BACKGROUND OF THE INVENTION
The present invention relates to oxidation-reduction catalysts that include strontium cobaltite materials.
Transition-metal oxides (TMOs) have been studied for energy technologies because of their physical properties. In particular, owing to the ionic and electronic conductivity offered from the flexibility of transition metal's charge states, multi-valent TMOs have attracted attention for potential applications as catalysts.
Many energy storage and sensor devices rely on atoms that are convertible from one valence state to another. For example, catalytic gas converters use platinum-based metals to transform harmful emissions such as carbon monoxide into nontoxic gases by adding oxygen. Less expensive oxide-based alternatives to platinum usually require very high temperatures, at least 600-700° C., to trigger the redox reaction, making such materials impractical in conventional applications.
SUMMARY OF THE INVENTION
The present invention provides an epitaxially stabilized strontium cobaltite catalyst adapted to rapidly transition between oxidation states at substantially low temperatures.
In one embodiment there is provided an article having a thin film of the epitaxially stabilized strontium cobaltite catalyst.
In one embodiment there is provided a method of transitioning a strontium cobaltite catalyst from a first oxidation state to a second oxidation state comprising the steps of: providing a substrate supporting at least a thin film or layer of the catalyst in the first oxidation state, elevating the temperature of the catalyst, and providing a vacuum atmosphere to the catalyst in the first oxidation state. The temperature is elevated to within a range of 210° C. to 320° C. for a first period of time to effect the transition of the catalyst to the second oxidation state.
Further, the method may provide the steps of reversing the transitioning of the catalyst from the first oxidation state to the second oxidation state by transitioning the catalyst from the second oxidation state to the first oxidation state by: elevating the temperature of the catalyst in the second oxidation state to and providing oxygen for a second period of time, wherein the second period of time is sufficient to effect the transition of the catalyst to the first oxidation state.
In one embodiment there is provided an epitaxially stabilized thin film of strontium cobaltite adapted to transition between SrCoO2.5 and SrCoO3-δ phases.
In a further embodiment, the film of strontium cobaltite is stabilized on a substrate selected from perovskite (ABO3, A: alkaline or alkaline earth element, B: transition metal, and O: oxygen group including SrTiO3, and (LaAlO3)0.3—(Sr2AlTaO6)0.7.
These and other objects, advantages, and features of the invention will be more fully understood and appreciated by reference to the description of the current embodiment and the drawings.
BRIEF DESCRIPTION OF THE DRAWINGS
FIG. 1 is a schematic depiction of a strontium cobaltite thin film in the brownmillerite phase;
FIG. 2 is a schematic depiction of a strontium cobaltite thin film in the perovskite phase;
FIG. 3A is an x-ray diffraction scan pattern of brownmillerite-type SrCoO2.5;
FIG. 3B is an x-ray diffraction scan pattern of perovskite-type SrCoO3-δ;
FIG. 4A is an x-ray rocking curve scan of the brownmillerite phase of strontium cobaltite;
FIG. 4B is an x-ray rocking curve scan of the perovskite phase of strontium cobaltite;
FIG. 4C is a reciprocal space map of SrCoO2.5 on STO substrate;
FIG. 4D is a reciprocal space map of SrCoO3-δ on LSAT substrate;
FIG. 5A is a polarized x-ray absorption spectroscopy (XAS) graph of the oxygen K-edge peaks for two phases of strontium cobaltite;
FIG. 5B is a polarized x-ray absorption spectroscopy (XAS) graph of the cobalt L-edge peaks for two phases of strontium cobaltite;
FIG. 5C is an x-ray magnetic circular dichroism (“XMCD”) spectra of two phases of strontium cobaltite;
FIG. 6A is a graphical representation of the dependent magnetization of two phases of strontium cobaltite;
FIG. 6B is a graph of magnetic hysteresis loops for two phases of strontium cobaltite;
FIG. 6C is a graph of the resistivity of two phases of strontium cobaltite as a function of temperature;
FIG. 6D is a graph of the thermoelectromotive force of two phases of strontium cobaltite at 300K;
FIG. 7A depicts the phase transition of strontium cobaltite from the perovskite phase to the brownmillerite phase;
FIG. 7B depicts the phase transition of strontium cobaltite from the brownmillerite phase to the perovskite phase;
FIG. 8 is a graphical depiction of the energy differences between strontium cobaltite and strontium manganite with different oxygen contents.
FIG. 9A is a graphical representation of catalytic activity of brownmillerite strontium cobaltite as carbon monoxide consumption; and
FIG. 9B is a graphical representation of the catalytic activity as carbon dioxide production.
Before the embodiments of the invention are explained in detail, it is to be understood that the invention is not limited to the details of operation or to the details of construction and the arrangement of the components set forth in the following description or illustrated in the drawings. The invention may be implemented in various other embodiments and is capable of being practiced or being carried out in alternative ways not expressly disclosed herein. Also, it is to be understood that the phraseology and terminology used herein are for the purpose of description and should not be regarded as limiting. The use of “including” and “comprising” and variations thereof is meant to encompass the items listed thereafter and equivalents thereof as well as additional items and equivalents thereof. Further, enumeration may be used in the description of various embodiments. Unless otherwise expressly stated, the use of enumeration should not be construed as limiting the invention to any specific order or number of components. Nor should the use of enumeration be construed as excluding from the scope of the invention any additional steps or components that might be combined with or into the enumerated steps or components.
DESCRIPTION OF THE CURRENT EMBODIMENT
A catalyst in accordance with an embodiment of the invention is shown in FIGS. 1-2 and is generally designated 10 and 20, respectively. The catalyst 10 is shown in FIG. 1 as a strontium cobaltite (SrCoO2.5) thin film in a brownmillerite phase. SrCoO2.5 is orthorhombic (ao=5.5739, bo=5.4697, and co=15.745 angstrom (Å)), which can be represented as pseudo-tetragonal (at=3.905 and ct/4=3.9363 Å). The catalyst 20 is shown in FIG. 2 as strontium cobaltite (SrCoO3-δ) thin film in a perovskite phase. SrCoO3-δ is cubic with ac=3.8289 Å. Thin films of SrCoO2.5 and SrCoO3-δ may be grown epitaxially, for example by use of pulsed laser epitaxy on a variety of substrates. Examples of such substrates include, but are not limited to, perovskite (ABO3: where A is an alkaline or alkaline earth element, B is a transition metal and O is oxygen), SrTiO3, herein “STO” and (LaAlO3)0.3—(Sr2AlTaO6)0.7, herein “LSAT”. Thermo-mechanical degradation reduces the overall performance and lifetime of many perovskite oxides undergoing reversible redox reactions, such as those found in solid oxide fuel cells, rechargeable batteries, electrochemical sensors, oxygen membranes and catalytic converters. To mitigate this degradation, these reactions must occur at lower temperatures. However, high temperatures (>700° C.) are often required in conventional perovskites for fast catalysis and bulk diffusion. The strontium cobaltite oxygen sponge described herein can easily absorb and shed oxygen at as low a temperature as 200° C., which has been confirmed by switching between the crystalline phases of perovskite SrCoO3-δ and brownmillerite SrCoO2.5.
Directional terms, such as “vertical,” “horizontal,” “top,” “bottom,” “upper,” “lower,” “inner,” “inwardly,” “outer” and “outwardly,” are used to assist in describing the invention based on the orientation of the embodiments shown in the illustrations. The use of directional terms should not be interpreted to limit the invention to any specific orientation(s).
Referring to FIG. 3A, an x-ray diffraction (XRD) θ-2θ scan pattern of a brownmillerite-type SrCoO2.5 thin film grown on an STO substrate is shown. In FIG. 3B is shown an XRD θ-2θ scan pattern for perovskite-type SrCoO3-δ thin film grown on an LSAT substrate. While both phases may be grown epitaxially on either substrate STO (ac=3.905 Å) and LSAT (ac=3.868 Å) lower or fewer lattice mismatches are obtained when SrCoO2.5 films are grown on an STO substrate and when SrCoO3-δ films are grown on an LSAT substrate. The brownmillerite phase thin film may be readily stabilized in a simple molecular oxygen atmosphere while the perovskite phase thin film may require epitaxially stabilization in a more oxidizing atmosphere, for example, with the addition of mixed gases of ozone and oxygen.
Epitaxial SrCoO2.5 and SrCoO3-δ thin films (40-60 nm in thickness) were grown on (001) STO and (001) LSAT substrates by pulsed laser epitaxy (KrF, λ=248 nm). The films were grown at 750° C. in 0.013 mbar of O2 for the SrCoO2.5 and 0.267 mbar of O2+O3 (5%) for the SrCoO3-δ. The laser fluence was fixed at 1.7 J/cm2. The sample structure and crystallinity were characterized by high-resolution four-circle XRD (X'Pert, Panalytical Inc.). The Z-contrast images were obtained using a Nion Ultra STEM 200 operated at 200 keV.
FIGS. 3A and 3B reveal well defined peaks and distinct thickness fringes, which demonstrate the chemically sharp interface and flat surface. X-ray rocking curve scans further confirm excellent crystallinity (Δω<0.05°) and reciprocal space mapping confirmed that both films were coherently strained on the substrate as shown See FIGS. 4A through 4D. These rocking curves are from the 008 Bragg peak of SrCoO2.5 and the 002 Bragg peak of SrCoO3-δ. The curves reveal full width at half maxima (FWHM) of 0.04° and 0.05°, respectively. Typical FWHM in ω scan of the 002 STO peak is ˜0.015°, emphasizing the superior crystallinity confirmed by this analysis. Due to a larger mismatch between SrCoO3-δ and STO, the crystallinity of SrCoO3-δ is poorer than that of SrCoO3-δ on LSAT. In particular, the brownmillerite phase of strontium cobaltite exhibits a doubling of the c-axis lattice constant originating from the alternate stacking of octahedral and tetrahedral sub-layers along the c-axis (as shown in FIG. 1). FIGS. 4C and 4D are reciprocal space maps (“RSMs”) around the 103 STO with SrCoO2.5 (FIG. 4C) and SrCoO3-δ around 103 LSAT reflection (FIG. 4D), The in-plan lattice constants of both SrCoO2.5 and SrCoO3-δ films were coherently matched to those of the substrates. In addition, the c lattice constants from the RSMs were consistent with those measured from the θ-2θ scans, i.e. c/2=3.93 Å for SrCoO2.5 and c=3.79 Å for SrCoO3-δ, The c-axis lattice constant of SrCoO3-δ was smaller than the bulk value due to the substrate induced tensile strain.
Confirmation of the two chemically distinct phases was provided by observing the details of cobalt valence state of epitaxial SrCoO2.5 and SrCoO3-δ thin films by polarized x-ray absorption spectroscopy (XAS). This technique provides information on the oxidation state of Cobalt which plays a deterministic role in the magnetic, electronic and catalytic properties of the materials. Oxygen stoichiometry in the SrCoO2.5 and SrCoO3-δ thin films were qualitatively characterized by monitoring the Oxygen K-edge peaks as shown in FIG. 5A. The peak of SrCoO3-δ (solid line) at 527 eV clearly indicates different oxygen content as compared to SrCoO2.5. It may be expected that the pre-peak intensity of Oxygen K-edge decreases as δ approaches 0.5 in polycrystalline SrCoO3-δ. The weaker pre-peak may provide evidence that the SrCoO2.5 films were grown with the robust Co+3 valence state. The Oxygen K-edge pre-peak for the SrCoO3-δ film shows higher intensity compared to that for polycrystalline SrCoO2.82, which corresponds to the largest oxidation state reported in the literature, indicating that the analyzed SrCoO3-δ films are highly oxidized with a δ<0.18. Moreover, cobalt L-edge spectra may provide a rigorous determination of the cobalt valence state by measuring the empty cobalt 3d electronic state directly as shown in FIG. 5B (where data for SrCoO3-δ is shown by a solid line and data for SrCoO2.5 is shown by a dashed line). The shift in L3-edge toward the higher energy (>0.7 eV) in SrCoO3-δ indicates the cobalt ions in SrCoO3-δ are in a higher valence state. Together with the clear metallic behavior shown in FIG. 6C, discussed herein below, and the shift of the cobalt L2-edge peak from bulk SrCoO2.88, the highly oxygenated state of SrCoO3-δ is determined to be about δ≦0.1. Such highly oxygenated crystalline strontium cobaltite has not been grown previously without post treatment.
The distinct chemical valence difference between the SrCoO2.5 and SrCoO3-6 phases produced distinct magnetic and electric properties. Element-resolved measurements of the net magnetic moment using x-ray magnetic circular dichroism (XMCD) showed a large ferromagnetic signal in the SrCoO3-δ film, see FIG. 5C. The ferromagnetic state was further supported by the field-dependent XMCD data revealing that about 70% of the XMCD signal at 5 T was retained at 0.1 T. On the other hand, SrCoO2.5 displayed much less an XMCD signal even at 5 T, implying SrCoO2.5 is at a ferromagnetic ground state. As shown in FIGS. 6A and 6B, the magnetization measurement confirmed the XMCD results (SrCoO2.5 is represented by the solid lines and SrCoO3-δ is represented by the dotted lines). FIG. 6A represents the temperature dependent magnetization of SrCoO2.5 and SrCoO3-δ thin films at 1000 Oe. The SrCoO3-δ epitaxial film is ferromagnetic below ˜250K. The Curie temperature (Tc) is slightly lower than that from a single crystalline bulk SrCoO3 (˜305K). The lower Tc may be due to the substrate induced tensile strain. FIG. 6B shows the hysteresis loops for both SrCoO2.5 and SrCoO3-δ at 10K. The saturation magnetism of (Ms) of SrCoO3-δ at 10K was ˜2.3 μB/Co, slightly smaller than the bulk value with an intermediate spin state that can be attributed to tensile strain as well. These results provide experimental confirmation of the theoretically predicted spin state in SrCoO3-δ. In contrast to SrCoO3-δ, no significant XMCD or SQUID signal was observed from the SrCoO2.5 film. In addition, M(H) curves recorded at 10 and 250 K (data not shown) did not show a discernible difference, supporting the notion that the SrCoO2.5 epitaxial films described herein are antiferromagnetic.
There are also different electron transport properties between the two phases. Referring to FIG. 6C (where SrCoO2.5 is represented by the solid lines and SrCoO3-δ is represented by the dotted lines) there is shown a difference of more than four orders of magnitude in the resistivity between SrCoO2.5 and SrCoO3-δ at room temperature. An observation of metallic ground state in the epitaxial SrCoO3-δ film described herein represent the successful stabilization of Co+4 in the films, as the insulator-to-metal transition occurs at δ˜0.1. SrCoO3-δ shows a lower resistivity than that of polycrystalline samples which may indicate minimal oxygen deficiency and high quality form direct epitaxy. In contrast, the SrCoO2.5 film exhibited highly insulating properties. The calculated thermal activation energy of SrCoO2.5 was about 0.19 eV, which is similar to the value expected for bulk samples (0.24 eV). The difference in transport measurement is also shown in thermopower value (S) which makes it possible to distinguish electronic ground state. The thermopower (S) reflects the electronic ground state, i.e. insulator or metal, a comparison of the S-value of SrCoO2.5 with that of SrCoO3-δ at room temperature, as shown in FIG. 6D, reveals an S-value of SrCoO2.5 (S=+254 μVK) that is significantly greater than that of SrCoO3-δ (S=+9.6 μVK) and confirms the insulating property of SrCoO2.5. The observed S-values for both SrCoO3-δ and SrCoO2.5 are positive which indicates p-type conduction, consistent with bulk results. These results indicate the high sensitivity of electron transport properties as related to changes the oxygen content of strontium cobaltite.
The valence (or oxidation) state and magnetism in SCO was elucidated by XAS and XMCD at beamline 4-ID-C of the Advanced Photon Source, Argonne National Laboratory. Magnetic property was characterized with a 7 T Superconducting Quantum Interference Device (“SQUID”) magnetometer (Quantum Design). Temperature dependent DC transport measurements with van der Pauw geometry were performed with a 14 T Physical Property Measurement System (PPMS) (Quantum Design). The thermopower values were also measured by a conventional steady state method using two Peltier devices under the thin films to give a temperature difference (ΔV˜10K). Unlike the other measurements to eliminate the different substrate contributions during the thermopower measurements, both SrCoO3-δ and SrCoO2.5 films used were grown on (001) STO substrates.
Reversible redox reactions were monitored for a SrCoO3-δ-to-SrCoO2.5 conversion and SrCoO2.5-to-SrCoO3-δ conversion. These reactions were observed by XRD, in which several parameters may be controlled, such as but not limited to, gas type, flow rate and pressure. During the SrCoO3-δ-to-SrCoO2.5 conversion a high-resolution four-circle XRD (X'Pert, Panalytical Inc.) with a domed hot stage (DHS 900, Anton Paar) was used. The inside of the dome was evacuated with a mechanical pump to a base pressure of 0.0013 mbar. For the SrCoO2.5-to-SrCoO3-δ conversion, a power XRD with reactor chamber (XRK 900, Anton Paar) was used to pressurize the inside of the heating chamber to about 5 bar of O2. The temperature ramping rate was between about 30-60° C. per minute with an average scan time of about 2 to 2.5 minutes. To demonstrate the phase reversal processes, real-time readings of temperature dependent XRD θ-2θ scans with epitaxial films on LSAT in vacuum and with an oxygen atmosphere were recorded.
The ability to control the oxygen content in a material is significant for both physics and for technological application of multivalent oxides. Reversible oxidation-reduction (“redox”) reactions were directly observed between the epitaxially stabilized perovskite and brownmillerite phases of strontium cobaltite without destruction or degradation of the parent material. Results of direct probing of reversible redox activity are shown in FIGS. 7A and 7B. Real time temperature dependent XRD θ-2θ scans around the 002 LSAT reflection are shown. As shown in FIG. 7A, for the reduction process (i.e., SrCoO3-δ-to-SrCoO2.5), the 002 peak indicating the SrCoO3-δ film begins to disappear at about 175° C., and a complete transition to SrCoO2.5 is observed at about 210° C. The transition from one phase to another may be confirmed by the elongated c-axis lattice constant, as one can see in FIGS. #A and 3B, due to the oxygen vacancy ordering. The resulting c-axis orientation, i.e. oxygen vacancy channels aligned parallel to the interface in the oxygen reduced film on LSAT can be understood by the lower mismatch (lattice mismatch=0.96%) as compared to the a-orientation (lattice mismatch=1.76%) with the vacancy channels aligned vertically.
For the oxidation process (SrCoO2.5-to-SrCoO3-δ), complete oxidation may be achieved at about 350° C. in 5 bar of O2, as shown in FIG. 7B. In addition to the surprisingly low temperature for the topotactic phase transformation process (i.e., oxygen content change), the phase conversion was notably fast as the entire XRD scanning for the phase conversion took less than 10 minutes and is described herein above as between 2 and 2.5 minutes. Typically, conventional high pressure annealing processes at similar temperatures or by room temperature electrochemical approaches have taken >10 hours to complete. Moreover, the oxidation pressure is at least several hundred times lower than that of the annealing approach previously known. Phase conversion time and oxygen pressure for the epitaxial thin films were observed by the methods described herein as low as one minute and 0.67 bar, respectively. The lower limits of the pressure and temperature were determined by a step-by-step change of the post-annealing temperature subsequent to the growth of the thin film without breaking the vacuum.
The energy barrier for the phase transition of FIGS. 7A and 7B may be quantified by a computational thermodynamic approach. The calculated Gibbs free energy difference between the brownmillerite (SrCoO2.5) and the perovskite (SrCoO3-δ) at different oxygen contents (0<δ<0.3) as a function of temperature is represented in FIG. 8. Energy barriers of widely studied manganites, i.e. SrMnOX (“SMO”), which also forms SrMnO2.5 and SrMnO3-δ phases. The overall magnitude of the energy difference between SrCoO2.5 and SrCoO3-δ is smaller, in particular at low temperatures, than that of SMO at a given oxygen content. The energy difference between brownmillerite SrCoO2.5 and SrCoO3-δ decreases drastically, by at least about 30%, with a small deviation (e.g., δ=0.1) from perfect stoichiometry, while the maximum change in SMO is only about 20% at the same oxygen content. The reduced energy barrier for the phase transition between SrCoO2.5 and SrCoO3-δ with a small deviation from the stoichiometric SrCoO3 at low temperatures may be readily manipulated with non-thermodynamic factors, e.g., kinetics, induced strain, and surface-to-volume ratio. However, the energy difference between the SrMnO2.5 and SrMnO3-δ phases remained large regardless of the change in oxygen stoichiometry and temperature. These thermodynamic considerations provide considerable advantages of the strontium cobaltite over other conventional perovskites for rapid topotactic phase control. As can be observed in FIG. 8, the formation of SrCoO2.5 is more energetically favorable than SrCoO3-δ, whereas formation of SrMnO3-δ is more favorable than SrMnO2.5. Interestingly, the Gibbs free energy difference for highly oxygenated SrCoO3-δ (δ<0.3) increases as the temperature increases, while the opposite relationship exists for SMO. This observation of the thermodynamic phase stability of strontium cobaltite provides crucial information needed for understanding the topotactic processes in multivalent cobaltites.
Thermodynamic descriptions for SrCoO2.5 and SrCoO3-δ were taken from the thermodynamic modeling of strontium cobaltite for which model parameters were critically and self-consistently evaluated to reproduce both phase equilibrium and thermochemistry data in their bulk form. All the equilibrium phases in the thermodynamic modeling of strontium cobaltite other than SrCoO2.5 and SrCoO3-δ were suspended in the thermodynamic calculation to compute only the energy difference between SrCoO2.5 and SrCoO3-δ. Thermo-Calc software was used to minimize the individual Gibbs energies of the strontium cobaltite phases at given temperatures and oxygen contents. Thermodynamic descriptions for SrCoO2.5 and SrCoO3-δ were obtained from the 2010 Ph.D. thesis of J. E. Saal at the University of Pennsylvania. SrCoO2.5 and SrMnO2.5 were modeled as stoichiometric, while the perovskite phases were modeled as solution phases. The latter's designated oxygen sub-lattices allowed for mixing between oxygen and vacancy to achieve hypo-stoichiometry.
Fast, reversible redox activity is useful to provide catalytic activity at relatively low temperatures. Carbon monoxide oxidation probe reactions were conducted to elucidate the use of strontium cobaltite as a heterogeneous catalyst. Due to the extremely small surface areas of the epitaxial films, a custom designed micro-reactor was used with inlet gas streams of CO (0.1 mbar) and O2 (0.1 mbar). The effluent gas mixture was analyzed in-line with a gas-chromatograph and mass spectroscopy detector. For this reaction, the oxygen activity in the reactor was low; thus the epitaxial SrCoO3-δ film was unstable and the epitaxial SrCoO2.5 film on LSAT was chosen for the catalytic study. Catalytic activity can be connected to both the consumption of CO and the production of CO2. The conversion of the inlet CO gas is shown in FIG. 9A. As compared with the clean reactor, where there was no sample, clear conversion of CO was observed above about 320° C., increasing at a significant rate. At the same temperatures, an uptake in CO2 production was observed as shown in FIG. 9B. Because substantial activity can be measured at relatively low temperatures (above 320° C.) from a sample with a surface area of approximately 0.5 cm2, stabilized SrCoO2.5 is a good catalyst for many other redox reactions.
Gas-phase catalysis measurements were made with a custom micro-reactor with a volume <50 mL. A high level of reactor cleanliness was achieved by limiting material within the reactor to fused quartz, stainless steel and fluoropolymer seals. Heating was conducted by passing light from a halogen bulb through a fused quartz platform to the backside of the sample. The inlet gas streams consisted of 300 ppm CO and 300 ppm O2, both mixed with a helium balance (resulting in a partial pressure of 0.1 mbar for each gas), and a throughput of 5 SCCM. The temperature was programmed at 30° C. intervals each held for 16 minutes. At each interval the initial 10 minutes were reserved to allow the system to reach steady state conditions. Afterwards, a 20 μl aliquot of the gas stream was injected into the gas chromatograph mass spectrometer (PERKINS ELMER®). A carbon packed capillary column was used to separate the CO and residual N2 in the gas sample. The total ion chromatogram was mass separated to isolated peaks from CO and CO2 and then integrated. The concentration was determined with the integrated values and calibrated values from known gas mixtures. Detection limits for the CO conversion were set by the surfaces within the reactor, as determined by running the temperature program without a sample loaded and measuring the CO levels. In contrast, the detection limit for CO2 was set by the sensitivity of the mass spectrometer (˜4 ppb). The reproducibility of the conversion trends was substantiated with measurement of an additional SrCoO2.5 epitaxial thin film.
The above description is that of current embodiments of the invention. Various alterations and changes can be made without departing from the spirit and broader aspects of the invention as defined in the appended claims, which are to be interpreted in accordance with the principles of patent law including the doctrine of equivalents. This disclosure is presented for illustrative purposes and should not be interpreted as an exhaustive description of all embodiments of the invention or to limit the scope of the claims to the specific elements illustrated or described in connection with these embodiments. For example, and without limitation, any individual element(s) of the described invention may be replaced by alternative elements that provide substantially similar functionality or otherwise provide adequate operation. This includes, for example, presently known alternative elements, such as those that might be currently known to one skilled in the art, and alternative elements that may be developed in the future, such as those that one skilled in the art might, upon development, recognize as an alternative. Further, the disclosed embodiments include a plurality of features that are described in concert and that might cooperatively provide a collection of benefits. The present invention is not limited to only those embodiments that include all of these features or that provide all of the stated benefits, except to the extent otherwise expressly set forth in the issued claims. Any reference to claim elements in the singular, for example, using the articles “a,” “an,” “the” or “said,” is not to be construed as limiting the element to the singular.
Patent Application
20150148218
