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May 30, 2019 - STCH-active oxide is reduced at a high temperature (TRE, typically, 1300−1500 °C) in low oxygen environments, driving the production...
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Cite This: Inorg. Chem. 2019, 58, 7705−7714

Phase Identification of the Layered Perovskite CexSr2−xMnO4 and Application for Solar Thermochemical Water Splitting Debora R. Barcellos,§ Francisco G. Coury,§ Antoine Emery,† Michael Sanders,§ Jianhua Tong,‡ Anthony McDaniel,⊥ Christopher Wolverton,† Michael Kaufman,§ and Ryan O’Hayre*,§ §

Metallurgical and Materials Engineering Department, Colorado School of Mines, Golden, Colorado, United States Department of Materials Science and Engineering, Northwestern University, Evanston, Illinois, United States ‡ Department of Materials Science and Engineering, Clemson University, Clemson, South Carolina, United States ⊥ Sandia National Laboratory, Livermore, California, United States

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†

S Supporting Information *

ABSTRACT: Ruddlesden−Popper (layered perovskite) phases are attracting significant interest because of their unique potential for many applications requiring mixed ionic and electronic conductivity. Here we report a new, previously undiscovered layered perovskite of composition, CexSr2−xMnO4 (x = 0.1, 0.2, and 0.3). Furthermore, we demonstrate that this new system is suitable for solar thermochemical hydrogen production (STCH). Synchrotron radiation X-ray diffraction and transmission electron microscopy are performed to characterize this new system. Density functional theory calculations of phase stability and oxygen vacancy formation energy (1.76, 2.24, and 2.66 eV/O atom, respectively with increasing Ce content) reinforce the potential of this phase for STCH application. Experimental hydrogen production results show that this materials system produces 2−3 times more hydrogen than the benchmark STCH oxide ceria at a reduction temperature of 1400 °C and an oxidation temperature of 1000 °C. typically, 1300−1500 °C) in low oxygen environments, driving the production of oxygen vacancies. The oxide is then cooled to a lower temperature (TOX, typically, 800−1000 °C), where steam is introduced, and the driving force for reoxidation leads to the dissociation of the hydrogen−oxygen bonds, thereby reoxidizing the STCH material and liberating hydrogen gas. The suitability of a prospective oxide material for STCH application is intrinsically tied to the energy required to create/ fill oxygen vacancies during this cyclic reduction/oxidation process. Ceria, the benchmark material for STCH, represents the upper limit in the range of viable oxygen vacancy formation energy (Ev) at about 5 eV/O atom.11 Any material having an Ev at or above this value will likely require impractically high temperatures (above 1550 °C) for meaningful reduction to occur. The cubic perovskite SrxLa1−xMnyAl1−yO312 (SLMA), having an Ev between 1.5 and 2.5 eV/O atom,13 depending on exact composition, represents the potential lower limit for STCH materials according to Meredig et al.,7,14 because Ev also directly impacts the driving force for oxide reoxidation in the presence of steam. In essence, Ev must be low enough to allow a reasonable level of nonstoichiometry during reduction at

1. INTRODUCTION Perovskite oxides are of great scientific interest due to their compositional versatility and the resulting breadth of physical properties they exhibit. Perovskite oxides are widely studied for numerous applications including solid oxide fuel cells,1 colossal magnetoresistors,2 multiferroics,3 chemical looping,4 and solar thermochemical hydrogen production (STCH).5 Currently, there are about 254 confirmed, experimentally synthesized inorganic perovskites.6 Recent density functional theory (DFT) and machine learning (ML) screening efforts suggest that there are yet hundreds more stable perovskite or perovskite-related oxides remaining to be discovered and synthesized.6−9 Theoretical tools are increasingly used to predict new materials, and DFT can also provide insights into newly discovered compounds. In this paper, we report one such new perovskite-related compound, Ce x Sr 2−x MnO 4 (CSM), a K2NiF4-type10 layered perovskite Ruddlesden− Popper phase. Our discovery of CSM was triggered by a search for new STCH-active materials. Here, we provide detailed analysis validating its structure and show that CSM indeed possesses intriguing thermochemical water-splitting behavior that makes it potentially attractive for STCH application. In a two-step thermochemical water-splitting process, a STCH-active oxide is reduced at a high temperature (TRE, © 2019 American Chemical Society

Received: December 14, 2018 Published: May 30, 2019 7705

DOI: 10.1021/acs.inorgchem.8b03487 Inorg. Chem. 2019, 58, 7705−7714

Article

Inorganic Chemistry reactor-attainable conditions (i.e., TRE < 1400 °C) but must be high enough to overcome the O−H bond strength of water to ensure sufficient driving force for water splitting during the reoxidation step.15 In the search for new STCH-active materials, perovskite and perovskite-related oxides provide fertile hunting grounds because of their compositional flexibility and tolerance for large degrees of nonstoichiometry. A- and B-site cation selection, site-substitution (doping), order/disorder effects, and structural distortions can all be used to alter the thermodynamics of oxygen vacancy formation. The trade-off between the energy requirements for reduction and oxidation, and the demonstrable ability to tune Ev,13,16 is crucial to the continued search for suitable STCH candidates. Barcellos et al. recently showed that the perovskite oxide BaCe0.25Mn0.75O3 (BCM) attains good STCH performance.17 However, BCM forms a line compound; i.e., it can only form with an exact composition, a Ce/Mn ratio of 0.25/0.75, likely due to size constraints on the A- and B-sites, which eliminates the possibility for redox tunability via substitutional doping. Motivated by the success of BCM, but with the goal of identifying related perovskites possessing greater compositional tunability, we hypothesized that Ce could be analogously doped on the B-site of SrMnO3. However, we instead obtained an unexpected structure, part of the Ruddlesden−Popper family AO(ABO3)n (where n = 1),18 known as a K2NiF4-type10 layered perovskite, where Ce shares the A-site with Sr. Although Ce occupies the A-site rather than the B-site in this layered perovskite, it nevertheless shows remarkable activity for solar thermochemical water splitting. Layered perovskites are also of great interest in areas such as SOFCs19,20 and oxygen transport membranes21 thanks to increased oxygen ion mobility22,23 due to the rock salt interlayers formed by AO (where the A cation in the rock salt layer is the same as the A cation in the ABO3 perovskite layer). The AO layers may potentially serve as highways for oxygen ion transport because they can accommodate interstitial oxygen.24,25 Several RP phase materials used as SOFC cathodes have previously been shown to be thermochemically active for water splitting (LaSrCoO4 and LaSrFeO426) with similar performance and behavior to SLMA. The layered perovskite we discovered here, CexSr2−xMnO4 (Figure 1), has never been reported in the literature, motivating the detailed structural analysis from synchrotron radiation X-ray diffraction and high-resolution transmission electron microscopy (HR-TEM) reported in this paper. DFT calculations and preliminary WS experiments suggested that this phase has the potential for water splitting. More extensive testing on purified phase samples corroborate these encouraging observations, highlighting the promise of this RP layered perovskite structure type for STCH and opening new avenues of exploration for the next generation of water-splitting materials.

Figure 1. Illustration of the CexSr2−xMnO4 Ruddlesden−Popper layered perovskite crystal structure. Purple octahedra are MnO6 and green/yellow atoms denote the A-site positions which are randomly shared between Sr and Ce atoms.

Once the gel was formed, it was transferred to a drying oven at 160 °C to dry completely. The resulting charcoal-like samples were calcined first at 800 °C for 10 h and then subsequently at 1400 °C for 5 h. Although the goal of these syntheses was to produce a range of Bsite doped SrCeyMn1−yO3 compounds, in all cases (expect for the y = 0 and y = 1 end-member compositions) a mixture of SrCeO3 and/or SrMnO3, CeO2, and the new A-site doped layered perovskite CexSr2−xMnO4 was obtained instead (further details in the Results and Discussion section). Subsequently, we targeted pure-phase synthesis of this new CexSr2−xMnO4 layered perovskite phase. Three CexSr2−xMnO4 samples (x = 0.1, 0.2, and 0.3) were synthesized following the evaporation-to-dryness procedure of Karita et al.,28,29 where nitrates of the metal precursors, in this case strontium nitrate (Alfa Aesar 99%), cerium nitrate hexahydrate (Alfa Aesar 99%), and manganese nitrate tetrahydrate (Alfa Aesar 99%) in the stoichiometric ratios were stirred in deionized water and heated at 180 °C until dry. The samples were calcined at 350 °C for 4 h to allow for the decomposition of the nitrates and then at 1000 °C for 10 h. A second calcination took place at 1400 °C for 10 h. 2.2. Characterization. Powder X-ray diffraction (XRD) patterns were measured on a PANalytical PW3040 diffractometer surveying the 2θ range between 20° and 120° at a step size of 0.008° with Cu− Kα radiation of λ = 1.540598 and 1.54439 Å.

2. EXPERIMENTAL SECTION 2.1. Synthesis. Initial experiments attempted B-site doping of SrMnO3 by Ce. Five samples spanning the nominal SrCeyMn1−yO3 compositional space were synthesized with y = 0, 0.3, 0.5, 0.7, and 1 using the sol−gel modified Pechini synthesis route described by Shang et al.27 Stoichiometric amounts of strontium nitrate (Alfa Aesar 99%), cerium nitrate hexahydrate (Alfa Aesar 99.5%), and manganese acetate tetrahydrate (Alfa Aesar 98+%) were stirred in deionized water, with ethylenediaminetetraacetic acid (EDTA), citric acid, and ammonium hydroxide for pH adjustment while heating at 90 °C. 7706

DOI: 10.1021/acs.inorgchem.8b03487 Inorg. Chem. 2019, 58, 7705−7714

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lations,41 because of the use of DFT+U (Ce and Mn) and the phase transition between 0 and 300 K (O).39 Compounds that are stable have a zero hull distance, whereas compounds that are unstable have a positive hull distance. The mixed perovskite CeaSrbMn1‑aO3 compounds were generated by mixing Ce and Mn on the Mn site of the cubic SrMnO3 perovskite with the special quasirandom structure approach (SQS).42 SQS are structures generated so that their short-range correlation functions (Π̅ f (σ )) match those of a random alloy, thus allowing the calculation of random mixing despite periodic boundary conditions. SQS were generated using the Monte Carlo semi quasirandom structure (mcsqs) as implemented in the Alloy Theoretic Automated Toolkit (ATAT).43 The structures of CexSr2‑xMnO4 were created by taking LaSrMnO4 from the ICSD (collection code number 153645) and substituting Ce for La. The SQS approach was used to model the cation mixing on the (Ce−Sr) A-site. The oxygen vacancy formation energies were calculated by using a supercell approach (20 formula units) according to

For structural characterization of the new phase (CexSr2−xMnO4), three samples were probed by synchrotron radiation through the mailin program offered by the Advanced Photon Source (APS) 11-BM at Argonne National Laboratory (λ = 0.412628 Å powder XRD). Rietveld refinement of the resulting diffraction patterns was conducted using the free software GSAS-II30 by indexing the patterns to the tetragonal structure of space group I4/mmm (139). Selected samples were also analyzed by transmission electron microscopy (TEM) using an FEI Talos operating at 200 kV. The microscope was equipped with four energy dispersive X-ray spectroscopy (EDS) detectors optimized for compositional mapping. The TEM samples were prepared by focused ion beam (FIB) on an FEI Helios equipped with a dual-beam system that uses an electron gun as well as a Ga gun for the milling steps. The milling was performed at 30 kV, but a final thinning step at 2 kV was performed to minimize surface damage. Images obtained from high-resolution electron microscopy (HRTEM) were processed using the CRISP software31 following the standard procedure described elsewhere32 to identify the unit cell in the image by the Fourier Transform, averaging over a region of several unit cells and imposing the known symmetry. Only the Fourier components with a resolution smaller than the first crossover of the image were used for the image at the Scherzer defocus condition. The regions from the image that yielded the smallest phase residuals were used, which also coincided with the thinnest regions of the crystal. These processed images were compared with Multislice simulations performed using the jEMS software for different defocus and thickness values. Once the thickness was discovered by finding the best match between the simulated and experimental images, the other defocus images were compared with the respective defocus simulated image for the same thickness. 2.3. DFT Calculations. DFT calculations were performed using the Vienna ab initio simulation package (VASP)33,34 with projectoraugmented wave method (PAW)35 potentials with the Perdew− Burke−Ernzerhof generalized gradient approximation36 to the XC functional. DFT+U37,38 was used with Hubbard U added to the Mn d and Ce f states (U = 3.8 and 5 eV, respectively, J = 0). U and J values are used to describe the on-site Coulomb interactions of localized electrons, which are not typically well described in DFT. The strength of the on-site interaction is represented by an on-site Coulomb interaction U and an on-site exchange J. These corrections are often important for compounds with localized d and f electrons. The stability of the compounds was calculated with respect to all phases in the Sr−Ce−Mn−O system present in the Open Quantum Materials Database (OQMD).39 We note that all compounds containing cerium were not simply taken from OQMD but were recalculated (in the present work) with the Ce pseudopotential containing 12 valence electrons (i.e., including Ce f as valence). The total number of phases in the OQMD for this quaternary system is 230. From these computed energies, the ground state convex hull was computed to assess the stability of phases in this quaternary system. Among the 230 total phases, 23 were stable before adding the CeaSrbMncOd phases. The stability, or convex hull distance, for any arbitrary CeaSrbMncOd compound is defined as CeaSrbMncOd ΔHstab = ΔHfCeaSrbMncOd − ΔHfCH

ΔEvO = E(Ce20xSr20(2 − x)Mn20O79) + μO − E(CexSr20(2 − x)Mn20O80

(3)

where E(Ce20xSr20(2−x)Mn20O79) and E(Ce20xSr20(2−x)Mn20O80) (x = 0.1, 0.2, 0.3) are the DFT total energies of the defected and pristine supercells, respectively, and μO is the chemical potential of oxygen, which was detailed above. As not all oxygen atoms are equivalent, we performed several calculations where a symmetrically distinct oxygen atom is removed. We report the oxygen vacancy formation energies of the vacancy with the lowest energy, i.e., the oxygen site that is the most likely to reduce first. 2.4. Testing for STCH. Temperature-programmed reduction (TPR) was conducted to measure the extent of reduction as a function of temperature, as a preliminary indicator of a materials’ potential for STCH. About 50 mg of sample was loaded into a Pt crucible and heated in a Setaram SetSys Evo thermogravimetric analyzer (TGA) to the reduction temperature (TRE) of 1350 °C in ultra high purity (UHP) nitrogen at a ramp rate of 10 °C/min. After a dwell of 1 h, the sample was cooled down to the oxidation temperature (TOX) of 1000 °C at the same temperature ramp rate and air was introduced. An oxygen sensor connected to the exhaust of the system monitored the oxygen content to confirm that the mass loss was due to oxygen release. The mass profile was converted into moles of atomic oxygen per mole of sample (or δ as in the ABO3‑δ) and was compared for all the samples tested. Water-splitting performance tests were performed in the stagnation flow reactor (SFR) at Sandia National Laboratories (configurations of the system and experiment are detailed elsewhere44,45). Briefly, a powder sample on a zirconia platform was maintained at the oxidation temperature (1000 °C) by a tubular furnace in argon atmosphere at 75 Torr. The sample was heated at the controlled rate of 10 °C/s to 1400 °C by a laser. The reduction temperature was held for 330 s, and the amount of oxygen released during reduction was measured in a calibrated mass spectrometer. The samples were then immediately cooled back down to the oxidation temperature by turning off the laser and steam (40 vol % balanced in argon) was introduced to the reactor. The amount of hydrogen produced during the oxidation step and oxygen evolved from the samples during the reduction step were measured by a mass spectrometer, and the detected amounts of hydrogen and oxygen were normalized by the mass of sample.

(1)

ΔHCH f

is the convex hull (CH) energy at the CeaSrbMncOd where aSrbMncOd composition and ΔHCe is the formation energy of the f CeaSrbMncOd compound calculated in eq 2) as ΔHfCeaSrbMncOd = E(CeaSrbMncOd ) − aμCe − bμSr − cμMn − dμO (2)

3. RESULTS AND DISCUSSION Motivated by the perovskite BaCe0.25Mn0.75O3, first synthesized by Fuentes et al.46 and found to be a promising STCH candidate by Barcellos et al.,17 we attempted to synthesize five samples in the SrCeyMn1−yO3 system (y = 0, 0.3, 0.5, 0.7, and 1; denoted SMO, SCM3070, SCM5050, SCM7030, and SCO respectively) via sol−gel with the goal of having the Ce cation

where E(CeaSrbMncOd) is the DFT total energy of the CeaSrbMncOd compound, and μCe, μSr, μMn, and μO are the chemical potentials of strontium, cerium, manganese, and oxygen, respectively. The chemical potential of Sr is the DFT 0 K energy of the element. For Ce, Mn, and O, the chemical potentials are fitted to experimental data using the FERE technique,40 as well as the approach of Jain et al. to combine Generalized Gradient Approximation (GGA) and GGA+U calcu7707

DOI: 10.1021/acs.inorgchem.8b03487 Inorg. Chem. 2019, 58, 7705−7714

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layered perovskite solid solutions and find that these materials could be suitable for WS based on oxygen vacancy formation energies in the range of 1.8−2.7 eV (results summarized in Table 2). While the reported limit for La into the

sharing the perovskite B-site with Mn. The Goldschmidt tolerance factor calculations (0.99, 0.96, and 0.93 respectively) indicate that it should be possible to form cubic SrCe0.3Mn0.7O3, SrCe0.5Mn0.5O3, and SrCe0.7Mn0.3O3 respectively. However, powder X-ray diffraction patterns obtained from the calcined samples reveal that the target compositions are not formed, and all samples show the presence of mixed phases. Moreover, Rietveld refinement summarized in Table S1 indicates the presence of an unexpected phase common to all the Ce-containing samples. This phase is identified as belonging to the space group I4/mmm (139), known as a K2NiF4-type10 Ruddlesden−Popper phase, with a structure similar to that of SrLaCoO4.47 We turned to DFT calculations to corroborate the relative stability of this previously unknown phase within the Sr−Ce− Mn−O system. The DFT-predicted relative phase stabilities and decomposition reactions are summarized in Table 1. The

Table 2. DFT Stability and Oxygen Vacancy Formation Energy Calculations on CSM1, CSM2, and CSM3

formula SrMnO3 SrCe0.25Mn0.75O3

−0.026 0.057

SrCe0.5Mn0.5O3

0.062

SrCe0.75Mn0.25O3

0.059

SrCeO3

0.018

decomposition reactions stable 0.250 CeO2 + 0.250 Sr4Mn3O10 0.450 CeO2 + 0.050 SrO2 + 0.500 Ce0.1Sr1.9MnO4 0.475 CeO2 + 0.025 SrO2 + 0.250 Ce0.1Sr1.9MnO4 + 0.250 Sr2CeO4 0.500 CeO2 + 0.500 Sr2CeO4

Ce0.1Sr1.9MnO4 Ce0.2Sr1.8MnO4 Ce0.3Sr1.7MnO4 CeSrMnO4

−0.012 0.000 −0.007 0.061

decomposition reactions stable stable stable 0.500 Ce2O3 + 0.500 Sr2Mn2O5

oxygen vacancy formation energy [eV/O atom] 1.76 2.24 2.66

LaxSr2−xMnO4 structure is x = 1,48 our experimental results suggest that less Ce can be incorporated in the CexSr2−xMnO4 structure. DFT calculations corroborate this observation, as CeSrMnO4 is calculated to be unstable. Preliminary WS experiments (Figure S2 and Table S2) performed on the mixed-phase samples with the largest amounts of the CexSr2−xMnO4 phase (SCM5050 and SCM7030), as well as on the end members, indicated that this K2NiF4-type phase is in fact thermochemically active. To facilitate further study, we attempted to isolate the new phase in high purity via an evaporation-to-dryness synthesis approach, as further described in the Experimental Section. This relatively simple synthesis approach utilizes nitrate precursors (rather than carbonates) for all cations. This factor likely contributes to its success in producing a phase-pure compound as the nitrate precursor decompositions occurs at much lower temperatures than corresponding carbonate decompositions, thereby kinetically promoting the formation of the K2NiF4-type structure. Ultimately, three nearly phasepure samples, Ce 0.1 Sr 1.9 MnO 4 (CSM1), Ce 0.2 Sr 1.8 MnO 4 (CSM2), and Ce0.3Sr1.7MnO4 (CSM3), were successfully synthesized by the evaporation-to-dryness method. Synchrotron radiation XRD was employed for structural analysis, and the obtained diffractograms are shown below in Figure 2a−c for the samples CSM1, CSM2, and CSM3, respectively. XRD reveals minor secondary phase peaks that increase with increasing Ce content. These peaks do not correspond to any known Ce, Sr, or Mn oxide phases, but since they increase with increasing Ce addition, they are possibly an unknown Ce-rich phase or potentially higher order RP phases (which do not significantly change the main peaks but add small reflection peaks). However, the amount of secondary phase is minimal, even for the CSM3, and is most likely due to inefficient mixing during synthesis, and/or incomplete calcination. There is also some variation in the main peak intensities between the three compositions. While there are no reported polymorphs of the parent Sr2MnO4 phase, many Mncontaining oxides exhibit polymorphism, including several different hexagonal and orthorhombic distortions in SrMnO3.49,50 These distortions could account for the variations; however, their presence does not appear to appreciably affect the WS performance. Rietveld refinement was performed with GSAS-II30 using the Powder Diffraction File number 04-002-4381 on Sr0.8La1.2CoO4, modified using the software VESTA51 to

Table 1. DFT Calculations of Stability, Decomposition Reactions, and Oxygen Vacancy Formation Energy of the BSite Mixed Perovskites SrCeyMn1−yO3 (y = 0, 0.25, 0.5, 0.75, 1)a stability [eV/ atom]

formula

stability [eV/ atom]

oxygen vacancy formation energy [eV/O atom] 1.42

a

The stability is calculated based on the distance from ground state convex hull and the decomposition reactions indicate the stable phases on convex hull.

DFT calculated stabilities in Table 1 indicate that the SrMnO3 perovskite is stable, whereas B-site mixing of Ce into this compound results in perovskite phases that are not on the convex hull (have a positive convex hull distance) and hence unstable. For these unstable phases, we also indicate in Table 1 the stable phases on the convex hull for each composition. These decomposition reactions, or convex hull phases, show that one of the competing, stable phases is the layered perovskite, Sr1.9Ce0.1MnO4. In agreement with experimental observations, the DFT calculations also predict the stability of SrMnO3 perovskite and mixed Sr2‑xCexMnO4 layered perovskites, and the instability of mixed SrCe1−xMnxO3 perovskites. The calculated stabilities indicate that the SrMnO3 perovskite is stable, whereas doping this compound with Ce on the B-site results in an instability and decomposition into other phases, one of them being the layered perovskite. The confirmation of the formation and stability of the layered perovskite motivated additional DFT calculations to further examine the stability of this phase as a function of the Ce content. DFT calculations confirmed the stability of CexSr2−xMnO4 single-phase solutions across a range of Ce concentrations (e.g., x = 0.1, 0.2, 0.3). In addition, we have computed the oxygen vacancy formation energy of these 7708

DOI: 10.1021/acs.inorgchem.8b03487 Inorg. Chem. 2019, 58, 7705−7714

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Inorganic Chemistry

Figure 2. Powder X-ray diffraction using synchrotron radiation (λ = 0.412628 Å) on (a) Ce0.1Sr1.9MnO4 (CSM1), (b) Ce0.2Sr1.8MnO4 (CSM2), and (c) Ce0.3Sr1.7MnO4 (CSM3). Observed (red), calculated (black), difference plot (blue), and expected phase peaks (vertical marks).

ment confirms that Ce is being incorporated into the structure and that CSM3 has the largest amount of Ce compared to CSM2 and CSM1. Furthermore, the lattice parameters generally increase with increasing Ce content, from CSM1 to CSM3. While the physical origin of this increase is still under investigation and Ce3+ ions are smaller than the replaced Sr2+, this substitution requires charge compensation, which is most likely accommodated by a corresponding change of Mn4+ to Mn3+ (which produces an ionic radius increase), thereby causing a slight net increase in the lattice parameter. A thorough transmission electron microscopy (TEM) study using a combination of electron diffraction and high-resolution transmission electron microscopy (HRTEM) was performed on sample CSM1, which is the closest to pure phase, to further characterize its structure. The first step of the TEM analysis consisted of taking a series of selected area diffraction patterns

account for the different A-site cation ratios and cation species Ce and Mn instead of La and Co, respectively. Using the layered perovskite nomenclature, (AO)(A1‑xA′xBO3)n where n = 1, the Ce cations were modeled as sharing the perovskite Asites randomly. The refinement analysis results are shown below in Table 3. The goodness of fitting (GOF) for the refinements are reasonable but reflect increasing discrepancies as the amount of secondary phases increases. Only the isotropic atomic displacement parameter, more commonly known as the temperature factor (Uiso), is reported, since neutron diffraction is necessary to differentiate anisotropic U factors. Further details of the crystal structure investigation(s) may be obtained from the Fachinformationszentrum Karlsruhe, D-76344 Eggenstein-Leopoldshafen (Germany), on quoting the depository number numbers CSD-1873754, CSD-1873755, and CSD-1873756. The site fraction refine7709

DOI: 10.1021/acs.inorgchem.8b03487 Inorg. Chem. 2019, 58, 7705−7714

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Inorganic Chemistry Table 3. Final Structural Parameters for Ce0.1Sr1.9MnO4, Ce0.2Sr1.8MnO4, and Ce0.3Sr1.7MnO4 atom sitea

Ce (1)

Ce0.1Sr1.9MnO4b x/a 0.0 y/b 0.0 z/c 0.3567(4) Uiso 0.009(6) frac 0.084(2) site Ce0.2Sr1.8MnO4c x/a 0.0 y/b 0.0 z/c 0.3574(7) Uiso 0.006(2) frac 0.099(8) site Ce0.3Sr1.7MnO4d x/a 0.0 y/b 0.0 z/c 0.3567(4) Uiso 0.008(7) frac 0.138(4) site

Sr (1)

Mn (1)

O (1)

O (2)

0.0 0.0 0.3567(4) 0.009(6) 0.915(8)

0.0 0.0 0.0 0.007(3) 1.000

0.0 0.5 0.0 0.008(2) 1.000

0.0 0.0 0.1541(5) 0.015(2) 1.000

0.0 0.0 0.3574(7) 0.005(2) 0.900(2)

0.0 0.0 0.0 0.004(1) 1.000

0.0 0.5 0.0 0.005(2) 1.000

0.0 0.0 0.1533(1) 0.011(1) 0.974(9)

0.0 0.0 0.3567(4) 0.008(3) 0.861(6)

0.0 0.0 0.0 0.006(1) 1.000

0.0 0.5 0.0 0.014(9) 1.000

0.0 0.0 0.1577(9) 0.006(1) 1.000

Figure 3. Indexed SADPs (left) and CBED patterns (right) of the three main zone axes of the Ce0.1Sr1.9MnO4 layered perovskite: (a) [001], (b) [100], and (c) [110]. The yellow lines represent the mirror symmetries present in each CBED pattern, and they confirm the I4/ mmm space group. The diffraction patterns are not necessarily at the same scale.

a Space group: I4/mmm (139). bFor Ce0.1Sr1.9MnO4, a = 3.814(8) Å, c = 12.45(5) Å; V = 181.2(5) Å3, GOF = 2.76 cFor Ce0.2Sr1.8MnO4, a = 3.827(8) Å, c = 12.45(8) Å; V = 182.5(4) Å3, GOF = 3.73 dFor Ce0.3Sr1.7MnO4, a = 3.825(2) Å, c = 12.46(8) Å; V = 182.4(3) Å3, GOF = 4.92

imposed in both zone axes, which forces phase and amplitude restrictions on symmetrically related reflections, and an image was reconstructed from these modified structure factors. This procedure, valid since the symmetry of the zone axis is known and has been proved by CBED, corrects several aberrations from the microscope and the image acquisition process. These experimental images were compared to multislice simulations produced by the jEMS software, with information derived from the refined atomic positions of the synchrotron XRD patterns. The multislice calculations were performed by varying the defocus and the sample thickness. On the basis of image comparison, the thickness of the sample was determined to be around 28 nm; the simulated and experimental images for different defocus conditions can be seen in Figure 4. The simulated images match well with the experimental ones; the heavier Ce and Sr atoms can be easily distinguished, especially on the [100] projection, in which the layered structure is confirmed. EDS analysis was also performed in the TEM to verify that the chemical composition of the layered perovskite was consistent with expectations. The Sr to Ce ratios measured in samples CSM1, CSM2, and CSM3 were 16.9, 8.9, and 5.9, respectively, close to the nominal expected values of 19, 9, and 5.7. Table S3 in the Supporting Information shows the EDS results. The DFT calculated oxygen vacancy formation energies of the CSM family (1.8, 2.2, and 2.7 eV/O atom for CSM1, CSM2, and CSM3 respectively, see Table 2) are within the target range of 1.5−5 eV/O atom required for potential STCH materials.7,13 In addition, layered perovskites also potentially enable enhanced oxygen ion mobility (e.g., Ivanov et al.53,54 have attributed the increased mobility of the layered perovskite LaxSr2‑xMnO4 compared to the regular LaxSr1−xMnO3 to the

(SADP) and convergent beam electron diffractions (CBED). The combination of these two techniques can confirm the symmetry of a given phase. SADP and CBED patterns of all three major zone axis, namely, the [001], [100], and [110], are given in Figure 3. The CBED pattern from the [001] zone axis reveals a 4 mm symmetry, while the [100] and [110] patterns both reveal a 2 mm symmetry. The diffraction pattern in Figure 3a is rotated 90° relative to the diffraction patterns in b and c. The SADPs displayed in Figure 3 also confirm the I centering of the phase, due to the systematic absence of the h + k + l = (2n − 1) reflections. No other systematic absences were observed; therefore, the space group is symmorphic. Overall, the electron diffraction results confirm that the space group of the perovskite is I4/mmm. The measured lattice parameters by TEM are a = 3.8 Å and c = 12.5 Å, which match the values measured from synchrotron XRD (Table 2), further confirming this is indeed the phase present in the CSM1 sample. The CSM2 and CSM3 samples were also analyzed by TEM electron diffraction, confirming the presence of the same phase (Figure S3 in the Supporting Information). The layered structure was confirmed using HRTEM by a combination of image processing and simulation using a defocus series. HRTEM images were acquired in the two shortest projection axes, the [100] and the [110], under different defocus conditions. Two processing steps were performed on the images using the software CRISP.52 The first imposed an average over hundreds of unit cells to remove the random noise from the HRTEM images, while the second consisted of taking a Fourier transform of the image to extract its Fourier coefficients, which in turn were related to the structure factor of the sample. The 2 mm symmetry was 7710

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Inorganic Chemistry

Figure 4. Simulated and experimental images of the two shortest projections of the layered perovskite under two different defocus conditions. The projections of the unit cells on the same orientations are displayed for comparison.

Figure 5. Water-splitting results for two complete cycles at TRE = 1400 °C for 330 s and at TOX = 1000 °C for 1200 s. Blue lines are oxygen peaks evolved during reduction, and red lines are hydrogen peaks recorded during the reoxidation (water splitting) step. The peaks were integrated to calculate the total amount of hydrogen and oxygen produced, with the results listed in the legend. For the oxygen values, the first value indicates the oxygen produced by the initial reduction, with the second value corresponding to the subsequent cycles average. The hydrogen value refers to the average hydrogen production. Note that, while theoretically, the hydrogen production should be exactly twice that of the oxygen produced in the subsequent reduction step, the oxygen values are typically over-represented due to integration errors and access to additional reduction toward equilibrium that is not achieved during the first reduction cycle.

1400 °C and a TOX of 1000 °C with 40 vol % steam (Figure 5). CSM2 shows the greatest per-cycle H2 productivity (247 μmol/g) because it balances a moderate extent of reduction with adequate driving force (as quantified by the oxygen vacancy formation energy) to approach complete reoxidation in steam. CSM1 yields a larger extent of reduction but cannot fully reoxidize in steam due to the relatively low oxygen vacancy formation energy. Thus, its full redox capacity is not utilized. CSM3 has a relatively small extent of reduction (and therefore a small redox capacity) because it has the highest oxygen vacancy formation energy of the three compositions. In general, under the same testing conditions, the hydrogen production performance of CSM3 is comparable to that of BCM, at 181 μmol/g, and much better than ceria’s 71 μmol/ g.17 The behavior of all three compositions is generally consistent with that of STCH materials having the Ev values

layered aspect of this structure). These characteristics make CSM an excellent potential candidate for STCH application. To confirm the potential for water splitting seen in the preliminary testing of the mixed-phase samples, TPR was conducted on the three phase-pure samples (shown in Figure S4). All three samples show complete redox reversibility under the TPR test cycle conditions. Consistent with the oxygen vacancy formation energy calculations, CSM1, having the smallest amount of Ce (and the lowest oxygen vacancy formation energy), has the largest extent of reduction (δ), at 0.18. CSM3, which contains the largest amount of Ce (and therefore the highest oxygen vacancy formation energy), has the smallest δ (0.09). Motivated by these promising TPR results, we tested all three phase-pure samples’ water-splitting capability using Sandia National Laboratories SFR test facility using a TRE of 7711

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Inorganic Chemistry predicted by our DFT calculations. This agreement represents an important advancement in expanding the use of DFT calculations to probe for new STCH materials.

Christopher Wolverton: 0000-0003-2248-474X Notes

The authors declare no competing financial interest.

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4. CONCLUSIONS Motivated by the development of BCM as a successful perovskite material for STCH, we expanded our search for new Ce- and Mn-containing perovskites. We replaced Ba with Sr on the A-site to permit greater structural and compositional flexibility, since BCM is a line-compound material that lacks this compositional tunability. In attempting to synthesize the target SrCeyMn1−yO3 compositional family, however, we instead encountered the unexpected formation of a layered perovskite CexSr2−xMnO4, previously unreported in the literature for this compositional family. DFT calculations played a key role in the iterative process of identifying and corroborating the stability of the new phase. Nearly phase-pure Ce0.1Sr1.9MnO4, Ce0.2Sr1.8MnO4, and Ce0.3Sr1.7MnO4 compositions were successfully synthesized by the evaporation-todryness method, and the structures were characterized by synchrotron radiation X-ray diffraction and transmission electron microscopy. The oxygen vacancy formation energy calculated from DFT in parallel with the stability calculations flagged all three compounds as good candidates for water splitting, values that were found to be consistent with results from TPR. The samples were then tested for water splitting, and two of the compositions were shown to produce more hydrogen than ceria and BCM under identical test conditions. This new phase and its promising performance further motivates the search for perovskite-related compounds with unique and potentially valuable STCH properties. The DFT stability calculations and the calculated oxygen vacancy formation energy predictions of WS performance highlight the possibility of harnessing computation to accelerate the development of new STCHactive materials.

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ACKNOWLEDGMENTS The authors would like to acknowledge the support of the 11BM team in the use of the Advanced Photon Source at Argonne National Laboratory supported by the U.S. Department of Energy, Office of Science, Office of Basic Energy Sciences, under Contract No. DE-AC02-06CH11357. The authors gratefully acknowledge research support from the HydroGEN Advanced Water Splitting Materials Consortium, established as part of the Energy Materials Network under the U.S. Department of Energy, Office of Energy Efficiency and Renewable Energy, Fuel Cell Technologies Office, under Award Number DE-EE0008087 and Award Number DEEE0008089 (supporting DFT calculations). The electron microscopy work was supported by the Brazilian Centro Nacional de Pesquisa (CNPq) [Scholarship Process Number 233746/2014-5] and the Center for Advanced Non-Ferrous Structural Alloys (CANFSA), a National Science Foundation Industry/University Cooperative Research Center (I/UCRC) [Award No. 1624836], at the Colorado School of Mines (CSM), in Golden, CO, USA. Finally, we are grateful for the financial support [Scholarship Process Number BEX: 1345213-4] provided by CAPES (Coordination for the Improvement of Higher Education Personnel) and the Science Without Borders Program granted by the Brazilian government.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.inorgchem.8b03487. Initial XRD, Rietveld analysis, and water-splitting data (for SMO, SCM3070, SCM5050, SCM7030, and SCO), TEM/EDS compositional information (for CSM1-3), and TPR results comparing CSM1-3 to ceria and BCM (PDF) Accession Codes

CCDC 1873754−1873756 contain the supplementary crystallographic data for this paper. These data can be obtained free of charge via www.ccdc.cam.ac.uk/data_request/cif, or by emailing [email protected], or by contacting The Cambridge Crystallographic Data Centre, 12 Union Road, Cambridge CB2 1EZ, UK; fax: +44 1223 336033.

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REFERENCES

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Michael Sanders: 0000-0001-6366-5219 Jianhua Tong: 0000-0002-0684-1658 7712

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