reduction (Figures 10 and 11) and in situ XRD (Figures 12 and 13) were applied. On the TPR-H 2 curves of unmodified ceria, two regions of hydrogen consumption can be distinguished: in the temperature ranges of 100-600 • C (a peak maximum at T = 500 • C) and 600-1000 • C (a peak maximum at T = 880 • C), which are due to the reduction of cations Ce 4+ localized on the surface and in the volume of particles, respectively [52,98]. The reduction profile of (CeM) 1−x Ni x O y samples depends upon the amount of introduced nickel ( Figure 10a) as well as the type and content of M ( Figure 10b).

Activation of Samples
As a rule, to activate nickel reforming catalysts, samples were reduced at a temperature of 600-800 °C for 1-3 h [37,96,97]. In order to study the features of the reduction of (CeM)1−xNixOy samples, methods of temperature-programed hydrogen reduction ( Figures  10 and 11) and in situ XRD (Figures 12 and 13) were applied. On the TPR-H2 curves of unmodified ceria, two regions of hydrogen consumption can be distinguished: in the temperature ranges of 100-600 °C (a peak maximum at T = 500 °C) and 600-1000 °C (a peak maximum at T = 880 °C), which are due to the reduction of cations Ce 4+ localized on the surface and in the volume of particles, respectively [52,98]. The reduction profile of (CeM)1−xNixOy samples depends upon the amount of introduced nickel ( Figure 10a) as well as the type and content of M ( Figure 10b). For Ce1−xNixOy-series, three regions of hydrogen consumption can be observed: lowtemperature (T < 250 °C), medium (250 < T < 500 °C), and high-temperature (T > 550 °C). In the low-temperature region, weak consumption of H2 is observed, the intensity and position of which is practically independent on the Ni content. The two peaks with maxima at ~150 °C and 230 ± 10 °C can be identified. In the medium-temperature region, the intensity of H2 consumption increases with an increase in the nickel content. The peak is wide, without clearly defined maxima; its position slightly shifts towards higher temperatures with increasing x. In the high-temperature region, there is one peak, the maximum of which shifts from 880 to 780 °C as x increases from 0 to 0.8.
As shown by XRD, Raman spectroscopy, and TEM methods (Figures 6-9, Table 2), Ni 2+ cations in Ce1−xNixOy are stabilized in a CeO2-based solid solution and fine NiO particles of ~5 nm in size. The share of the latter increases with the growth of x and becomes significant only at x ≥ 0.7. For comparison, Figure 10 also shows the TPR data for samples where Ni 2+ ions are stabilized only in the form of well-crystallized NiO particles with an average size of ~50 nm. The first sample is a sample with x = 0.3 obtained by the impregnation method and calcined at 500 °C; the second is a sample with x = 0.2 obtained by the polymerizable complex method and calcined at 900 °C. The reduction profiles of these samples are similar and are described by three hydrogen consumption regions with maxima at 230, 350, and 830 °C. The fraction of [Ni-Ce-O] structures in such non-dispersed samples is insignificant. Note also that, in the impregnated samples, with an increase in the Ni content, the intensity of H2 consumption increases only in the medium-temperature region [30].  Figure 10b). Additionally, for the Al-containing system, a peak at 2Ɵ = 26.4° was observed during reduction, which can be attributed to the oxygen-deficient CeOy phase. Its low intensity indicates a small amount of this phase.    The evolution of the unit cell volume of a CeO2-based solid solution during reduction exhibits a complex pattern ( Figure 13). With thermal treatment, thermal expansion of the lattice is expected, while a sudden change in the slope may be associated with a chemical process [36,70,81,105]. Indeed, for the (CeM)1−xNixOy samples, there is a region where the expected growth of the unit cell volume due to thermal expansion does not occur. Note that this behavior is typical only for Ni-containing samples. For nickel-free oxides (dot lines in Figure 13), a continuous increase in the unit cell volume with increasing temperature is observed. The studied temperature range (up to 700 °C) is insufficient for the reduction of Ce 4+ cations in the volume of crystallites. Therefore, the decrease in the lattice expansion rate may be associated with a change in the lattice composition due to Ni 2+ exsolution from CeO2 and the complicated activities of vacancies formed according to reaction (12) during the reduction process. The mobility of oxygen vacancies, their dimerization, clustering, or migration from the lattice to the surface to fill low-energy spaces, affects the structural parameters of samples [106].
With the exception of the Al-containing sample, the size of crystallites of CeO2-based solid solution increases during the thermal treatment in H2 ( Figure 13). The degree of crystallite growth is higher for the M-free sample (6.2 → 21.3 nm) and increases in the series M La (4.4 → 7.8 nm) < Mg (3.6 → 11.2 nm). For an Al-containing sample, the size of solid solution crystallites remains at the level of 2-3 nm. Accordingly, the introduction of M into the composition of the samples effectively prevents sintering. This effect increases in the series Mg→La→Al. The XRD data for (CeM)1−xNixOy samples after activation at 800 °C for 1 h in a 30%H2/70%Ar flow is presented in Table 3 and Figure 14. For Ce 1−x Ni x O y -series, three regions of hydrogen consumption can be observed: lowtemperature (T < 250 • C), medium (250 < T < 500 • C), and high-temperature (T > 550 • C). In the low-temperature region, weak consumption of H 2 is observed, the intensity and position of which is practically independent on the Ni content. The two peaks with maxima at 150 • C and 230 ± 10 • C can be identified. In the medium-temperature region, the intensity of H 2 consumption increases with an increase in the nickel content. The peak is wide, without clearly defined maxima; its position slightly shifts towards higher temperatures with increasing x. In the high-temperature region, there is one peak, the maximum of which shifts from 880 to 780 • C as x increases from 0 to 0.8.
As shown by XRD, Raman spectroscopy, and TEM methods (Figures 6-9, Table 2), Ni 2+ cations in Ce 1−x Ni x O y are stabilized in a CeO 2 -based solid solution and fine NiO particles of~5 nm in size. The share of the latter increases with the growth of x and becomes significant only at x ≥ 0.7. For comparison, Figure 10 also shows the TPR data for samples where Ni 2+ ions are stabilized only in the form of well-crystallized NiO particles with an average size of~50 nm. The first sample is a sample with x = 0.3 obtained by the impregnation method and calcined at 500 • C; the second is a sample with x = 0.2 obtained by the polymerizable complex method and calcined at 900 • C. The reduction profiles of these samples are similar and are described by three hydrogen consumption regions with maxima at 230, 350, and 830 • C. The fraction of [Ni-Ce-O] structures in such non-dispersed samples is insignificant. Note also that, in the impregnated samples, with an increase in the Ni content, the intensity of H 2 consumption increases only in the medium-temperature region [30].
Based on the analysis of literature data [12,[98][99][100][101] and the results of studying samples by structural methods, the H 2 consumption in the low-temperature region indicates the reduction of [Ni-O-Ce] structures localized on the surface of the solid solution or at the interface between fine NiO particles and CeO 2 crystallites: where V is an oxygen vacancy. The H 2 consumption in the medium-temperature region indicates the reduction of Ni 2+ cations in the bulk of CeO 2 -based solid solution as well as NiO particles. The third region of hydrogen uptake is associated with the reduction of Ce 4+ cations localized in the volume of particles. This process becomes easier as x increases. Samples of the Ce 1−x Ni x O y -series are characterized by close values of the specific surface area (Table 1) and, as a consequence, a comparable content of easily reduced surface [Ni-O-Ce] structures in their composition, regardless of the different nickel contents. This type of structure is the main one at x ≤ 0.1. On the contrary, the fraction of the second type of Ni 2+ increases with increasing x: for x < 0.7 it is mainly [Ni-O-Ce] structures in the bulk of CeO 2 -based solid solution, while at x ≥ 0.7 the contribution of NiO particles grows. In any case, for M-free samples with a wide range of x = 0.01-0.8, the reduction of Ni 2+ ends at a temperature not higher than 500 • C.
The introduction of M into the composition of the samples greatly changes their reducing properties (Figure 10b). For (CeM) 1−x Ni x O y with M = Al or La, the profile of the TPR curve is similar to that of the sample without M, but it is shifted to the high-temperature region (∆T = 100-200 • C). Considering the structural data (Table 2) (Table 2), it can be assumed that in the low-temperature region, surface and bulk [Ni-O-Ce] structures are reduced, while in the high-temperature region, NiO-MgO solid solution undergoes reduction [102]. Therefore, the feature of (CeM) 1−x Ni x O y (M = Al, La, and Mg) series is that a significant part of the Ni 2+ is reduced at T = 500-700 • C. This trend increases with increasing M/Ce molar ratio and for Mg-containing samples. An increase in the Ni 2+ reduction temperature for (CeM) 1−x Ni x O y indicates the increasing strength of the metal-support interaction [23]. For (CeM) 1−x Ni x O y (M = Al, La) samples with a high M/Ce molar ratio, the idea that the formation of joint highly dispersed surface Al-Ni-O or La-Ni-O structures was reduced in the high-temperature region cannot be ruled out [60,103].
Total hydrogen consumption for (CeM) 1−x Ni x O y samples during TPR increases with increasing x, which is associated with an increase in the Ni content, and in the series Mg→Al→La, which indicates an increase in the reducibility of samples ( Figure 11). The lower hydrogen uptake of the Mg-containing system may be due to the presence of a hard-to-reduce NiO-MgO solid solution. Figure 12 shows in situ XRD data for the reduction of (CeM) 1−x Ni x O y samples in a 30% H 2 /70% He mixture as a function of temperature. When the samples are reduced, the CeO 2 -based phase is retained, and the formation of metallic Ni is observed. For M-free samples, traces of Ni 0 phase are observed at 500 • C (Figure 12a) which corresponds to the temperature of full reduction of Ni 2+ cations according TPR data (Figure 10a). This correlates with the data that Ni exsolution, atomic diffusion of Ni toward the surface, and the nucleation of Ni particles start and proceed at lower temperatures than anticipated from their appearance in the crystallized phase in XRD [104]. For (CeM) 0.8 Ni 0.2 O y (M = Al, La, Mg) the temperature of the formation of the metallic nickel phase depends upon the type of M. The temperature at which reflections of metallic nickel in the diffraction pattern clearly appear, increases from 500 to 650 • C in the series Mg→Al→La ( Figure 12). For Al and La-containing samples at this temperature (T > 600 • C), the reduction of Ni 2+ cations is practically completed (Figure 10b Figure 10b). Additionally, for the Al-containing system, a peak at 2Ɵ = 26.4° was observed during reduction, which can be attributed to the oxygen-deficient CeOy phase. Its low intensity indicates a small amount of this phase. Ɵ = 26.4 • was observed during reduction, which can be attributed to the oxygen-deficient CeO y phase. Its low intensity indicates a small amount of this phase.
The evolution of the unit cell volume of a CeO 2 -based solid solution during reduction exhibits a complex pattern ( Figure 13). With thermal treatment, thermal expansion of the lattice is expected, while a sudden change in the slope may be associated with a chemical process [36,70,81,105]. Indeed, for the (CeM) 1−x Ni x O y samples, there is a region where the expected growth of the unit cell volume due to thermal expansion does not occur. Note that this behavior is typical only for Ni-containing samples. For nickel-free oxides (dot lines in Figure 13), a continuous increase in the unit cell volume with increasing temperature is observed. The studied temperature range (up to 700 • C) is insufficient for the reduction of Ce 4+ cations in the volume of crystallites. Therefore, the decrease in the lattice expansion rate may be associated with a change in the lattice composition due to Ni 2+ exsolution from CeO 2 and the complicated activities of vacancies formed according to reaction (12) during the reduction process. The mobility of oxygen vacancies, their dimerization, clustering, or migration from the lattice to the surface to fill low-energy spaces, affects the structural parameters of samples [106].
With the exception of the Al-containing sample, the size of crystallites of CeO 2 -based solid solution increases during the thermal treatment in H 2 ( Figure 13). The degree of crystallite growth is higher for the M-free sample (6.2 → 21.3 nm) and increases in the series M La (4.4 → 7.8 nm) < Mg (3.6 → 11.2 nm). For an Al-containing sample, the size of solid solution crystallites remains at the level of 2-3 nm. Accordingly, the introduction of M into the composition of the samples effectively prevents sintering. This effect increases in the series Mg→La→Al.
The XRD data for (CeM) 1−x Ni x O y samples after activation at 800 • C for 1 h in a 30% H 2 /70% Ar flow is presented in Table 3 and Figure 14. It can be seen that the activated samples of the Ce 1−x Ni x O y -series contain a CeO 2 -based phase and the phase of metallic nickel. For the CeO 2 -based phase, an increase in the unit cell parameter (∆a is equal to 0.0006 → 0.0019 at x = 0.1 → 0.8) is observed, which may indicate the release of nickel cations from the structure as well as an increase in the Ce 3+ /Ce 4+ molar ratio. The average size of ceria crystallites amounts to 25-50 nm, which is significantly higher than in fresh samples. The average size of the formed Ni 0 crystallites depends on the nickel content and increases from 8.5 to 50 nm with increasing x from 0.1 to 0.8. present in X-ray amorphous forms. For M=Mg, the changes resulting from activation are similar to those found for the M-free series. A distinguishing feature of this series is the presence of the magnesium oxide phase with an average crystallite size of 12 nm. Note that the lattice parameter of this phase indicates that the selected activation conditions ensure the reduction of the NiO-MgO solid solution present in the fresh sample. The absence of the Ni 0 -phase reflections in the XRD patterns of the samples with M/Ce >1 (M = Al or La) after activation indicates its high dispersion.  Nonetheless, the specific surface area of (CeM)1−xNixOy samples decreases as a result of high-temperature treatment during activation. To a lesser extent, SBET declines in samples with an initially lower surface. For example, at M/Ce = 1, SBET decreases by factors of 7.2 (80 → 11 m 2 /g), 2.8 (48 → 17 m 2 /g), and 2.4 (36 → 15 m 2 /g) for M = Al, La and M = Mg, respectively. Differences in the textural properties of the samples become less noticeable. Generally, all values of the specific surface area of activated (CeM)0.8Ni0.2Oy samples are within the range of 20 ± 10 m 2 /g (Table 3). Only for x ≥ 0.7 SBET decreases to ~5 m 2 /g.   After activation, the samples retained a spongy morphology ( Figure 16) and mesoporosity. The average diameter of pores is within the wide range of 11-55 nm (Table 3). Nonetheless, the specific surface area of (CeM) 1−x Ni x O y samples decreases as a result of high-temperature treatment during activation. To a lesser extent, S BET declines in samples with an initially lower surface. For example, at M/Ce = 1, S BET decreases by factors of 7.2 (80 → 11 m 2 /g), 2.8 (48 → 17 m 2 /g), and 2.4 (36 → 15 m 2 /g) for M = Al, La and M = Mg, respectively. Differences in the textural properties of the samples become less noticeable. Generally, all values of the specific surface area of activated (CeM) 0.8 Ni 0.2 O y samples are within the range of 20 ± 10 m 2 /g ( Table 3). Only for x ≥ 0.7 S BET decreases to~5 m 2 /g. Thus, in the (CeM) 1−x Ni x O y samples, as a result of activation, exsolution from fluorite structure and reduction of Ni 2+ cations occur, and Ni 0 particles are formed. The temperature range of Ni 0 phase formation is higher for samples with a higher Ni content and increases with an increase in the molar ratio M/Ce and in the series M-free < Al = La < Mg. This indicates an increase in the Ni-support interaction upon the introduction of M into the sample. In the case of M = Al or La, the oxide matrix is the Ce-M-O joint phase, which is in contact with Ni 0 nanoparticles. The situation is opposite in the case of M = Mg: there are coexisting CeO 2 , MgO, and Ni 0 phases. In situ XRD study in the temperature range up to 700 • C shows that the resistance to sintering of the CeO 2 -based phase is higher in the case of doping M = Al or La, i.e., when a CeO 2 -based solid solution continues to exist after activation. After activation at 800 • C, the dispersion of the Ni 0 phase at the same content is comparable at M/Ce ≤ 1 for all studied M. There are two types of species: nanocrystallites (2-4 nm) and nanoparticles with an average size of 12-15 nm (x = 0.2). This is a fairly low nickel particle size at its content of 10 wt.%. For comparison, at the same Ni content and temperature of reduction, the Ni 0 crystallite size varies from 19 to 33 nm for impregnated 10Ni/CeO 2 samples [71,105] and from 20 to 46 nm for 10Ni/CeO 2 , prepared using the hydrothermal method [36]. As x increases from 0.1 to 0.8, the average particle size increases from 8 to 50 nm. At higher M/Ce molar ratio, the influence M becomes stronger. At M/Ce > 1, in contrast to Mg, the presence of Al or La in the composition of (CeM) 1−x Ni x O y samples ensures a higher dispersion of the Ni 0 phase as a result of activation. Thus, the activated (CeM) 1−x Ni x O y samples have similar textural and morphological properties but different structural characteristics and resistance to sintering due to the different strength of Ni-support interaction, which can affect their functional performance.

Functional Properties
The functional properties (activity, stability, and anti-coking ability) of the (CeM) 1−x Ni x O y samples were studied in the methane bi-reforming reaction when tested in screening and stability modes. Figure 17 shows the temperature dependences of product concentrations during the bi-reforming of methane over (CeM) 1−x Ni x O y samples. In the case of x = 0.05 (Figure 17a), the CH 4 and CO 2 conversion starts at 600 • C, but the quantity of reforming products-CO and H 2 -is negligible (<10 vol.%). As the temperature rises, concentrations of methane and carbon dioxide decrease, while those of CO and hydrogen increase. This tendency is more pronounced for samples with x ≥ 0.2 (Figure 17b-e), for which higher values of H 2 and CO concentrations are achieved in the whole temperature range 600-850 • C. In the case of M = Al or Mg, the process parameters at 850 • C are close to thermodynamic equilibrium results (Figure 17f).
increase in x. This may be due to an increase in the average particle size of the Ni active component (8 → 50 nm, Table 3), and, accordingly, a decrease in the fraction of available surface-active sites. The surface atoms/bulk atoms ratio rapidly decreases with increasing particle size: at a particle size of 3 nm, 50% of atoms or ions are on the surface, while at a size of 10 nm it is ~15% and at a size of 25 nm it is ~5% [107]. It can be concluded that at x = 0.2, a compromise is reached between sufficiently high catalyst productivity and a notquite-low specific rate of hydrogen formation.  The parameters of the bi-reforming of methane over (CeM) 1−x Ni x O y samples depend on the values of x ( Figure 18) and type of M ( Figure 19). It was found that with an increase in x from 0.05 to 0.8, both the conversion of reactants and the yield of products increased, reaching the following values at 750 • C: X(CH 4 ) 32 → 86%, X(CO 2 ) 46 → 73%, Y(H 2 ) 45 → 92%, and Y(CO) 43 → 85%. Catalyst productivity also grew with increasing x: from 8.1 to 12.0 L(H 2 )/(g cat ·h) which may be connected with a higher content of Ni and, accordingly, a higher content of active Ni species per gram of catalyst. However, the efficiency of H 2 formation per mole of Ni decreases from 17.7 to 1.1 mole H2 /(mole Ni ·min) (Figure 18b). It can be seen that, in contrast to the conversion and yield indicators, the specific indicator (calculated per mole of the Ni active component) of the catalyst activity decreases with an increase in x. This may be due to an increase in the average particle size of the Ni active component (8 → 50 nm, Table 3), and, accordingly, a decrease in the fraction of available surface-active sites. The surface atoms/bulk atoms ratio rapidly decreases with increasing particle size: at a particle size of 3 nm, 50% of atoms or ions are on the surface, while at a size of 10 nm it is~15% and at a size of 25 nm it is~5% [107]. It can be concluded that at x = 0.2, a compromise is reached between sufficiently high catalyst productivity and a not-quite-low specific rate of hydrogen formation.   The performance of (CeM) 0.8 Ni 0.2 O y in the bi-reforming of methane exceeds that of Ce 0.8 Ni 0.2 O 1.8 ( Figure 19). An exception is the La-containing sample, in the presence of which, despite the high values of the conversion of the reagents, a reduced yield of hydrogen is observed. This may indicate the accumulation of carbonaceous deposits according to the reactions (8-9). Indeed, after screening tests (stepwise temperature-rise mode, 650→850 • C) the samples contain coke, the content of which increases in the series 0.8 wt.% (M = Al) ≈ 0.9 wt.% (M = Mg) < 2.7 wt.% (M-free) < 6.9 wt.% (M = La). At constant type of M, changing the M/Ce ratio has little effect on catalyst performance in screening tests. There is a slight tendency to a decrease in the hydrogen yield with an increase in the M/Ce molar ratio.
Note that the activity in screening tests of the studied (CeM) 0.8 Ni 0.2 O y samples, prepared using the PC method, is comparable with that of impregnated samples of the same composition [108], while the degree of coking is significantly lower (Figure 20). For both cases, the Mg-containing samples are the most resistant to coking. Table 4 and Figure 21 demonstrate the effect of Ni molar fraction (x), type of M, molar ratio M/Ce, and preparation method on the parameters of the bi-reforming of methane at 800 • C during stability tests. A decrease in process performance over time is observed when M-free samples are used as catalysts or at M = La. The degree of sample deactivation, calculated as the relative decrease in H 2 yield over 24 h of time on stream, is at the maximum (35%) for a sample with a low Ni molar fraction (x = 0.05). For the La-series, it does not exceed 10% (Table 4). On the contrary, in the case of M = Al or Mg, the process values practically do not change with time on stream. As in the screening tests, in the stability tests, the M/Ce molar ratio has little effect on the activity of the samples. Among the studied types of M, the yield of hydrogen increases in the series M-free < La < Mg < Al ( Figure 21). Better performance of methane reforming is observed for samples obtained by the PC method rather than by impregnation (Figure 21d). The performance of (CeM)0.8Ni0.2Oy in the bi-reforming of methane exceeds that of Ce0.8Ni0.2O1.8 ( Figure 19). An exception is the La-containing sample, in the presence of which, despite the high values of the conversion of the reagents, a reduced yield of hydrogen is observed. This may indicate the accumulation of carbonaceous deposits according to the reactions (8)(9). Indeed, after screening tests (stepwise temperature-rise mode, 650→850 °C) the samples contain coke, the content of which increases in the series 0.8 wt.% (M = Al) ≈ 0.9 wt.% (M = Mg) < 2.7 wt.% (M-free) < 6.9 wt.% (M = La). At constant type of M, changing the M/Ce ratio has little effect on catalyst performance in screening tests. There is a slight tendency to a decrease in the hydrogen yield with an increase in the M/Ce molar ratio.
Note that the activity in screening tests of the studied (CeM)0.8Ni0.2Oy samples, prepared using the PC method, is comparable with that of impregnated samples of the same composition [108], while the degree of coking is significantly lower (Figure 20). For both cases, the Mg-containing samples are the most resistant to coking. The performance of (CeM)0.8Ni0.2Oy in the bi-reforming of methane exceeds that of Ce0.8Ni0.2O1.8 ( Figure 19). An exception is the La-containing sample, in the presence of which, despite the high values of the conversion of the reagents, a reduced yield of hydrogen is observed. This may indicate the accumulation of carbonaceous deposits according to the reactions (8)(9). Indeed, after screening tests (stepwise temperature-rise mode, 650→850 °C) the samples contain coke, the content of which increases in the series 0. Note that the activity in screening tests of the studied (CeM)0.8Ni0.2Oy samples, prepared using the PC method, is comparable with that of impregnated samples of the same composition [108], while the degree of coking is significantly lower (Figure 20). The performance of (CeM)0.8Ni0.2Oy in the bi-reforming of methane exceeds that of Ce0.8Ni0.2O1.8 ( Figure 19). An exception is the La-containing sample, in the presence of which, despite the high values of the conversion of the reagents, a reduced yield of hydrogen is observed. This may indicate the accumulation of carbonaceous deposits according to the reactions (8)(9). Indeed, after screening tests (stepwise temperature-rise mode, 650→850 °C) the samples contain coke, the content of which increases in the series 0. Note that the activity in screening tests of the studied (CeM)0.8Ni0.2Oy samples, prepared using the PC method, is comparable with that of impregnated samples of the same composition [108], while the degree of coking is significantly lower (Figure 20). The performance of (CeM)0.8Ni0.2Oy in the bi-reforming of methane exceeds that of Ce0.8Ni0.2O1.8 ( Figure 19). An exception is the La-containing sample, in the presence of which, despite the high values of the conversion of the reagents, a reduced yield of hydrogen is observed. This may indicate the accumulation of carbonaceous deposits according to the reactions (8)(9). Indeed, after screening tests (stepwise temperature-rise mode, 650→850 °C) the samples contain coke, the content of which increases in the series 0. Note that the activity in screening tests of the studied (CeM)0.8Ni0.2Oy samples, prepared using the PC method, is comparable with that of impregnated samples of the same composition [108], while the degree of coking is significantly lower (Figure 20). The performance of (CeM)0.8Ni0.2Oy in the bi-reforming of methane exceeds that of Ce0.8Ni0.2O1.8 ( Figure 19). An exception is the La-containing sample, in the presence of which, despite the high values of the conversion of the reagents, a reduced yield of hydrogen is observed. This may indicate the accumulation of carbonaceous deposits according to the reactions (8)(9). Indeed, after screening tests (stepwise temperature-rise mode, 650→850 °C) the samples contain coke, the content of which increases in the series 0. Note that the activity in screening tests of the studied (CeM)0.8Ni0.2Oy samples, prepared using the PC method, is comparable with that of impregnated samples of the same composition [108], while the degree of coking is significantly lower (Figure 20).  Studied samples of optimal composition provide high and stable conversion of methane (X CH4 ≈ 90%) and CO 2 (X CO2 ≈ 85%), as well as a high hydrogen yield (Y H2 > 90%) that is comparable to or better than the literature data for the bi-reforming of methane [109][110][111][112][113][114]. Namely, during the bi-reforming of methane at 800 • C over impregnated Ni/CeO 2 catalyst, X CH4 is equal to 50% [109], at 800 • C over co-precipitated Ni/CeO 2 it is 80% [109], at 900 • C over Ni-Ce-Fe/Al 2 O 3 it is 90% [110], at 800 • C over Ni/SBA-15 it is 70% [111], at 800 • C over Ni/MgAl 2 O 4 it is 90% [112], at 800 • C over Ni/Ce 0.6 Zr 0.4 O 2 it is 60% [113], and at 700 • C over Ni/ZrO 2 it is 45-70%, depending on the type of ZrO 2 [114].
The study of spent catalysts by thermal analysis shows that an insignificant weight loss (0.2-0.4 wt.%) occurs in the low-temperature region (T < 200 • C) due to the desorption of water and volatile intermediate products ( Figure 22). Further, at a temperature of 300-500 • C, the weight of the sample increases, which is associated with the oxidation of the Ni 0 active component. The process is accompanied by an exothermic effect, the maximum of which shifts from 345 to 420 • C in the series Al < La < Mg < M-free. This indicates the decrease in Ni stability to re-oxidation in the presence of M, which correlates with improvement to the metal-support interaction. This effect is more pronounced for M = Al. There is weight loss in the region of 500-800 • C. The exception is M-free samples with x < 0.3, for which no change in weight is observed in the high-temperature region. The amounts of this weight loss vary from 0.2 to 9 wt.% and can be assigned to the oxidation of carbonaceous deposits. Higher weight values due to the burnout of carbon deposits are observed for samples with M = La. The carbon accumulation rate decreases with the decrease of the molar ratio of M/Ce and in the series La→Al→Mg. Among (CeMg) 0.8 Ni 0.2 O y samples, the minimum rate of formation of carbonaceous deposits was observed for the Ce 0.6 Mg 0.2 Ni 0.2 O 1.6 sample ( Table 4). It is equal to 0.08 mgC/g cat ·h, which is significantly lower than the specific rate of H 2 formation: 6.8 × 10 −5 mole C /(mole Ni ·min) vs. 5.0 mole H2 /(mole Ni ·min).  The study of spent catalysts by thermal analysis shows that an insignificant weight loss (0.2-0.4 wt.%) occurs in the low-temperature region (T < 200 °C) due to the desorption of water and volatile intermediate products ( Figure 22). Further, at a temperature of 300-500 °C, the weight of the sample increases, which is associated with the oxidation of the Ni 0 active component. The process is accompanied by an exothermic effect, the maximum of which shifts from 345 to 420 °C in the series Al < La < Mg < M-free. This indicates the decrease in Ni stability to re-oxidation in the presence of M, which correlates with improvement to the metal-support interaction. This effect is more pronounced for M = Al. There is weight loss in the region of 500-800 °C. The exception is M-free samples with x < 0.3, for which no change in weight is observed in the high-temperature region. The amounts of this weight loss vary from 0.2 to 9 wt.% and can be assigned to the oxidation of carbonaceous deposits. Higher weight values due to the burnout of carbon deposits are observed for samples with M = La. The carbon accumulation rate decreases with the decrease of the molar ratio of M/Ce and in the series La→Al→Mg. Among (CeMg)0.8Ni0.2Oy samples, the minimum rate of formation of carbonaceous deposits was observed for the Ce0.6Mg0.2Ni0.2O1.6 sample (Table 4). It is equal to 0.08 mgC/gcat•h, which is significantly lower than the specific rate of H2 formation: 6.8 × 10 −5 moleC/(moleNi•min) vs. 5.0 moleH2/(moleNi•min). The main pathways of carbon formation are cracking (6) and Boudouard reaction (7). According to thermodynamic data, the selectivity of coke formation during the bi-reforming of methane decreases with increasing temperature [21]. Therefore, carbon formation occurs more intensively in screening tests than in stability tests: 6.9 vs. 3.2 and 0.9 vs. 0.4 wt.% for M = La and Mg and at M/Ce = 1. In comparison to the literature data, the carbon formation level for spent samples with M = Al or Mg is low. For example, after the bi-reforming of methane, 12% Ni/Al 2 O 3 catalyst includes 50 wt.% of carbonaceous deposits (C) [110], 12% Ni-5% Ce-5% Fe/Al 2 O 3-2 wt.% of C [110], 5% Ni/MgAl 2 O 4-3 wt.% of C [112], 10% Ni/SBA-15-6.4 and 3% B-10% Ni/SBA-15-1.5 wt.% of C [111], for 10% Ni/ZrO 2 the content of C varies from 0.2 to 47.9 wt.% for different types of ZrO 2 support [114]. Therefore, a significant advantage of the developed materials is their high resistance to coking even when a low steam/CO 2 molar ratio (H 2 O/CO 2 = 0.5) is used in the initial reaction mixture. This stability of the catalyst to coking may be due to several factors. First, this is connected with the anti-coking properties of ceria support [49,115], and second, the large size of Ni crystallites, on which, as was shown in [36], the coking rate is lower than for nickel particles with a size of 20-30 nm. In addition to the influence of M on the rate of coke formation, its influence on the mode of change in the textural and structural properties of the samples (CeM) 1−x Ni x O y as a result of the reaction is expected.
provement to the metal-support interaction. This effect is more pronounced for M = Al. There is weight loss in the region of 500-800 °C. The exception is M-free samples with x < 0.3, for which no change in weight is observed in the high-temperature region. The amounts of this weight loss vary from 0.2 to 9 wt.% and can be assigned to the oxidation of carbonaceous deposits. Higher weight values due to the burnout of carbon deposits are observed for samples with M = La. The carbon accumulation rate decreases with the decrease of the molar ratio of M/Ce and in the series La→Al→Mg. Among (CeMg)0.8Ni0.2Oy samples, the minimum rate of formation of carbonaceous deposits was observed for the Ce0.6Mg0.2Ni0.2O1.6 sample (Table 4). It is equal to 0.08 mgC/gcat•h, which is significantly lower than the specific rate of H2 formation: 6.8 × 10 −5 moleC/(moleNi•min) vs. 5.0 moleH2/(moleNi•min). The main pathways of carbon formation are cracking (6) and Boudouard reaction (7). According to thermodynamic data, the selectivity of coke formation during the bi-reforming of methane decreases with increasing temperature [21]. Therefore, carbon formation occurs more intensively in screening tests than in stability tests: 6.9 vs. 3.2 and 0.9 vs. 0.4 wt.% for M = La and Mg and at M/Ce = 1. In comparison to the literature data, the carbon formation level for spent samples with M = Al or Mg is low. For example, after the bireforming of methane, 12% Ni/Al2O3 catalyst includes 50 wt.% of carbonaceous deposits (C) [110], 12% Ni-5% Ce-5% Fe/Al2O3-2 wt.% of C [110], 5% Ni/MgAl2O4-3 wt.% of C [112], 10% Ni/SBA-15-6.4 and 3% B-10% Ni/SBA-15-1.5 wt.% of C [111], for 10% Ni/ZrO2 the content of C varies from 0.2 to 47.9 wt.% for different types of ZrO2 support [114]. Therefore, a significant advantage of the developed materials is their high resistance to coking even when a low steam/CO2 molar ratio (H2O/CO2 = 0.5) is used in the initial reaction mixture. This stability of the catalyst to coking may be due to several factors. First, this is connected with the anti-coking properties of ceria support [49,115], and second, the large size of Ni crystallites, on which, as was shown in [36], the coking rate is lower than for nickel particles with a size of 20-30 nm. In addition to the influence of M on the rate of coke formation, its influence on the mode of change in the textural and structural properties of the samples (CeM)1−xNixOy as a result of the reaction is expected.

Characterization of Spent Samples
The harsh conditions of the catalytic reaction (high temperatures, presence of water vapor, etc.) inevitably affect the properties of the samples. Table 5 presents the textural properties and phase composition of spent (CeM)1−xNixOy samples.

Characterization of Spent Samples
The harsh conditions of the catalytic reaction (high temperatures, presence of water vapor, etc.) inevitably affect the properties of the samples. Table 5 presents the textural properties and phase composition of spent (CeM) 1−x Ni x O y samples.
The specific surface area of Ce 1−x Ni x O y samples decreases from 20 to 5 m 2 /g, while for the (CeM) 1−x Ni x O y samples it remains at the level of the values for the samples after activation (15 ± 5 m 2 /g). This indicates a positive effect of M on the textural properties of the samples and a lower rate of thermal sintering of (CeM) 1−x Ni x O y samples. Spent samples lose the porous spongy structure that was observed for fresh ( Figure 5) and activated ( Figure 16) samples and consist of large aggregates of irregularly shaped particles ( Figure 23). The morphology of the spent M-free sample looked particularly dense. The average pore diameter increases, reaching 20-50 nm in size (Table 5). XRD data show that spent Ce 1−x Ni x O y samples contain CeO 2 and Ni 0 phases ( Figure 24, Table 5), which is similar to the phase composition observed after sample activation ( Figure 14, Table 3). However, the phases become less dispersed (50 nm vs. 15-25 nm) as a result of sintering under the reaction conditions, which is typical for supported metal catalysts [31,116]. A similar picture is observed for the (CeM) 1−x Ni x O y samples. A distinctive feature of the La-containing series is the preservation of a CeO 2 -based solid solution after the reaction. Its average crystallite size increases from 20-25 to 50 nm. In contrast to M-free, Al-and La-containing samples, at M = Mg, the sample composition includes three phases (CeO 2 , MgO, and Ni 0 ) formed at the activation stage (Tables 3 and 5). They grow larger as a result of the reaction. Nevertheless, the ceria phase has an average crystallite size of 25 nm, which is smaller than those for M-free and La-containing samples. XRD data show that spent Ce1−xNixOy samples contain CeO2 and Ni 0 phases ( Figure  24, Table 5), which is similar to the phase composition observed after sample activation ( Figure 14, Table 3). However, the phases become less dispersed (50 nm vs. 15-25 nm) as a result of sintering under the reaction conditions, which is typical for supported metal catalysts [31,116]. A similar picture is observed for the (CeM)1−xNixOy samples. A distinctive feature of the La-containing series is the preservation of a CeO2-based solid solution after the reaction. Its average crystallite size increases from 20-25 to 50 nm. In contrast to M-free, Al-and La-containing samples, at M = Mg, the sample composition includes three phases (CeO2, MgO, and Ni 0 ) formed at the activation stage (Tables 3 and 5). They grow larger as a result of the reaction. Nevertheless, the ceria phase has an average crystallite size of 25 nm, which is smaller than those for M-free and La-containing samples.
According HRTEM study, for all samples, in addition to large, well-crystallized big Ni particles, Ni nanoparticles of 1.5-5 nm in size were observed on the oxide surface. The particles were stabilized due to epitaxy on CeO2 (a coincidence of the crystallographic directions [111] CeO2 and [100] Ni is observed) (Figure 25). EDX-mapping of spent samples demonstrates that distribution behavior of M between phases of an individual oxide and CeO2-based solid solution is controlled by the type of M. After 24 stability tests, La 3+ cations remain in the solid solution, Al 3+ cations only partially remain in the fluorite structure and form Al-enriched crystallites of needle shape, while Mg 2+ cations prefer to exist outside of the CeO2-based solid solution.  Thus, the high activation temperature and subsequent harsh reaction conditions smoothed out the initial textural and structural differences between samples with different values of x. After 24 h during the bi-reforming of methane, spent Ce1−xNixOy samples were characterized by the same very-low SBET values (3-5 m 2 /g) and the phase composition, which is CeO2 and Ni 0 with a CSR of 50 nm. Nevertheless, with increasing x, first, a higher hydrogen productivity is achieved, and hence a higher hydrogen concentration in the mixture of reaction products, and second, lower deactivation rates are observed. However, the number of carbonaceous deposits is increasing: samples with a lower degree of deactivation (x = 0.3) have a higher rate of coke accumulation. It can be assumed that these differences in the activity and stability of the Ce1−xNixOy samples are more related to the differences in their anti-sintering properties, rather than in their resistance to coke formation. In situ XAS spectroscopy indicates [117] that under reducing conditions the size and morphology of Ni particles change, as they become flattened and strongly stabilized  Thus, the high activation temperature and subsequent harsh reaction conditions smoothed out the initial textural and structural differences between samples with different values of x. After 24 h during the bi-reforming of methane, spent Ce1−xNixOy samples were characterized by the same very-low SBET values (3-5 m 2 /g) and the phase composition, which is CeO2 and Ni 0 with a CSR of 50 nm. Nevertheless, with increasing x, first, a higher hydrogen productivity is achieved, and hence a higher hydrogen concentration in the mixture of reaction products, and second, lower deactivation rates are observed. However, the number of carbonaceous deposits is increasing: samples with a lower degree of deactivation (x = 0.3) have a higher rate of coke accumulation. It can be assumed that these differences in the activity and stability of the Ce1−xNixOy samples are more related to the differences in their anti-sintering properties, rather than in their resistance to coke formation. In situ XAS spectroscopy indicates [117] that under reducing conditions the size and morphology of Ni particles change, as they become flattened and strongly stabilized Thus, the high activation temperature and subsequent harsh reaction conditions smoothed out the initial textural and structural differences between samples with different values of x. After 24 h during the bi-reforming of methane, spent Ce 1−x Ni x O y samples were characterized by the same very-low S BET values (3-5 m 2 /g) and the phase composition, which is CeO 2 and Ni 0 with a CSR of 50 nm. Nevertheless, with increasing x, first, a higher hydrogen productivity is achieved, and hence a higher hydrogen concentration in the mixture of reaction products, and second, lower deactivation rates are observed. However, the number of carbonaceous deposits is increasing: samples with a lower degree of deactivation (x = 0.3) have a higher rate of coke accumulation. It can be assumed that these differences in the activity and stability of the Ce 1−x Ni x O y samples are more related to the differences in their anti-sintering properties, rather than in their resistance to coke formation. In situ XAS spectroscopy indicates [117] that under reducing conditions the size and morphology of Ni particles change, as they become flattened and strongly stabilized on the partially reduced CeO 2 surface which promotes their stability under reaction. Taking into account the absent of carbonaceous deposits for spent samples with x ≤ 0.2, the decrease in activity in the initial period of the reaction before reaching a steady state conversion seems to be mainly associated with phase sintering and a decrease in the number of available active sites. By improving the dispersion of Ni to isolated cations in assynthesized samples, the Ni particle growth mechanism switches from crystal migration to atomic migration, which leads to higher particle growth rates [31]. Therefore, the sintering process is more pronounced for samples with low values of x, which have higher initial nickel dispersion. Among Ce 1−x Ni x O y samples, the sample with x = 0.2 was chosen as the optimal one due to its ability to provide high and stable conversion of CH 4 and CO 2 to synthesis gas, high H 2 productivity, resistance to the formation of carbon deposits, and a moderate deactivation rate.
The introduction of M into the composition of the samples improves their activity during the bi-reforming of methane. In the case of hydrogen yield, this tendency increases in the series La < Mg < Al. In addition, (CeM) 1 samples are characterized by coke formation, the rate of which increases in the series Mg < Al < La. This is in good correlation with the structural properties of materials. To be specific, after activation at M = La, the solid solution is preserved, at M = Al, the Al 3+ cations are only partly retained in the CeO 2 structure, and at M = Mg, the full destruction of solid solution occurs, and CeO 2 and MgO phases co-exist. Given that the M-free sample has better resistance to carbon deposits, this ability is reduced for CeO 2 -based solid solutions due to a decrease in the oxygen mobility at a high doping concentration [118].

Sample Preparation
The procedure for the synthesis of (CeM) 1−x Ni x O y (M = Al, La, Mg) solid solutions by the polymerizable complex method was adapted from our previous studies [59,60]. The molar ratio M/Ce varied within 0-4; the mole fraction of Ni in the sample (x) varied from 0.01 to 0.8. The molar fraction of oxygen (y) was calculated based on the principle of electrical neutrality and the assumption that the cations were in the form of Ce 4+ , Al 3+ , La 3+ , Mg 2+ , and Ni 2+ . The scheme for the synthesis of catalysts is shown in Figure 1 First, appropriate amounts of metal salts were dissolved in water and mixed with a solution of CA in EG, obtained at 70 • C, with stirring. Then, the resulting solution was stirred at 60 • C for 1 h, and ED was dropwise added to promote the polyesterification reaction. After evaporation at 70 • C for 48 h and the formation of a polymer resin, the sample was thermally treated in a muffle furnace at 500 • C for 4 h in air in order to eliminate the organic matrix. The heat treatment details are given in [60].
For comparison, samples with the same chemical composition were obtained by incipient wetness impregnation of the corresponding Ce 1−z M z O y (M = Al, La, Mg, z = 0-1, y = 1-2) supports by an aqueous solution of nickel nitrate with the required concentration.
Supports Ce 1−z M z O y as reference samples were prepared following the procedure of the PC method used for (CeM) 1−x Ni x O y . These materials were calcined also at 500 • C in air.

Sample Characterization
The chemical composition of the samples was determined by X-ray fluorescence analysis in an ARL ADVANT'X analyzer (ThermoTechno Scientific, Ecublens, Switzerland) with an Rh anode of the X-ray tube.
Thermogravimetry (TG) and differential thermogravimetry (DTG) with differential thermal analysis (DTA) of polymeric gel and spent catalysts were performed (NETZSCH STA 449C, Selb, Germany) at a heating rate of 10 • C/min up to 900 • C in flowing air.
The N 2 adsorption-desorption isotherms for the samples were measured on an ASAP 2400 automated volumetric instrument (Micromeritics, Norcross, GA, USA) at -196 • C. The specific surface area (S BET ) was calculated using the BET method. The total pore volume (V p ) and average pore size (D p ) were calculated using the BJH method applied to the desorption branch of the isotherm.
The scanning electron microscopy (SEM) analysis of the samples was conducted using a JSM-6390LA (JEOL, Tokyo, Japan) electron microscope.
The X-ray diffraction (XRD) analysis of the samples ex situ was carried out on an HZG-4C diffractometer (Freiberger Prazisionmechanik, Freiberg, Germany) with CoK α radiation (λ = 1.79021 Å) at room temperature in the 2 alysts 2022, 12, x FOR PEER REVIEW 16 of 40 clearly appear, increases from 500 to 650 °C in the series Mg→Al→La ( Figure 12). For Al and La-containing samples at this temperature (T > 600 °C), the reduction of Ni 2+ cations is practically completed (Figure 10b). For the Mg series, on the contrary, at T = 500 °C, only a part of the Ni 2+ cations is reduced-in the composition of the [Ni-Ce-O] structures. Nickel cations in the composition of NiO-MgO are reduced in a higher temperature region ( Figure 10b). Additionally, for the Al-containing system, a peak at 2Ɵ = 26.4° was observed during reduction, which can be attributed to the oxygen-deficient CeOy phase. Its low intensity indicates a small amount of this phase. Ɵ Supports Ce1−zMzOy as reference samples were prepared following the procedure of the PC method used for (CeM)1−xNixOy. These materials were calcined also at 500 °C in air.

Sample Characterization
The chemical composition of the samples was determined by X-ray fluorescence analysis in an ARL ADVANT'X analyzer (ThermoTechno Scientific, Ecublens, Switzerland) with an Rh anode of the X-ray tube.
Thermogravimetry (TG) and differential thermogravimetry (DTG) with differential thermal analysis (DTA) of polymeric gel and spent catalysts were performed (NETZSCH STA 449C, Selb, Germany) at a heating rate of 10 °C/min up to 900 °C in flowing air.
The N2 adsorption-desorption isotherms for the samples were measured on an ASAP 2400 automated volumetric instrument (Micromeritics, Norcross, GA, USA) at -196°C. The specific surface area (SBET) was calculated using the BET method. The total pore volume (Vp) and average pore size (Dp) were calculated using the BJH method applied to the desorption branch of the isotherm.
The scanning electron microscopy (SEM) analysis of the samples was conducted using a JSM-6390LA (JEOL, Tokyo, Japan) electron microscope.
The X-ray diffraction (XRD) analysis of the samples ex situ was carried out on an HZG-4C diffractometer (Freiberger Prazisionmechanik, Freiberg Germany) with CoKα radiation (λ = 1.79021 Å) at room temperature in the 2Ɵ ranging from 10 to 80° with a step of 0.1 degrees and an accumulation time of 6-15 s. Diffraction peaks were identified using JCPDS powder diffraction databases. Then, according to the Bragg-Wolfe equation, interplanar distances and the parameters of the unit cell (a) were calculated. The accuracy of the determination of the parameters of the unit cell was ±0.003 Ǻ. The coherent scattering region (CSR) was calculated by the Selyakov-Scherrer method from the broadening of the diffraction peak (111), fixed phases having a cubic structure of the fluorite type; CSR (NiO)-peak (200) phases of NiO; CSR (Ni)-peak (200) phases of Ni 0 .
X-ray phase analysis of the samples in situ was conducted on an AXS D8 diffractometer (Bruker, Karlsruhe, Germany) with CuKa radiation (λ = 1.5406 Å) in the 2Ɵ ranging from 20 to 50° with a step of 0.05 degrees and an accumulation time in each point of 5 s. The heating and cooling rate was 12 °C/s, the exposure time of the diffraction pattern at a certain temperature was at least 30 min before each diffraction pattern was recorded. Parameters of the unit cell were determined using IK 2.1. software (BIC SB RAS, Novosibirsk, Russia) and JCPDS powder diffraction databases.
The Raman spectra of the samples were recorded using an excitation wavelength of 514.5 nm on a Renishaw Invia Raman spectrometer (Renishaw plc., Wotton-under-Edge, Gloucestershire, United Kingdom). The power of the laser radiation incident on the sample did not exceed 2 mW, and the accumulation time was 30 s.
Transmission electron microscopy (TEM) studies were carried out using JEM-2010 (JEOL, Tokyo, Japan), JEM-2200FS (JEOL, Tokyo, Japan) and Themis Z (Thermo Fisher Scientific, Waltham, MA, USA) electron microscopes operated at 200 kV. Images in Scanning-TEM (STEM) mode were acquired using high-angle annular dark field (HAADF) detectors. The local elemental analysis of the samples was studied by energy-dispersive X-ray (EDX) spectroscopy using a JEOL JED-2300 (JEM-2200FS) and Thermo Fisher Scientific Super-X (THEMIS Z) EDX detectors. The samples for the TEM studies were dispersed ultrasonically and deposited on copper grids covered with a holey carbon film.
The temperature-programed hydrogen reduction (H2-TPR) was performed on a setup equipped with a flow reactor and a thermal conductivity detector [119]. To eliminate exothermic effects, a 100 mg sample with a size of 250-500 µm was mixed with 100 mg of quartz of the same size. The sample was then pre-treated at 450 °C for 1 h in a stream of air. The reduction was carried out at a heating rate of 10 °C/min from 25 to 900 °C in a flow of 10%H2/90%Ar, 30 cm 3 /min.

Catalytic Activity Testing
. The coherent scattering region (CSR) was calculated by the Selyakov-Scherrer method from the broadening of the diffraction peak (111), fixed phases having a cubic structure of the fluorite type; CSR (NiO)-peak (200) phases of NiO; CSR (Ni)-peak (200) phases of Ni 0 .
X-ray phase analysis of the samples in situ was conducted on an AXS D8 diffractometer (Bruker, Karlsruhe, Germany) with CuK a radiation (λ = 1.5406 Å) in the 2 Catalysts 2022, 12, x FOR PEER REVIEW 16 of 40 clearly appear, increases from 500 to 650 °C in the series Mg→Al→La ( Figure 12). For Al and La-containing samples at this temperature (T > 600 °C), the reduction of Ni 2+ cations is practically completed (Figure 10b). For the Mg series, on the contrary, at T = 500 °C, only a part of the Ni 2+ cations is reduced-in the composition of the [Ni-Ce-O] structures. Nickel cations in the composition of NiO-MgO are reduced in a higher temperature region ( Figure 10b). Additionally, for the Al-containing system, a peak at 2Ɵ = 26.4° was observed during reduction, which can be attributed to the oxygen-deficient CeOy phase. Its low intensity indicates a small amount of this phase. Ɵ ranging from 20 to 50 • with a step of 0.05 degrees and an accumulation time in each point of 5 s. The heating and cooling rate was 12 • C/s, the exposure time of the diffraction pattern at a certain temperature was at least 30 min before each diffraction pattern was recorded. Parameters of the unit cell were determined using IK 2.1. software (BIC SB RAS, Novosibirsk, Russia) and JCPDS powder diffraction databases.
The Raman spectra of the samples were recorded using an excitation wavelength of 514.5 nm on a Renishaw Invia Raman spectrometer (Renishaw plc., Wotton-under-Edge, Gloucestershire, United Kingdom). The power of the laser radiation incident on the sample did not exceed 2 mW, and the accumulation time was 30 s.
Transmission electron microscopy (TEM) studies were carried out using JEM-2010 (JEOL, Tokyo, Japan), JEM-2200FS (JEOL, Tokyo, Japan) and Themis Z (Thermo Fisher Scientific, Waltham, MA, USA) electron microscopes operated at 200 kV. Images in Scanning-TEM (STEM) mode were acquired using high-angle annular dark field (HAADF) detectors. The local elemental analysis of the samples was studied by energy-dispersive X-ray (EDX) spectroscopy using a JEOL JED-2300 (JEM-2200FS) and Thermo Fisher Scientific Super-X (THEMIS Z) EDX detectors. The samples for the TEM studies were dispersed ultrasonically and deposited on copper grids covered with a holey carbon film.
The temperature-programed hydrogen reduction (H 2 -TPR) was performed on a setup equipped with a flow reactor and a thermal conductivity detector [119]. To eliminate exothermic effects, a 100 mg sample with a size of 250-500 µm was mixed with 100 mg of quartz of the same size. The sample was then pre-treated at 450 • C for 1 h in a stream of air. The reduction was carried out at a heating rate of 10 • C/min from 25 to 900 • C in a flow of 10% H 2 /90% Ar, 30 cm 3 /min.

Catalytic Activity Testing
The bi-reforming of CH 4 over prepared catalysts was studied in a fixed-bed flow quartz reactor with an inner diameter of 10 mm, at 1 atm, a gas flow rate of 200 ml N /min, and a molar ratio between reagents of CH 4 :CO 2 :H 2 O:He = 1:0.8:0.4:2.8. The molar ratio used was somewhat different from the stoichiometric CH 4 :CO 2 :H 2 O molar ratio of 1:0.5:0.5. A higher CO 2 /H 2 O ratio (0.8 vs. 0.5) was applied in order to increase carbon dioxide utilization, while a higher O/C ratio (1.1 vs. 1) was used to reduce the rate of coking. Water was supplied by a piston pump at the required rate to the evaporator, where water vapor was mixed with other components (CH 4 , CO 2 , He) and quantitatively introduced into the fixed-bed reactor.
For catalytic activity testing, a 500 mg sample with a grain size of 250-500 µm was used. Initially, the activation of catalyst precursors was performed in situ by thermal treatment at 800 • C for 1 h in a 30% H 2 /70% He flow.
The screening tests were performed in the stepwise temperature rise mode 650 → 850 • C. The heating rate was 10 degrees per minute; the holding time at each temperature was 40 min.
The stability tests were performed for 24 h at 800 • C. The analysis of reaction mixtures was performed using the online automatic gas chromatography system Kristall 2000 m (Yoshkar-Ola, Russia) with a flame ionization detector and thermal conductivity detector.
The catalyst performance was characterized by CH 4 conversion (X CH 4 ), CO 2 conversion (X CO 2 ), yield of H 2 (Y H 2 ), and yield of CO (Y CO ), which were calculated using the following formulas: where F i is the molar flow rate of reagent (i) at the inlet (in) and outlet (out) of the reactor. The thermodynamic equilibrium analysis was conducted using the software package IVTANTHERMO based on the minimization of the Gibbs free energy method [120,121].

Conclusions
Thus, (CeM) 1−x Ni x O y (M = Al, La, Mg) materials served as precursors of Ni catalysts during the bi-reforming of methane. Their genesis features as well as textural, structural, and redox characteristics and functional properties were regulated by variation in the type and content of M. In comparison to the M-free sample, for (CeM) 1−x Ni x O y with an optimal M/Ce ratio, the size of crystallites decreases, the defectiveness and the thermal stability of the material increases, and the temperature of formation of the active Ni 0 phase rises (∆T = 100-200 • C), while the temperature of its re-oxidation decreases (∆T = 70 • C). This all results in the improvement of the strength of the metal-support interaction and provides a more stable and higher yield of hydrogen.
In the case of M = Al, the as-synthesized samples have a homogenous phase composition of a solid fluorite-like solution with mesoporous texture and spongy morphology. Among studied M, the Al-series is characterized by the highest burnout temperature of the organic matrix during synthesis by the PC method, a higher specific surface area, and the smallest solid solution crystallite size (CSR 2.8-4.5 nm). For higher specific surface area (80-100 m 2 /g) the Al/Ce molar ratio should be equal to ≤1. In (CeAl) 1−x Ni x O y samples, the cations of Ni 2+ are stabilized in the structure of a solid solution and their exsolution under a reducing atmosphere leads to the formation of the Ni 0 phase at 600 • C. A feature of the Al-series is that, under activation and reaction conditions, a complete transformation of the oxide matrix occurs: CeO 2 -based solid solution → Ni 0 + cerium aluminate of perovskite structure (Al/Ce = 1). This is accompanied by a significant decrease in the specific surface area and pore volume. (CeAl) 1−x Ni x O y samples provide the highest H 2 yield for the bi-reforming of methane and moderate stability against coking.
In the case of M = La and synthesis by the PC method, the formation of a solid fluorite-like solution (CSR 4.0-6.0 nm) with the highest concentration of oxygen vacancies is observed, and only an insignificant fraction of nickel is stabilized as NiO nanoparticles (3-5 nm) on the sample surface. Among studied M, for samples of the La-series, the burnout of the organic matrix occurs at the lowest temperature. They are characterized by a rapid decrease in the specific surface area and an increase of the defectiveness of the structure with increasing La/Ce molar ratio. A feature of the La-series is that the exsolution of Ni from solid solution is highly hindered, and the Ni 0 phase appears only at a reduction temperature of 650 • C, while the Ce-La-O-solid solution remains. (CeLa) 1−x Ni x O y samples provide the lowest H 2 yield for the bi-reforming of methane and poor stability against coking.
In the case of M = Mg, a two-phase system is already observed for the freshly synthesized sample, which may be due to the low stability of the Mg-citrate complex and partial segregation of the MgO phase at the earliest stage of the complex oxide formation. Therefore, in addition to a solid fluorite-like solution, the formation of a NiO-MgO solid solution (CSR 3.6-8.5 nm) is observed. Due to this, samples have the lowest reducibility. After activation, a three-phase system is formed, including Ni 0 , MgO, and a CeO 2 -based fluoritelike phase. (CeMg) 1−x Ni x O y samples provide high H 2 yield during the bi-reforming of methane and the highest stability against coking.
Taking into account three factors (activity, stability, and the coke accumulation rate), the efficiency of methane bi-reforming catalysts increases in the following order: M-free < La < Al < Mg. The composition Ce 0.6 Mg 0.4 Ni 0.2 O 1.6 is considered the best. During the bi-reforming of methane over Ce 0.6 Mg 0.4 Ni 0.2 O 1.6 , a stable H 2 yield of 90% is achieved at a high level of CO 2 and CH 4 conversions (>85%).