Patent Application
20240034630
This disclosure relates to a method for the scalable production of diverse perimorphic materials, including stratified perimorphic materials comprising two or more perimorphic strata. More particularly, this disclosure relates to a low-cost, waste-reducing method for producing novel perimorphic materials wherein process materials may be recycled.
Compared to bulk materials, nanostructured materials may possess superior properties. Three-dimensional, ordered architectures constructed from nanostructured building blocks may facilitate the realization of these superior properties in bulk forms of the materials. These “architected” materials may be produced by synthesizing and arranging nanoscopic or microscopic building blocks into fine assemblies. In particular, porous materials with architected pore structures are appealing due to their low density, high specific surface area, and potential mechanical properties.
The '49195 Application teaches a scalable method (the “General Method”) of synthesizing carbon perimorphic frameworks using surface replication, or conformal replication of a templating surface, to direct the formation of the perimorphic architecture. By engineering the template materials, carbon frameworks possessing a variety of rationally engineered substructural and superstructural features have been synthesized. In some variants of the General Method, such as the “Preferred Method” taught in the '49195 Application, the template materials employed comprise magnesium oxide (MgO) templates derived from magnesium carbonate (MgCO3 xH2O) template precursor materials.
While the applications for carbonaceous perimorphic frameworks are numerous, it would be desirable to develop other perimorphic materials via template-directed surface replication procedures similar to those described in the '49195 Application. In particular, it would be desirable to synthesize perimorphic frameworks constructed from a range of materials that are stable in atomic monolayer or few-layer configurations. Examples of potentially useful framework compositions include sp2-hybridized boron nitride (BN), borophene (B), silicene (Si), boron carbonitride (BCxN), and various other ceramic compounds. In particular, it would be desirable to generate perimorphic frameworks comprising either electrically insulating or semiconducting elements or compounds, and likewise to generate these elements of compounds with the rational, architected morphologies that can be achieved via surface replication.
It would also be desirable to create three-dimensional frameworks comprising heterogeneous chemical compositions. In this way, different phases of the framework could fulfill different functions. As an example, carbonaceous frameworks might be shielded from thermal oxidation by sandwiching them between or encapsulating them within non-carbonaceous ceramic strata. As another example, carbonaceous frameworks might be coated by and serve as a functional support for a catalytic stratum. In particular, a perimorphic wall comprising multiple, distinct perimorphic strata would be desirable. Myriad useful heterostructured compositions have been identified in the graphene and graphene oxide literature, and it would be useful to be able to generate perimorphic frameworks from these diverse compositions, as well as new stratigraphically organized compositions that might be readily envisioned.
Additionally, it would be desirable to produce perimorphic materials, including perimorphic frameworks, using an approach that recycled both the template and the process liquids. Additionally, in certain variants, it would be desirable to recycle process gases. As described in the '49195 Application, conserving and resuing process materials can reduce the material inputs and outputs required for producing perimorphic materials, reducing cost and waste. If the General Method or the Preferred Method could be more generally employed to produce perimorphic materials that were not exclusively carbonaceous in composition, this would make the manufacture of these novel perimorphic materials more scalable and efficient.
The present disclosure demonstrates a method for synthesizing novel perimorphic materials of a number of chemical composition using surface replication techniques. The exemplary surface replication techniques demonstrated herein may be incorporated in the General Method. Therefore, these surface replication techniques expand the applicability of the General Method to include the scalable production of perimorphic frameworks of diverse chemical compositions. These novel perimorphic materials may be synthesized directly on MgO templates, or onto other perimorphic materials synthesized on MgO templates, and may be synthesized using the Preferred Method.
The present disclosure also demonstrates a method for synthesizing two-dimensional materials or arbitrary chemical composition directly on non-metallic templates, porous templates, and recyclable templates. In particular, a method is disclosed for synthesizing two-dimensional materials directly on thermally stable metal oxide compounds such as MgO, making it possible to engineer these two-dimensional materials in a variety of three-dimensional architectures. To demonstrate this, sp2-hybridized BN and BCxN perimorphic frameworks are synthesized via template-directed CVD on porous MgO templates. Analysis presented herein shows that these perimorphic frameworks comprises crosslinked, layered networks similar to the synthetic anthracitic carbon networks described in the '37435 Application. Other methods of deposition, including physical vapor deposition, and other gases may be used to deposit other two-dimensional materials on these highly stable, versatile templates. These other materials include monoelemental Xenes (such as borophene, silicene, germanene, stanene, phospherene, arsenene, antimonene, bismuthene, and tellurene) and compounds (such as various transition metal dischalcogenides), as well as doped variants.
The present disclosure also demonstrates a method for encapsulating a perimorphic framework by forming a gas-impermeable barrier phase around it. This barrier may be utilized to shield the encapsulated perimorphic framework from an external reactant, such as O2, or to seal the framework in a gas-evacuated internal state.
The present disclosure also demonstrates examples of novel perimorphic materials constructed from two-dimensional molecular structures such as sp2-hybridized BN and BCxN. These two-dimensional materials can therefore be fabricated into the same engineerable perimorphic architectures that have previously been demonstrated with graphenic carbon, including perimorphic architectures with controllably compact, ordered substructures and elongated, thin, equiaxed, hierarchical and hollow superstructures. Similar to the controllably flexible perimorphic frameworks that may be generated from graphenic carbon, controllably flexible perimorphic frameworks may also be generated from these other two-dimensional materials.
This disclosure also demonstrates examples of novel perimorphic materials constructed from two or more distinct perimorphic strata in order to obtain new functionality. These strata may comprise materials arranged in atomic monolayers, like graphenic carbon or sp2-hybridized BN, or materials with three-dimensional bonding structures, like silica or sp3-hybridized BN. The combination of carbon with other strata may also provide enhanced functionality. In particular, a practical example of shielding a graphenic perimorphic stratum from thermal oxidation via the addition of one or more thermally inert perimorphic strata is demonstrated in this disclosure. A supporting graphenic stratum may also be usefully combined with a catalytic stratum, such as a metallic or metal oxide adsorbate, to provide a high surface area catalyst, or to prevent charge carrier recombination (as in graphenic/TiO2 composites). Numerous applications for stratified perimorphic frameworks, which are analogous to three-dimensional architectures of two-dimensional/van der Waals heterostructures, will be readily apparent to those skilled in that art.
This disclosure also demonstrates examples of novel perimorphic materials that are either electrically insulating or semiconducting. Examples of electrically insulating perimorphic materials include silica-like and sp2-hybridized BN perimorphic frameworks. An example of a semiconducting perimorphic material is BCxN perimorphic frameworks, in which bandgaps can be varied by varying the carbon content.
This disclosure, in general, demonstrates examples of ceramics that are much lighter and varied in mechanical properties than their conventional bulk counterparts. In particular, these ceramics may be engineered to be flexible, and even crumpled or collapsed reversibly, like the carbon perimorphic frameworks that have previously been demonstrated. A wide range of ceramic alloys may be designed by performing thermal treatments of preceramics.
It is an object of this disclosure to render the previously disclosed General Method and Preferred Method more versatile. The ability to manufacture perimorphic materials with a variety of chemistries, while also conserving and reusing process materials, renders these methods and their variants more powerful and more broadly applicable. For example, there are a variety of mesoporous silicas made via sol-gel procedures using cetyl-trimethylammonium bromide (CTAB), n-dodecyl-trimethylammonium bromide (DTAB), or other consumable template materials; the present disclosure offers an alternative pathway that allows mesoporous silica-like materials to be made using recyclable template materials and process liquids. In particular, it is an object of this disclosure to provide a scalable means of creating three-dimensional, controllably compact architectures that are either noncarbonaceous or comprise multiple phases with distinct chemistries.
It is another object of this disclosure to demonstrate how a variety of elements and compounds, as well as stratified heterostructures of these elements and compounds, can be synthesized with templated, porous morphologies that have hitherto only been demonstrated for carbons. This includes mesoporous and macroporous substructures with engineerable cellular subunit geometries, as well as a variety of superstructural geometries. In particular, microscopic superstructures with elongated, flat, equiaxed, or hierarchical geometries can be synthesized.
It is another object of this disclosure to demonstrate how flexible perimorphic frameworks can be constructed from noncarbonaceous, two-dimensional materials. This is achieved via rational design of the superstructure and compactness of the perimorphic frameworks, as well as the thickness of the perimorphic walls.
It is another object of this disclosure to demonstrate how graphenic perimorphic frameworks can be modified by addition of other perimorphic strata. This may take the form of an overall encapsulation of the graphenic perimorphic framework, at the superstructural level, or a finer encapsulation of the perimorphic wall at the substructural level. These modifications can be utilized to shield graphenic carbons from thermal oxidation, enabling them to resist burning in high-temperature oxidizing environments, such as those that would be encountered in the presence or proximity of flames or at high velocities in the atmosphere.
It is another object of this disclosure to demonstrate how conventional and advanced ceramics, including ceramic alloys and composites, can be synthesized with engineered mesoporous or macroporous substructures and a variety of superstructural geometries. This may useful for applications that benefit from weight reduction, rapid mass transfer, high surface area, or increased toughness, as these properties may all be achievable via engineered porosity.
The Detailed Description of the present disclosure is organized according to the following sections:
Our objective in including References A through C in the present disclosure is to enable quick reference to the Detailed Descriptions of these related patent applications and to keep the exposition in the exposition of Sections I through IV as focused as possible.
Section V, or “Reference A,” is the Detailed Description from the specification of the '49195 Application, which teaches the scalable synthesis of carbonaceous perimorphic materials—a subset of the larger category of perimorphic materials. The Detailed Description included in Reference A has been modified in certain ways to harmonize its presence within the present disclosure and to avoid confusion that might otherwise arise from its inclusion. For example, where figures are cited in Reference A, their original numbering in the '49195 Application has been changed as needed to avoid redundancy. The designations of specific Furnace Schemes originally presented in the '49195 Application have been renamed in Reference A to avoid redundancy with the designated Furnace Schemes in other sections of the present disclosure. Measurements and data reported in Reference A were measured according to the analytical techniques specified in Reference A, not necessarily according to analytical techniques specified in other sections of the present disclosure. Lastly, sections in Reference A that were originally numbered with Roman numerals are designated with a single asterisk (e.g. what was originally designated Section I of the '49195 Application is designated Section I* in Reference A).
Section VI, or “Reference B,” is the Detailed Description from the specification of the '37435 Application, which teaches the synthesis of graphenic networks-a subset of the larger category of perimorphic materials. The Detailed Description included in Reference B has been modified in certain ways to harmonize its presence within the present disclosure and to avoid confusion that might otherwise arise from its inclusion. For example, where figures are cited in Reference B, their original numbering in the '37435 Application has been changed as needed to avoid redundancy. The designations of specific Furnace Schemes originally presented in the '37435 Application have been renamed in Reference B to avoid redundancy with the designated Furnace Schemes in other sections of the present disclosure. Measurements and data reported in Reference B were measured according to the analytical techniques specified in Reference B, not necessarily according to analytical techniques specified in other sections of the present disclosure. Lastly, sections in Reference B that were originally numbered with Roman numerals are designated with a double asterisk (e.g. what was originally designated Section I of the '37435 Application is designated Section I** in Reference B).
Section VII, or “Reference C,” is the Detailed Description from the specification of the '154 Application, which teaches surface replication on certain soluble templates-a subset of the larger category of templates that might be used according to the General Method. The Detailed Description included in Reference C has been modified in certain ways to harmonize its presence within the present disclosure and to avoid confusion that might otherwise arise from its inclusion. For example, where figures are cited in Reference C, their original numbering in the '154 Application has been changed as needed to avoid redundancy. The designations of specific Furnace Schemes originally presented in the '154 Application have been renamed in Reference C to avoid redundancy with the designated Furnace Schemes in other sections of the present disclosure. Measurements and data reported in Reference C were measured according to the analytical techniques specified in Reference C, not necessarily according to analytical techniques specified in other sections of the present disclosure. Lastly, sections in Reference C that were originally numbered with Roman numerals are designated with a triple asterisk (e.g. what was originally designated Section I of the '154 Application is designated Section I*** in Reference C).
We additionally note that the materials naming scheme that was utilized in the '49195 Application is adopted again in Sections I through IV of the present disclosure, and the numbering in the present disclosure is continued where the numbering in the '49195 Application left off. This has been done to facilitate easy reference to previously described materials and procedures and to avoid confusion. The continuity of the numbering does not imply that the present disclosure is a continuation or is supplemental to any previous disclosure.
The compatibility and combinability of many exemplary techniques, procedures, and materials set forth in Sections I through VII will be evident to knowledgeable practitioners of the art. For example, the Raman spectral features pertaining to an exemplary x-carbon demonstrated in Section VI might be readily obtained in a perimorphic framework comprising a hollow superstructure demonstrated in Section V. Taken together, Sections I through VII disclose a versatile and industrially scalable approach to producing architected, nanostructured materials, and they describe a diverse category of perimorphic materials.
In any cases where a contradiction may exist between the definition or description of a term or concept in References A through C and the definition or description of the same term or concept in Sections I through IV, the definition or description of that term or concept set forth in Sections I through IV should be understood as authoritative within the present disclosure. In any cases where a term or concept is defined or described in Sections A through C and is not contradicted by a corresponding definition or description in Sections I through IV, the definition or description set forth in Sections A through C should be understood as authoritative within the present disclosure.
A “stratified” perimorphic framework, as defined herein, comprises a multiphase framework in which the two or more distinct perimorphic strata can be identified within the perimorphic wall. When describing a stratigraphic pattern, the present disclosure describes the pattern with a string of letters in which each distinct stratum is represented by a letter, the position of a stratum in relation to other strata is represented by the position of its letter with respect to the other letters in the string, and compositionally similar strata are designated by the same letter. Hence, the string AB represents a perimorphic wall comprising two distinct and compositionally dissimilar strata, while the string BAB represents a perimorphic wall comprising three distinct strata, wherein one inner stratum is sandwiched between two outer strata, the outer strata being compositionally similar.
A “perimorphic stratum” (or “stratum”), as defined herein, comprises a distinct phase within a stratigraphically organized perimorphic wall. A perimorphic stratum typically shares a general alignment and topological similarity with the other perimorphic strata and with the perimorphic wall itself. As an example, a perimorphic wall might comprise a graphenic stratum positioned above or below a silica stratum. Even an all-carbon perimorphic wall may comprise distinct carbon strata, as described in the '580 Application.
“Pre-extraction replication,” as defined herein, comprises a surface replication technique that is performed prior to endomorphic extraction. Pre-extraction replication may be utilized to adsorb an adsorbate exclusively to one side of an existing perimorphic material, resulting in an A→*AB→*ABC buildup of the perimorphic wall (where A is synthesized on the templating surface, then B is synthesized on A, then C is synthesized on B, and then endomorphic extraction is performed).
“Post-extraction replication,” as defined herein, comprises a surface replication technique that is performed after endomorphic extraction. Post-extraction replication may be utilized to adsorb an adsorbate to both sides of an existing perimorphic material, resulting in an A→*BAB→*CBABC buildup of the perimorphic wall (where A is synthesized on the templating surface, then endomorphic extraction is performed, then the B strata are synthesized on both sides of A, and then the C strata are synthesized on both sides of the BAB stratigraphic arrangement).
Pre-extraction and post-extraction replication strategies can be combined. For example, an ABC stratification may be obtained via sequential pre-extraction replications. Subsequently, a post-extraction replication may be utilized to obtain a DABCD stratification.
“Stratigraphic occlusion,” as defined herein, comprises the use of one or more perimorphic strata to occlude another perimorphic stratum in a conformal configuration. One way to achieve stratigraphic occlusion is to use a post-extraction replication technique to obtain a BAB-type stratification, where a perimorphic stratum (A) is occluded via two conformally configured perimorphic strata (B).
“Stratigraphic encapsulation,” as defined herein, comprises the use of one or more perimorphic strata to encapsulate a perimorphic framework. One way to achieve stratigraphic encapsulation is to apply a perimorphic stratum around the periphery of an existing perimorphic framework.
A “two-dimensional” material, as defined herein, comprises a material with a bonding configuration that result in an atomic monolayer structure over small distances.
The “General Method” is the most basic form of the method and applies to the synthesis of perimorphic products of any chemical composition. It comprises a method for synthesizing a perimorphic product wherein substantial portions of the template material and the process liquid are conserved and may be reused. As such, the General Method may be performed cyclically. All variants of the method disclosed in the present disclosure comprise some variant of the General Method.
The General Method comprises a series of steps that is herein presented, for ease of description, in 4 stages (i.e. the Precursor Stage, Template Stage, Replication Stage, and Separation Stage). Each stage is defined according to one or more steps, as described below:
In practice, each step within these stages may itself comprise multiple, subsidiary steps. Additionally, each of the steps may occur concurrently with steps from another stage, such that in practice different stages may overlap in chronology. This can especially be expected in variants employing one-pot techniques. As a hypothetical example of this, a stock solution might be continuously sprayed alongside an adsorbate material into a furnace. In this hypothetical furnace, precursor particles might be precipited from the stock solution, template particles might be formed by heating of the precursor particles, and perimorphic material might be adsorbed to the template particles continuously and concurrently. This would correspond to steps assigned herein to the Precursor Stage, Template Stage, and Replication Stage, respectively.
Similarly, it is anticipated that in practice, many variants of the General Method may incorporate the steps described in the 4 Stages in different sequences. Also, in some variants, a step assigned by definition to one of the four stages herein might instead occur in a different stage. Such variants are anticipated herein and do not deviate from the inventive method, which is only presented herein as a discrete sequence of 4 stages for the sake of describing the overall cycle.
Ancillary processing steps (e.g. rinsing, drying, blending, condensing, spraying, agitating, etc.) may also be incorporated into the method at each stage. As a hypothetical example of this, a Replication Stage might involve coating a template material with a perimorphic material via a liquid-phase adsorption procedure, then filtering, rinsing and drying the resulting PC material. The incorporation of these processing steps in many variants will be obvious to those skilled in the art and, as such, they are not enumerated herein.
The inputs and outputs of the General Method are illustrated in
Variants of the General Method
The following discussion enumerates a number of ways in which the General Method may be variously implemented. The omission of variants from this discussion should not be interpreted as limiting, since an exhaustive list of ways in which the General Method may be implemented is not practical.
The General Method is intended to offer a means for cyclical production of perimorphic products while conserving process materials. In each cycle of the General Method, some portion of the process materials utilized are conserved and reused. In some variants, substantially all of the process materials utilized may be conserved and reused. In other variants, a portion of the process materials may be lost. One hypothetical example of this would be evaporative losses of process liquids from open tanks or wet filters.
In some variants of the General Method, process steps may correspond to batch processes. In other variants, process steps may correspond to continuous processes.
In some variants of the General Method, the solventless precipitation may comprise at least one of the following techniques: heating or cooling the stock solution to change the solubility of a solute in the stock solution; volatilizing a dissolved gas within the stock solution; depressurizing the stock solution; atomization of the stock solution; spray-drying the stock solution or spray pyrolysis.
In some variants of the General Method, a precursor structure may comprise at least one of the following: an elongated, thin, equiaxed, or hierarchical-equiaxed superstructure; an elongated superstructure with a length-to-diameter ratio greater than 200:1; an elongated superstructure with a length-to-diameter ratio between 50:1 and 200:1; a spheroidal or spherical superstructure; a hollow superstructure; a fragmentary superstructure comprising fragments of some other parent superstructure; a curved, fragmentary superstructure comprising fragments of a hollow superstructure.
In some variants of the General Method, a precursor structure may be precipitated around one or more other sacrificial structures, which may be present as inclusions in the precursor structure after its precipitation. In some variants, these inclusions in the precursor structure may be subsequently removed, resulting in voids.
In some variants of the General Method, a precursor structure may measure less than 1 μm along its major axis. In some variants, the precursor may measure between 1 μm and 100 μm along its major axis. In some variants, the precursor may measure between 100 μm and 1,000 μm along its major axis.
In some variants of the General Method, the precursor material may comprise at least one of the following: a hydrate; a metal hydroxide; a metal bicarbonate or carbonate; a Group I or Group II metal bicarbonate or carbonate; a mixture of salts. In some variants, the precursor may comprise MgCO3·xH2O in the form of at least one of: a hexahydrate, lansfordite, nesquehonite, hydromagnesite, dypingite, magnesite, and a nanocrystalline or amorphous structure.
In some variants of the General Method, the stock solution may comprise at least one of the following: metal cations and oxyanions; an aqueous metal bicarbonate solution; a Group I or Group II metal bicarbonate; an organic salt; Mg(HCO3)2. In some variants, the stock solution may comprise at least one of a dissolved gas, acid, and base. In some variants, the stock solution may be metastable.
In some variants of the General Method, the process liquid conserved in the Precursor Stage may comprise a distillate. In some variants, the distillate may be formed by condensing the process liquid vapor formed during spray-drying or spray-pyrolysis. In some variants, a process liquid conserved in the Precursor Stage may host solvated ions, the process liquid and ions together comprising a mother liquor.
In some variants of the General Method, the treatment performed on a precursor material in the Template Stage may comprise at least one of the following: heating the precursor, decomposing the precursor; partially or locally decomposing the precursor; decomposing the precursor surface; thermally decomposing the precursor; and oxidizing an organic phase present within the precursor structure. In some variants, the treatment may comprise at least one of flash-drying, spray-drying, spray pyrolysis, vacuum drying, rapid heating, slow heating, and sublimation. In some variants, a vapor released during the treatment may be conserved. In some variants, the vapor released may comprise at least one of CO2 and H2O. In some variants, treatment may comprise at least one of: coarsening the grain structure of the precursor or a decomposition product of the precursor; exposing to a reactive vapor; exposing to water vapor; sintering; and sintering with the assistance of dopants.
In some variants of the General Method, a template material may comprise at least one of the following: a metal hydroxide, a metal sulfate, a metal carbonate, a metal nitrate, a metal oxide, a Group I or II metal oxide, a transition metal, and MgO. In some variants, a template structure may comprise at least one of the following: macropores, mesopores, hierarchical porosity, subunits larger than 100 nm, subunits between 20 nm and 100 nm, and subunits between 1 nm and 20 nm.
In some variants of the General Method, a template structure may comprise at least one of the following: an elongated, thin, equiaxed, or hierarchical-equiaxed superstructure; an elongated superstructure with a length-to-diameter ratio greater than 200:1; an elongated superstructure with a length-to-diameter ratio between 50:1 and 200:1; a spheroidal or spherical superstructure; a hollow superstructure; a fragmentary superstructure comprising fragments of some other parent superstructure; and a curved, fragmentary superstructure comprising fragments of a hollow superstructure.
In some variants of the General Method, adsorbing the perimorphic material to the templating surface may comprise at least one of the following: liquid-phase application of the adsorbate, aerosolization of the adsorbate, physical vapor deposition, chemical vapor deposition, application of a liquid-state adsorbate, and application of a solid-state adsorbate. In some variants, vapor deposition of the adsorbate may comprise pyrolytic decomposition of a vapor at a temperature between 350° C. and 950° C.
In some variants, the adsorbate may comprise at least one of: an organic compound, a hydrocarbon compound, an organosilicon compound, an organometallic compound, a metalorganic compound, an organoboron compound, an organic nitrogen compound, a preceramic compound, a polymer, a graphenic network, a synthetic anthracitic network, an spx network, a helicoidal network, a carbonaceous material, an x-carbon, a z-carbon, a boron nitride, a boron carbonitride, an electrical conductor, an electrical insulator, and an electrical semiconductor. In some variants, the preceramic compound may comprise a silicon-bearing molecule, which may include at least one of polysiloxane, polysilsesquioxane, polycarbosiloxane, polycarbosilane, polysilylcarbodiimide, polysilsesquicarbodiimide, polysilsesquiazane, polysilazane, and metal-containing variants of these molecules, as well as others.
In some variants, the adsorbate may be altered after adsorption to the templating surface by at least one of the following processes: crystallization, sintering, grain growth, coalescence, decomposition, pyrolysis, polymerization, chemical functionalization, molecular grafting, chemical etching, activation, passivation, orbital rehybridization, maturation, and formation of a helicoidal network.
In some variants, the PC structure formed by adsorption of the adsorbate may comprise at least one of a single perimorphic phase, two or more distinct perimorphic phases, and two or more perimorphic phases arranged in distinct perimorphic strata. In some variants, the distinct strata may be applied via multiple, sequential surface replication procedures occurring prior to or following endomorphic extraction. In some variants, a perimorphic stratum may be sandwiched between two z-adjacent strata.
In some variants of the General Method, endomorphic extraction may utilize an extractant solution comprising a weak acid as an extractant. In some variants, an extractant solution may be formed by dissolving a process gas in process water. In some variants, the extractant solution may be an aqueous solution of H2CO3 formed by dissolving liquid or gaseous CO2 in process water. In some variants, endomorphic extraction may comprise a shuttling technique. In some variants, endomorphic extraction may be performed under conditions of elevated pressure or temperature.
In some variants of the General Method, the perimorphic separation may comprise at least one of decantation, hydrocyclones, settling, sedimentation, flotation, froth flotation, centrifugal separation, filtration, and liquid-liquid extraction. In some variants, the perimorphic separation may separate the perimorphic product from substantially all of the process liquid. In some variants, the perimorphic product may retain a residual portion of the process liquid. In some variants, the perimorphic product may be naturally buoyant due to its retention of internal gas. In some variants, the perimorphic product's internal gas may be expanded by reducing pressure of the surrounding process liquid, increasing the buoyancy of the perimorphic product and causing flotation. In some variants, a portion of the perimorphic product's internal gas may be exfiltrated by reducing pressure of the surrounding process liquid, followed by re-pressurizing the surround process liquid, such that hydrostatic pressure causes the process liquid to infiltrate the perimorphic product.
In some variants of the General Method, the stock solution generated by endomorphic extraction may be concentrated by dissolving additional solute(s) in the stock solution. In some variants, an additional solute may comprise at least one of a solid precipitated from stock solution and a process gas. In some variants, dissolution of the additional solute(s) in the stock solution may be dissolved by changing the temperature or pressure of the stock solution. In some variants, a concentrated aqueous Mg(HCO3)2 stock solution may be formed by precipitating a dilute stock solution to form an MgCO3
In some variants of the General Method, the perimorphic product may comprise a perimorphic framework. In some variants, the morphology of the perimorphic framework may comprise at least one of a native morphology, a non-native morphology, a crumpled morphology, a hollow morphology, a hierarchical morphology, macropores, mesopores, micropores, a spheroidal superstructural geometry, a prismatic superstructural geometry, a shell, a shell fragment, a noncellular space, and a labyrinthine pore structure.
In some variants, the perimorphic framework may comprise at least one of a hydrophobic material, a hydrophilic material, an amphiphilic material, and a Janus material comprising hydrophobic and hydrophilic surfaces. In some variants, the perimorphic framework may comprise at least one of a flexible, rigid, and elastic. In some variants, the perimorphic framework may contain an internal gas and may float when immersed in a liquid.
In some variants, the major axis of the perimorphic framework may measure at least one of less than 1 μm, between 1 μm and 100 μm, and between 100 μm and 1,000 μm. In some variants, the framework may comprise a BET surface area of at least one of 1,500 to 3,000 m2/g and between 10 to 1,500 m2/g. In some variants, the framework may comprise an elongated, thin, or equiaxed superstructure. In some variants, an elongated framework may comprise a length-to-diameter ratio between 50:1 and 200:1.
In some variants of the General Method, the perimorphic framework may comprise a carbonaceous phase comprising at least one of a carbonaceous material, a pyrolytic carbon, a graphenic network of carbon, n anthracitic network of carbon, an spx network of carbon, and a helicoidal network of carbon, an x-carbon, and a z-carbon. In some variants, the perimorphic framework may comprise functional groups including at least one of a carbon atom, an oxygen atom, a halogen atom, a metal atom, a boron atom, a sulfur atom, a phosphorus atom, and a nitrogen atom.
In some variants of the General Method, under 532 nm excitation, a carbonaceous phase of a perimorphic framework may comprise at least one of a Raman spectral ID/IG ratio of between 4.0 and 1.5; a Raman spectral ID/IG ratio between 1.5 and 1.0; a Raman spectral ID/IG ratio between 1.0 and 0.1; a Raman spectral ITr/IG ratio between 0.0 and 0.1; a Raman spectral ITr/IG ratio between 0.1 and 0.5; a Raman spectral ITr/IG ratio between 0.5 and 1.0; a Raman spectral I2D/IG ratio between 0 and 0.15; a Raman spectral I2D/IG ratio between 0.15 and 0.3; and a Raman spectral I2D/IG ratio between 0.30 and 2.0.
In some variants of the General Method, under 532 nm excitation, a carbonaceous phase of a perimorphic framework may comprise at least one of an unfitted Raman spectral D peak positioned between 1345 and 1375 cm−1; an unfitted Raman spectral D peak positioned between 1332 and 1345 cm−1; an unfitted Raman spectral D peak positioned between 1300 and 1332 cm−1; an unfitted Raman spectral G peak positioned between 1520 cm−1 and 1585 cm−1; an unfitted Raman spectral G peak positioned between 1585 cm−1 and 1600 cm−1; and an unfitted Raman spectral G peak positioned between 1600 cm−1 and 1615 cm−1.
In some variants of the General Method, the perimorphic framework may comprise a non-carbonaceous phase comprising a ceramic. In some variants, the ceramic phase may comprise at least one of: one or more post-transition metals, one or more metalloids, one or more reactive nonmetals, and a decomposition product of one or more preceramics. In some variants, the ceramic phase may comprise at least one of the following: silicon oxycarbide (Si—O—C), silicon carbide (Si—C), silicon nitride (Si—N), silicon boride (Si—B), silicon carbonitride (Si—C—N), and silicon boron carbonitride (Si—B—C—N), as well as metal-modified and various stoichiometric compositions of these.
In some variants, the ceramic phase of a perimorphic framework may comprise a nanostructured BN phase comprising at least one of sp2-hybridized states, sp3-hybridized states, a mixture of sp2- and sp3-hybridized states, a layered architecture, and structural dislocations that provide internal crosslinking between layers. In some variants, the BN phase of a perimorphic framework may comprise a synthetic anthracitic network, an spx network, and a helicoidal network.
In some variants, the substantially sp2-hybridized BN phase of a perimorphic framework may comprise one or more atomic monolayers. In some variants, two or more atomic BN monolayers may exhibit nematic alignment. In some variants, the BN phase, under 532 nm excitation, may comprise a single broad, unfitted Raman spectral band between 500 and 2500 cm−1. In some variants, the peak position of this band may be located between in at least one of the following ranges: 1300 cm−1 to 1400 cm−1, 1400 cm−1 to 1500 cm−1, and 1500 cm−1 to 1600 cm−1.
In some variants, the ceramic phase of a perimorphic framework may comprise a nanostructured BCxN phase comprising at least one of sp2-hybridized states, sp3-hybridized states, a mixture of sp2- and sp3-hybridized states, a layered architecture, and structural dislocations that provide internal crosslinking between layers. In some variants, the BCxN phase of a perimorphic framework may comprise a synthetic anthracitic network, an spx network, and a helicoidal network. In some cases, the BCxN phase of a perimorphic framework may comprise an engineerable electronic bandgap based on its fractional composition of carbon.
In some variants, the substantially sp2-hybridized BCxN phase may comprise one or more atomic monolayers. In some variants, two or more atomic BCxN monolayers may exhibit nematic alignment. In some variants, the BCxN phase, under 532 nm excitation, may comprise at least one of a G peak positioned between 1500 cm−1 and 1650 cm−1, a broad, unfitted Raman spectral band between 500 cm−1 and 2500 cm−1, a G peak associated with sp2 carbon and an underlying broad band associated with BN, and a substantially absent D peak associated with sp2 carbon rings.
In some variants, the nanostructured ceramic phase may comprise at least one of the following monoelemental atomic monolayers: borophene, silicene, germanene, stanene, phospherene, arsenene, antimonene, bismuthene, and tellurene. In some variants, the nanostructured ceramic phase may comprise substantially two-dimensional transition metal dischalcogenides.
In some variants, the nanostructured ceramic phase may comprise a metal oxide, or the oxide of a metalloid or reactive nonmetal. In some variants, the metal oxide may comprise a layered transition metal oxide. In some variants, the metal oxide may comprise a mixed metal oxide. In some variants, the metal oxide may comprise at least one of a catalyst and a photocatalyst.
In some variants, the perimorphic wall may comprise a nanostructured metallic phase. In some variants, the nanostructured metallic phase may comprise at least one of a Group I metal, a Group II metal, a transition metal, a transition metal alloy, Ni, Ni—Mo, a reduced decomposition product of a metallocene, an electroless coating, and a catalyst.
In some variants, the perimorphic wall may comprise two or more nanostructured phases. In some variants, the two or more phases may comprise distinct perimorphic strata. In some variants, electrically insulating, conducting, or semiconducting perimorphic strata may be alternated. In some variants, a perimorphic stratum may be sandwiched between two other perimorphic strata to shield it. In some variants, a carbonaceous perimorphic stratum may be shielded from thermal oxidation by one or more other perimorphic strata.
In some variants of the General Method, the perimorphic product may be subjected to further treatment after perimorphic separation. In some variants, the further treatment after perimorphic separation may comprise at least one of flash-drying, spray-drying, spray-pyrolysis, decomposition, chemical reaction, annealing, sintering, and chemical functionalization.
In some variants of the General Method, the Liquid Cycle may also incorporate the recapture and conservation of process liquid released or evaporated during the Template Stage, although this is not reflected as an output in
In some variants of the General Method, a Gas Cycle may be incorporated into the method. The inputs and outputs of the General Method with a Gas Cycle are illustrated in
The Preferred Method, described below, comprises variants of the General Method in which a MgCO3·xH2O template precursor material is derived from an aqueous Mg(HCO3)2 stock solution and a portion of the CO2 process gas is conserved via a Gas Cycle. The inputs and outputs of the Preferred Method are shown in
Certain variants of the Preferred Method may employ pressure modulations in order to form concentrated stock solutions and improve precipitation processes. Concentrated stock solutions may be associated with many benefits, including superior precipitation kinetics, reduced process water volumes, smaller vessels, and improved energy efficiency. Two exemplary ways that this can be done are illustrated in
In the first frame of
Another way that a concentrated stock solution may be obtained is by performing the endomorphic extraction in a pressurized reactor. A schematic showing this is illustrated in
In the course of describing procedures to generate the exemplary materials described in the subsequent sections, certain furnace schemes have been detailed. These schemes may be used for the exemplary Template Stage procedures detailed in Section V and for the exemplary Replication Stage procedures detailed in Section VI.
Scheme A: In Scheme A, a Thermcraft tube furnace modified to be a rotary furnace may be employed with a quartz tube. The furnace has a clam shell design with a cylindrical heating chamber of 160 mm diameter and 610 mm heated length. The furnace has a wattage of 6800 W with a maximum operating temperature of 1100° C. The quartz tube may be a 60 mm OD quartz tube containing an expanded middle section of 130 mm OD tube (the “belly”) positioned within the furnace's heating zone. The tube may be rotated. Quartz baffles inside the belly may facilitate agitation of the a powder sample during rotation. The furnace may be kept level (i.e. not tilted). The template powder sample may be placed inside the belly in the heating zone, with ceramic blocks inserted outside the belly on each side of the furnace's heating zone. Glass wool may be used to fix the position of the ceramic blocks.
For exemplary procedures performed using Scheme A, a material sample may be placed inside the belly, such that it agitated within the reactor. Loose fitting ceramic blocks located outside the belly section on each side of the furnace's heating zone allowing for gas flow and powder containment. Packed glass wool may be used to affix the position of the ceramic blocks while acting as a gas permeable layer. The ends of the tube may be fitted with two stainless-steel flanges to allow for gas to flow for the system.
Scheme B: An MTI rotary tube furnace with a quartz tube may be used. The furnace has a clam shell design with a cylindrical heated chamber having dimensions of 120 mm diameter and 440 mm heated length. The furnace has a wattage of 2500 W with a maximum operating temperature of 1150° C. The quartz tube may be 60 mm in OD. The tube may be substantially level. For exemplary procedures performed using Scheme B, a material sample may be placed within a ceramic boat. This may then be placed inside the quartz tube within the heating zone prior to the initialization of heating. Loose fitting ceramic blocks located outside the furnace's heating zone allow for gas flow. Packed glass wool may be used to affix the position of the ceramic blocks while acting as a gas permeable layer. The ends of the tube may be fitted with two stainless-steel flanges. If ammonia borane (H3NBH3) is used, the solid H3NBH3 may be placed in a ceramic boat just outside the upstream side of the furnace heating zone, enabling it to reach a temperature between 130° C. and 170° C. when the furnace reaches the set temperature.
Scheme C: A Lindberg Blue-M tube furnace with a quartz tube may be used. The quartz tube may be 150 mm in OD. The furnace has a clam shell design with a cylindrical heated chamber having dimensions of 190 mm diameter and 890 mm heated length. The furnace has a wattage of 11,200 W with a maximum operating temperature of 1200° C. The tube may be substantially level. For exemplary procedures performed using Scheme C, a sample may be placed within a ceramic boat. This may then be placed inside the quartz tube within the heating zone prior to the initialization of heating. Loose fitting ceramic blocks located outside the furnace's heating zone allow for gas flow. The ends of the tube may be fitted with two aluminum flanges to allow for gas flow through the system.
Scheme D: A Vulcan 3-550 Muffle furnace may be used. The furnace has a rectangular heated chamber having dimensions of 190 mm×240 mm×228 mm. The furnace has a wattage of 1440 W with a maximum operating temperature of 1100° C. For exemplary procedures performed using Scheme D, a material sample may be placed within a ceramic boat. This may then be placed inside the muffle furnace prior to the initialization of heating.
Scheme E: A TA Instruments Q600 TGA/DSC may be used. For exemplary procedures performed using Scheme E, a 90 μL alumina pan may be used to hold a material sample. Gas flow may be 100 sccm of a specified gas unless otherwise noted. The heating rate may be mentioned in the exemplary procedures where Scheme E is used.
A number of analytical techniques were utilized to characterize the procedures and materials presented herein. These are detailed below.
Solution concentrations were measured using electrolytic conductivity (“conductivity”). The conductivity is a measured response of a solution's electrical conductance. The electrical response of a solution may be correlated to the concentration of ions dissolved in the solution, and as ions in solution are precipitated, the conductivity value decreases. An analog to this measurement is total dissolved solids (“TDS”) which relates the conductivity measurement to a referenced ion concentration (typically potassium chloride), dependent on the salt compound dissolved.
Raman spectroscopy was performed using a ThermoFisher DXR Raman microscope equipped with a 532 nm excitation laser. For each sample analyzed, 16 point spectra were generated using measurements taken over a 4×4 point rectangular grid with point-to-point intervals of 5 μm. The 16 point spectra were then averaged to create an average spectrum. The Raman peak intensity ratios and Raman peak positions reported for each sample all derive from the sample's average spectrum. No profile fitting software was utilized, so the reported peak intensity ratios and peak positions relate to the unfitted peaks pertaining to the overall Raman profile.
Gas adsorption measurements were made using a Micromeritics Tristar II Plus. Nitrogen adsorption was measured at a temperature of 77 K across a range of pressure (p) values, where
Increments of pressure ranged from
Micromeritics MicroActive software was used to calculate the BET specific surface area derived from the BET monolayer capacity assuming the cross-sectional area σm(N2, 77 K)=0.162 nm2. Samples were preconditioned by degassing with continuously flowing dry nitrogen gas at 100° C. prior to analysis.
The pore size distribution (PSD) and cumulative volume of pores is another technique that may be performed from gas adsorption data to lend insight into the sintering behavior of particles. The data was collected by a Micromeritics Tristar II Plus, measuring nitrogen adsorption and desorption at 77 K between pressures of
with increments ranging from
Samples were preconditioned by degassing with continuously flowing dry nitrogen gas at 100° C. prior to analysis.
Micromeritics MicroActive software was used to calculate adsorption-desorption PSD and cumulative volume of pores by applying the Barrett, Joyner and Halenda (BJH) method. This method provides a comparative assessment of mesopore size distributions for gas adsorption data. For all BJH data, the Faas correction and Harkins and Jura thickness curve may be applied. The cumulative volume of pores, VPORE (cm3/g), may be measured for both adsorption and desorption portions of the isotherm.
There are a number of exemplary materials described in the present disclosure. To aid in identification and tracking of these exemplary materials, a material naming system has been adopted and is described below. All names of exemplary materials are bolded; N2 herein describes an exemplary material, while N2 refers to nitrogen gas.
Exemplary types of template precursor materials are denoted Sx, where S designates the first one or two letters of the template precursor (i.e. N for nesquehonite, L for lansfordite, Li for lithium carbonate, C for magnesium citrate, A for amorphous MgCO3·xH2O, H for hydromagnesite, M for magnesite, E for epsomite, and Ca for calcium carbonate) and where x designates different types of the precursor compound (e.g. H1 and H2 designate two different types of hydromagnesite precursors).
Exemplary types of template materials are named in the format SxTy. The Sx name component designates the precursor type that was utilized to create the template type SxTy, and the Ty name component designates a specific treatment that was utilized to create the template type SxTy. For example, N1T1 and N1T2 indicate two different template types formed from two different treatments on the precursor type N1. We note that while the full SxTy name denotes a specific template type, the Ty name component by itself is only specific with respect to a given Sx precursor type. For example, the treatments utilized to make the template types N1T1 and N2T1 were different, despite these template types sharing the same T1 name component.
Exemplary types of PC materials are named in the format SxTyPZ, where the SxTy name component designates the template type and the Pz name component designates a specific type of perimorphic material. For example, M3T1P1 and M3T1P2 indicate two different PC materials formed from the same M3T1 template material. The Pz name component within the SxTyPZ name is unique—i.e. each Pz name component specifies a unique type of perimorphic material, irrespective of the SxTy template type utilized to make the perimorph.
Exemplary types of perimorphic frameworks (i.e. the porous perimorphic product resulting from endomorphic extraction) are named in the format Pz, where the Pz name component is not prefaced with an SxTy template type. The Pz name component utilized by itself to designate a framework type matches the Pz name component of the SxTyPz PC material from which the framework type was derived.
The exemplary types of template precursor materials, template materials, PC materials, and perimorphic materials in this disclosure are enumerated in
This section details the generation of exemplary perimorphic materials at small scales using exemplary procedures. A full implementation of the General Method is not described in connection with every exemplary procedure, although such an implementation would be possible with every exemplary procedure. Additionally, it should be understood that many techniques or materials utilized in these exemplary procedures are merely intended to demonstrate the effects or properties of techniques or materials that might be utilized in larger-scale industrial implementations.
The '49195 Application teaches the synthesis of a number of exemplary types of template precursor materials, template materials, PC materials, and perimorphic frameworks. Taken together, these materials demonstrate the breadth of carbonaceous perimorphic products that can be synthesized, as well as the breadth of template precursor materials and template materials that may be utilized. In this section, we demonstrate how the General Method can also be used to create stratified perimorphic frameworks and perimorphic frameworks that are not strictly carbonaceous. The procedures and materials presented in this section are meant to be exemplary, since a wide variety of procedures and materials may be readily envisioned and used without deviating from the invention. The types of perimorphic materials synthesized in the below examples are summarized in
Examples P24, P25: In two exemplary procedures, a stratified perimorphic material (P24) and a silica-like perimorphic material (P25) may be synthesized, depending on the choice of atmosphere during a final thermal treatment performed after surface replication.
For the purpose of demonstration, a P23-type carbonaceous perimorphic material may first be synthesized via surface replication on porous, ex-magnesite MgO template structures in a way consistent with the General Method (and the Preferred Method). This Precursor Stage and Template Stage procedures are described in the '49195 Application and Reference A, and the surface replication parameters may be found in
Next, the P23-type carbon may be chemically functionalized. To do this, an aqueous paste containing a P23-type carbon content of approximately 13% by weight may be made. A 144 g quantity of the paste, containing approximately 18.7 g of carbon, may be added to 500 g deionized water in a beaker and stirred using an overhead Cowles blade mixer to suspend the carbon. This mixture may then be transferred from the beaker into the reservoir of a high-shear rotor-stator homogenization processor (IKA Magic Lab, or “ML”). The mixture in the reservoir may be mixed using an overhead Cowles blade mixer to keep the particles adequately suspended in the reservoir. Residuals in the beaker may be rinsed with 50 g deionized water, and the residuals and rinsate may be added to the ML reservoir. The external thermal control system on the ML may be used to maintain the mixture at a temperature of 5° C., and the rotor-stator speed may be set to 15,000 RPM. Using these settings, the mixture may be circulated for 30 minutes, maintaining the mixture temperature at 5° C. The mixture may have a pH of approximately 10. At this point, 2.5 g of aqueous HCl may be added over 15 seconds. Then, 15.5 g of aqueous NaOCl (14.5% concentration) may be added over 15 seconds. The mixture may have a pH of approximately 2.35. The mixture may be run for an additional 15 minutes at a temperature of 5° C. At this point, 2 g of aqueous H2O2 (35% concentration) may be added. Next, the mixture may be removed from the ML. Residuals in the ML may be rinsed with deionized water, and the residuals and rinsate may be added to the mixture. The mixture may then be filtered, rinsed with deionized water, then rinsed with ethanol, resulting in an ethanol paste of oxidized carbon perimorphic material.
Next, a 92 g quantity of the ethanol paste containing a carbon content of approximately 15.9 g may be diluted with 400 g of ethanol in a beaker and stirred using an overhead Cowles blade mixer to suspend the carbon. This mixture may then be transferred from the beaker into the ML reservoir. The mixture in the reservoir may be mixed using an overhead Cowles blade mixer to keep the particles adequately suspended in the reservoir. The rotor-stator speed may be set to 15,000 RPM and the mixture may be allowed to remain at approximately room temperature. A 15.8 g quantity of 3-[2-(2-aminoethyl)amino] propyl trimethoxysilane (AEAPTMS) may be added to the reservoir over 1 minute. This may be followed by the addition of 95 g of deionized water and 1.6 g of NaOH, bringing the pH to approximately 10.6. The mixture may be circulated for 30 min, then removed from the ML and transferred to a beaker. Residuals in the ML may be rinsed with 100 g deionized water, then 50 g of ethanol, and the residuals and rinsate may be added to the mixture. The beaker may be magnetically stirred and heated for the next 150 minutes, its temperature ranging from approximately 78° C. to 93° C. During this time, the inner sides of the beaker may be rinsed twice, using 50 g of deionized water each time and raising the mixture's boiling point. At this point, the heating may be turned off. The mixture may then be filtered and rinsed with ethanol, resulting in an ethanol paste. The paste may be dried at 60° C. to form a powder, the frameworks comprising AEAPTMS-functionalized carbon.
Next, the powder of AEAPTMS-functionalized carbon may be subjected to a post-replication thermal treatment. In both Example P24 and Example P25, the treatment may be performed in a TGA instrument, as described in Furnace Scheme E detailed in Section III. In Example P24, the treatment may be performed under flowing Ar, while in Example P25 the treatment may be performed under flowing air. In each treatment, the powder sample may be heated from room temperature to a final temperature of 900° C. at a heating rate of 20° C./min. Upon reaching 900° C., the sample may be cooled back down to room temperature. The oxidizing atmosphere of the thermal treatment utilized in Example P25 causes the carbon perimorphic stratum and the organic phase of the polysiloxane to be completely oxidized, resulting in a silica-like perimorphic material, which is the brownish-white powder shown in
An SEM micrograph of silica-like P25-type frameworks are shown in
N2 desorption analysis shows that the silica-like P25-type frameworks possess a non-native cellular substructure. In
The N2 adsorption analysis also reveals that the silica-like frameworks have an average surface area of 1273 m2/g. This is a considerably higher than the P23-type carbon frameworks' average surface area of 461 m2/g and the AEAPTMS-functionalized frameworks' average surface area of 463 m2/g. This reflects the elimination of the carbonaceous perimorphic stratum in Example P25, in which an oxidizing thermal treatment was employed. The silica-like perimorphic stratum remaining after the thermal treatment is thinner than the eliminated carbon stratum.
Similar procedures may be used to create stratified or silica-like perimorphic frameworks with other engineered features. Other exemplary templates and procedures described in References A and B may be readily utilized in concert with the approach described in Example P24 and Example P25. Also, similar procedures may be used to obtain stratigraphic encapsulation of a perimorphic material. One way of obtaining this result is to form a preceramic stratum such as an inorganic polymer on an existing perimorphic stratum, to pyrolyze the preceramic stratum in order to form a ceramic stratum, and then to sinter or melt the ceramic stratum, such that a continuous ceramic phase is formed around the underlying perimorphic material. Stratigraphic encapsulation may be used, for example, to shield a carbonaceous perimorphic framework from oxidation in high-temperature oxidizing environments.
As an example, elongated perimorphic materials are shown in
Based on this demonstrated impermeability to O2 gas, we can conclude that if pyrolysis and sintering had been performed in a vacuum, the framework would have been encapsulated in an internally evacuated state, and would then have been sealed with respect to surrounding air. In certain perimorphic architectures, especially hierarchical superstructures with large central cavities, where the mass of an encapsulated gas becomes significant in relation to the mass of the framework, the framework's apparent density may be reduced by evacuating the gas from the framework's internal pores, then encapsulating the entire framework in this state of vacuum or partial vacuum. Hence, obtaining encapsulation of the perimorphic material in a vacuum may be useful.
Perimorphic frameworks and strata with a variety of morphologies and polymer-derived ceramic compositions may be readily obtained using pathways similar to those described in Example P24 and Example P25.
Examples P26, P27: In another exemplary procedure, boron nitride (BN) perimorphic materials and stratified perimorphic materials comprising BN and carbon strata may be synthesized.
For the purpose of demonstration, a P7-type carbonaceous perimorphic material may first be synthesized via surface replication on porous, ex-hydromagnesite MgO template structures in a way consistent with the General Method (and the Preferred Method). The Precursor Stage and Template Stage procedures are described in the '49195 Application and Reference A, and the surface replication parameters may be found in
Next, BN perimorphic strata may be adsorbed to each side of the P7-type carbon perimorphic frameworks, creating a BAB stratigraphic arrangement. To do this, a 40 mg quantity of P7-type carbon frameworks may be placed in a ceramic boat, which may be placed in a tube furnace, according to Scheme B, as detailed in Section III. A ceramic boat containing 2 g of ammonia borane complex (H3NBH3) may be placed in the quartz tube just outside the furnace's heated zone, such that when the furnace reaches a temperature of 700° C., the H3NBH3 reaches a temperature between 130° C. and 170° C. The furnace may then be heated to a temperature of 700° C. at a heating rate of 20° C./min under Ar flowing at 2000 sccm. Upon reaching 700° C., the furnace may be maintained at 700° C. for 60 minutes, then allowed to cool to room temperature under continued Ar flow.
Next, the powder may be placed in a muffle furnace, according to Scheme D, as detailed in Section III. The furnace may be heated to a temperature setting of 800° C. under air and then held at this temperature for 1 hour, then allowed to cool to room temperature.
The resulting powder may comprise two phases. The first phase comprises stratified perimorphic frameworks comprising a BAB stratigraphic arrangement, where B represents the outer strata of BN and A represents the inner stratum of carbon, and this phase comprises a type of perimorphic material identified herein as P26. The P26-type phase is optically black, as shown in the optical micrographs in
This stratigraphic arrangement is further confirmed via Raman spectroscopy. Each spectrum in
Overall, comparing the lineshape of this P26 spectrum to that of the P7-type carbon alone (
The formation of the distinct P26 and P27 powder phases is due to the static-bed CVD procedure used to grow the BN in Examples P26 and P27. With no agitation to facilitate solid-gas mixing, the carbon frameworks near the surface of the static bed were substantially covered with the BN adsorbate on both sides of the perimorphic wall during surface replication. This resulted in the stratigraphic occlusion of the carbon, which was sandwiched between two BN strata, and the formation of a stratified perimorphic material near the surface of the static bed. However, the carbonaceous perimorphic frameworks farther from the surface of the static bed were incompletely covered with the BN adsorbate during surface replication due to gas diffusion constraints. These unoccluded carbonaceous frameworks were then completely thermally oxidized during the 800° C. treatment.
Because of the ability to grow BN on carbon substrates, and vice-versa, a multistage replication procedure can be utilized to create various stratigraphic arrangements of BN and carbon. For example, if the BN growth procedure were performed in a pre-extraction replication procedure (i.e. prior to endomorphic extraction), the resulting stratigraphic arrangement would have been AB, as opposed to BAB. Another pre-extraction replication step growing carbon on the BN stratum could have been performed to create an ABA stratigraphic arrangement. Any number of steps, chemistries, and stratigraphic arrangements are possible, and it may be useful to create alternating electrically conducting, semiconducting and insulating strata.
Perimorphic frameworks with various chemical compositions and phases may be of interest for weight-sensitive ceramic applications, ceramic applications in which unusual mechanical properties, such as flexibility or pseudoelasticity, are desired, and ceramic applications in which high thermal stability or thermal shock resistance is desired. Retention of a carbon stratum within the perimorphic wall may be desired not only for its own functionality in applications, but also because it may stabilize the perimorphic architecture of other ceramic strata during high-temperature production and service. Exposed carbon strata may be easily chemically functionalized, whereas certain ceramics may be more difficult to functionalize, so carbon strata may be used for functionalization purposes, also. This is similar to the concept presented in the '580 Application, in which a disordered, easily oxidized perimorphic stratum or “skin” is formed over a less disordered, graphitic perimorphic stratum that is not as easily oxidized.
Other perimorphic compositions that will be desirable include transition metal dichalcogenides (“TMDCs”) and stratified perimorphic materials including multiple TMDCs or carbon and TMDCs. Also, stratified compositions involving carbon and metal oxides such as TiO2 would be desirable for a number of applications, such as photoanodes. The General Method may be used to generate these compositions in the form of controllably compact, perimorphic frameworks with engineered superstructural and substructural architectures. Hence, the method is not limited to carbon perimorphic frameworks, or even single-phase frameworks, but may be used to synthesize perimorphic materials comprising diverse chemistries and combinations of chemistries. It may be applied to numerous heterostructures and composites known in the art and/or described herein. It also may apply to structures not yet known or discovered.
Example P28: In another exemplary procedure, BN perimorphic materials comprising synthetic anthracitic networks may be synthesized directly on a template material via chemical vapor deposition. This synthesis demonstrates that, in a way that is analogous to the formation of synthetic anthracitic networks from graphenic carbon, BN anthracitic networks may be synthesized from graphenic BN via surface defect-catalyzed BN lattice nucleation and free radical-driven BN lattice growth. Hence, BN perimorphic frameworks can be synthesized via template-directed surface replication procedures in a way that is analogous to the template-directed synthesis of carbonaceous perimorphic frameworks. This being the case, the General Method (and the Preferred Method) can be utilized to synthesize these perimorphic materials and other perimorphic materials that are formed according to analogous nucleation and growth mechanics.
Additionally, like free radical-driven carbon growth processes, the formation of BN spx networks via free radical-driven BN growth processes should be optimized by modulating the amount of hydrogen gas in the gas medium during growth. This will prevent hydrogen from being too rapidly released from the growing BN domains and enable the tectonic interfaces to rearrange into configurations that maximize the edge-to-edge sp2 and sp3 grafting of BN domains. Maturation of these BN spx precursor networks via annealing can then be utilized to transform them into BN helicoidal networks that are substantially sp2-hybridized-again according to mechanics analogous to the maturation of carbonaceous spx precursor networks and the formation of carbonaceous helicoidal networks.
For the purpose of demonstration, N2T1-type porous MgO template structures may first be synthesized via thermal decomposition of N-type nesquehonite template precursor structures in a way consistent with the General Method (and the Preferred Method). This synthesis is described in the '49195 Application and Reference A.
Next, in an exemplary Replication Stage procedure, the N2T1-type template material may be utilized to direct the chemical vapor deposition of a disordered BN. To do this, a 176 mg quantity of N2T1-type template structures may be placed in a ceramic boat, which may be placed in a tube furnace, according to Scheme B, as detailed in Section III. A ceramic boat containing 1.0 g of ammonia borane complex (H3NBH3) may be placed in the quartz tube just outside the furnace's heated zone, such that when the furnace reaches a temperature of 900° C., the H3NBH3 reaches a temperature between 130° C. and 170° C. The furnace may initially be purged with Ar flowing at 2000 sccm for 30 minutes at room temperature. This may be followed by heating to a temperature of 900° C. at a heating rate of 20° C./min under Ar flowing at 2000 sccm. Upon reaching 900° C., the furnace may be maintained at 900° C. for 60 minutes, then allowed to cool to room temperature under continued Ar flow.
Next, endomorphic extraction may be performed, as it would be in the Separation Stage of the General Method or Preferred Method. This may be done in an aqueous H2CO3 extractant solution. After dissolution of the endomorphic MgO, the BN perimorphic frameworks may be filtered, rinsed and dried. In a full implementation of the General Method, the aqueous H2CO3 solution may be generated using the retained process water from the Precursor Stage precipitation, and the aqueous Mg(HCO3)2 solution may be utilized as the solution stock for precipitating an MgCO3 xH2O template precursor material like N2-type nesquehonite. In this way, both the template material and process liquid are conserved for cyclical use. CO2 process gas may also be beneficially conserved and utilized to regenerate the extractant solution. The type of BN perimorphic frameworks resulting from this process is identified herein as P28.
The native or near-native morphological state of the P25-type BN perimorphic frameworks, as shown in
Similar to perimorphic walls constructed from deposition of carbonaceous graphene, perimorphic walls constructed from the deposition of other two-dimensional molecular structures like sp2-hybridized BN may be thinned or thickened via a shorter or longer CVD procedure, respectively. Thinning them results in a more flexible framework, as shown in
Example P29: In another exemplary procedure, boron carbonitride (BCxN) perimorphic materials comprising synthetic anthracitic networks may be synthesized directly on a template material via chemical vapor deposition. This synthesis demonstrates that, in a way that is analogous to the formation of synthetic anthracitic networks from graphenic carbon and graphenic BN, BCxN anthracitic networks may be synthesized from graphenic BCxN via surface defect-catalyzed BCxN lattice nucleation and free radical-driven BCxN lattice growth. Hence, BCxN perimorphic frameworks can be synthesized via template-directed surface replication procedures in a way that is analogous to the template-directed synthesis of carbonaceous perimorphic frameworks. This being the case, the General Method (and the Preferred Method) can be utilized to synthesize these perimorphic materials and other perimorphic materials that are formed according to analogous nucleation and growth mechanics.
Additionally, like free radical-driven carbon growth processes, the formation of BCxN spx networks via free radical-driven BCxN growth processes should be optimized by modulating the amount of hydrogen gas in the gas medium during growth. This will prevent hydrogen from being too rapidly released from the growing BCxN domains and enable the tectonic interfaces to rearrange into configurations that maximize the edge-to-edge sp2 and sp3 grafting of BCxN domains. Maturation of these BN spx precursor networks via annealing can then be utilized to transform them into BN helicoidal networks that are substantially sp2-hybridized-again according to mechanics analogous to the maturation of carbonaceous spx precursor networks and the formation of carbonaceous helicoidal networks.
For the purpose of demonstration, an N2-type template precursor material may first be synthesized in a way consistent with the General Method (and the Preferred Method). This synthesis is described in the '49195 Application and Reference A.
Next, the template precursor material may be thermally treated. This may be performed according to Scheme B in a tube furnace, as detailed in Section III. For this treatment, approximately 7.88 g of the N2-type powder may be placed in the tube furnace. A ceramic boat containing 1.30 g of ammonia borane complex (H3NBH3) may be placed in the quartz tube just outside the furnace's heated zone, such that when the furnace reaches a temperature of 700° C., the H3NBH3 reaches a temperature between 130° C. and 170° C. After sealing the tube, an Ar gas flow of 2000 sccm may be initiated. Under flowing Ar, the furnace may be heated from room temperature to a temperature setting of 700° C. at a heating rate of 20° C./min. During this thermal treatment, CO2 gas may be released. In a full implementation of the General Method, the CO2 process gas released during the decomposition of the template precursor material may be recaptured and conserved using conventional techniques. The type of porous MgO template material resulting from this thermal process is identified herein as N2T7. The elongated superstructures of these porous templates are derived from N2-type nesquehonite template precursor particles, which are also elongated.
Approximately 2 minutes after the furnace reaches the temperature setting of 700° C., white fumes may be evolved and observed in the tube and in the exhaust. This indicates the initial vaporization of the H3NBH3. Approximately 5 minutes after the furnace reaches the temperature setting of 700° C., a 64 sccm flow of C3H6 may be initiated, and this condition may be maintained for 5 min, such that H3NBH3 and C3H6 are flowing simultaneously. The C3H6 gas flow may then be terminated and the H3NBH3 flow may continue for an additional 5 minutes. The furnace may then be allowed to cool to room temperature under continued Ar flow.
The bed of powder resulting from this procedure comprises a light phase and a dark phase on the surface. The bed of powder in the boat, as retrieved from the furnace, is shown in
Next, endomorphic extraction of the N2T7P29 phase may be performed, as it would be in the Separation Stage of the General Method (or Preferred Method). This may be done in an aqueous H2CO3 extractant solution. After dissolution of the endomorphic MgO, the BN perimorphic frameworks may be filtered, rinsed and dried. In a full implementation of the General Method, the aqueous H2CO3 solution may be generated using the retained process water from the Precursor Stage precipitation, and the aqueous Mg(HCO3)2 solution may be utilized as the stock solution for precipitating an MgCO3·xH2O template precursor material like N2-type nesquehonite. In this way, both the template material and process liquid are conserved for cyclical use. CO2 process gas may also be beneficially conserved and utilized to regenerate the extractant solution. The type of BCxN perimorphic frameworks resulting from this process is identified herein as P29.
The Raman spectrum of the light phase associated with BCxN perimorphic material can be contrasted with the dark phase, as shown in
The Detailed Description begins with initial section of “Terms and Concepts” that provides language and concepts for describing and understanding the invention. Subsequent sections are organized according to the following four stages of the method: the “Precursor Stage,” the “Template Stage,” the “Replication Stage,” and the “Separation Stage.” A number of exemplary procedures and materials pertaining to each of the 4 stages are demonstrated. Many potential variants of each stage might be readily conceived by those knowledgeable in the art and combined to form numerous variants without deviating from the method.
The Detailed Description is organized according to the following sections:
A “template,” as defined herein, is a potentially sacrificial structure that imparts a desired morphology to another material formed in or on it. Of relevance for surface replication techniques are the template's surface (i.e. the “templating surface”), which is positively replicated, and its bulk phase (i.e. the “templating bulk”), which is negatively replicated. The template may also perform other roles, such as catalyzing the formation of the perimorphic material. A “templated” structure is one that replicates some feature of the template.
A “perimorph” or “perimorphic” material is a material formed in or on a solid-state or “hard” template material.
“Surface replication,” as defined herein, comprises a templating technique in which a template's surface is used to direct the formation of a thin, perimorphic wall of adsorbed material, the wall substantially encapsulating and replicating the templating surface upon which it is formed. Subsequently, upon being displaced, the templating bulk is replicated, in negative, by an endocellular space within the perimorphic wall. Surface replication creates a perimorphic framework with a templated pore-and-wall architecture.
A “perimorphic framework” (or “framework”), as defined herein, is the nanostructured perimorph formed during surface replication. A perimorphic framework comprises a nanostructured “perimorphic wall” (or “wall”) that may range from less than 1 nm to 100 nm in thickness but is preferably between 0.6 nm and 5 nm. Insomuch as it substantially encapsulates and replicates the templating surface, the perimorphic wall can be described as “conformal.” Perimorphic frameworks may be made with diverse architectures, ranging from simple, hollow architectures formed on nonporous templates to labyrinthine architectures formed on porous templates. They may also comprise different chemical compositions. A typical framework may be constructed from carbon and may be referred to as a “carbon perimorphic framework.”
An “endomorph,” as defined herein, comprises a template as it exists within a substantially encapsulating perimorphic phase. Therefore, after the perimorphic phase has been formed around it, the template may be described as an endomorph, or as “endomorphic.”
A “perimorphic composite,” or “PC” material, as defined herein, is a composite structure comprising an endomorph and a perimorph. A PC material may be denoted x@y, where x is the perimorphic element or compound and y is the endomorphic element or compound. For example, a PC structure comprising a carbon perimorph on an MgO endomorph might be denoted C@MgO.
The term “positive” is used herein to describe a space that is occupied by a solid mass. The space occupied by the endomorph (i.e. the “endomorphic space”) in a perimorphic composite is an example of positive space. A nonporous template comprises only positive space. Exempting the space occupied by its thin wall, a perimorphic framework comprises no positive space.
The term “negative” is used herein to describe a space that is unoccupied by a solid or liquid mass. A negative space may be empty, gas-filled, or liquid-filled. The pores inside an unimpregnated, porous template comprise negative space. A porous template comprises both positive and negative space. Exempting the space occupied by its thin wall, a perimorphic framework comprises only negative space.
The term “cellular” is used herein to describe the pore-and-wall morphology associated with perimorphic frameworks. A “cell” or “cellular subunit” comprises a specified endocellular pore and region of the perimorphic wall around the pore.
The term “endocellular” is used herein to describe a negative space in a perimorphic framework that is formed by the displacement of the endomorph from the perimorphic composite. Like the endomorph whence it derives, the endocellular space is substantially encapsulated by the perimorphic wall.
The term “exocellular” is used herein to describe a negative space in a perimorphic framework that is inherited from the pore space of the perimorphic composite, which is in turn inherited from the pore space of a porous template. We note that an exocellular space, despite the “exo-” prefix, maybe located substantially inside a perimorphic framework.
A perimorphic framework's endocellular and exocellular spaces are substantially separated by the perimorphic wall. However, the ability to displace the endomorph from the template composite implies that the wall is somewhere open or an incomplete barrier, since a perfectly encapsulated endomorph could not be displaced. Therefore, while a perimorph is herein described as substantially encapsulating a templating surface, the encapsulation may nevertheless be incomplete or subject to breach.
The term “native” is used herein to describe the morphological state of a perimorphic structure in the perimorphic composite. A “native” feature comprises a feature that is substantially in its native state, and we may refer to a structure as “natively” possessing some feature (e.g. a perimorphic wall that is natively 1 nm thick). After displacement of the endomorph from the perimorphic composite, the perimorph may either substantially retain its native characteristics, or it may be altered.
The term “non-native” is used herein to describe a morphological state of a perimorphic structure that is substantially altered from its native morphological state (i.e. its original state in the perimorphic composite). This alteration may occur at the substructural or superstructural levels. For example, during evaporative drying of an internal liquid, a perimorphic wall may be pulled inward by the liquid, collapsing a portion of the endocellular space. A framework's deformation into a non-native, collapsed morphology may be reversible—i.e. the framework may be able to substantially recover its native morphology.
The term “labyrinth” or “labyrinthine” is used herein to describe a network of interconnected pores in a template or a perimorphic framework. A labyrinth may be endocellular or exocellular. A perimorphic framework formed on a porous template may natively comprise endocellular and exocellular labyrinths; therefore, a framework formed on porous templates may be described as a “labyrinthine framework.” The endocellular and exocellular labyrinths of a labyrinthine framework, while not overlapping, may be interwoven. Labyrinthine frameworks comprise a preferred class of perimorphic frameworks.
A “template precursor,” or “precursor,” as defined herein, is a material from which a template is derived via some treatment that may comprise decomposition, grain growth, and sintering. A template may retain a pseudomorphic resemblance to the template precursor; therefore, engineering the precursor may offer a way to engineer the template. The precursor is formed within a process liquid and is derived from a stock solution.
The term “superstructure” is herein defined as the overall size and geometry of a porous template or perimorphic framework. A perimorphic framework's superstructure may be inherited from the morphology of the template precursor. The superstructure of a perimorphic framework is important because the overall size and geometry of a framework will influence its properties, including how it interacts with other particles. Some superstructures may facilitate the drying of a wet paste of perimorphic frameworks into a fine powder, whereas other superstructures may cause a wet paste to dry into macroscopic granules, which may require subsequent grinding. Superstructures may comprise the following shapes:
The term “substructure” is herein defined as the localized morphology—i.e. the internal architecture—of a porous template or perimorphic framework. Certain porous templates or perimorphic frameworks have a substructure comprising repeating, conjoined substructural units, or “subunits.” Different substructures may be characterized by subunits of different shapes, sizes, and spacings from one another.
The term “noncellular” is used herein to describe a negative space inside that is not considered herein to be templated, nor to be part of a perimorphic framework, but that is nevertheless substantially surrounded by and located within a framework. Noncellular space is not templated space because it does not correspond strictly to a template's positive space, negative space, or surface, and it is only present when surface replication is not able to occur on some portion—typically an inaccessible interior region—of a templating surface.
Noncellular space may be desirable for density reduction in certain applications and may be engineered using a combination of rational template engineering and diffusion-limited synthesis techniques. In particular, large template precursors may be used to create large templates, which with minimal sintering may combine long diffusion pathways with small pores. Rational design of the surface replication parameters may also help. For instance, during CVD, low concentrations of the carbonaceous vapor may be more easily scavenged by reactive sites and prevented from penetrating throughout the porous substructure of the porous template.
Another way density reduction can be achieved is via the use of porous template precursor materials. This results in an exocellular internal porosity that is preferable to noncellular space because it is more engineerable and does not require diffusion constraints. Porous template precursors may be made with the use of blowants (e.g. hollow microspheres produced during spray-drying) or via the use of sacrificial materials when making the template precursors (e.g. synthesizing the template precursor around sacrificial micelles or polymers).
The concept of “compactness” herein relates to the area of perimorphic wall contained within a given volume of a perimorphic framework—i.e. a volumetric surface area. A framework with a more compact substructure will possess a finer, denser arrangement of perimorphic wall within a given volume, whereas a framework with a less compact substructure will possess a coarser, more spatially diffuse arrangement of perimorphic wall within a given volume. Porous templates, and the labyrinthine frameworks formed on them, may be engineered to have different levels of compactness. Compactness comprises a measure of a framework's mesoscale crosslinking—i.e. crosslinking not at the molecular scale, but at a higher scale, where crosslinking is derived from the topology of the templating surface.
A perimorphic framework's compactness and pore phases may be modulated by engineering the template's positive and negative spaces. For example, a porous MgO structure produced by decomposing a magnesium carbonate precursor has a positive space comprising a network of conjoined, MgO crystallites. Its negative space comprises a porous network running between the MgO crystallites and throughout the structure. It is well-known that the crystallites may grow at elevated temperatures, coarsening the grain structure. The same process may also lead to growth and coarsening of the pores. This coarsening of the positive and negative spaces will reduce the porous MgO template's surface area, and therefore reduce the compactness of a perimorphic framework formed on the template. At the same time that the template is coarsening, it will be densifying, and its densification will reduce the amount of exocellular space in a perimorphic framework formed on the template.
The coarseness of a template may be important for many reasons. For example, enlarging a template's pores and reducing its surface area may permit faster, deeper diffusion of a reactive vapor throughout the template's pores during CVD. The perimorphic walls synthesized in such processes may be more uniform in thickness if adequate diffusion kinetics can be achieved.
A perimorphic framework's compactness and pore phases may also be modulated by selecting different template precursors. Different precursors will have different fractions of labile mass. A template's negative space will depend on how much of the template precursor's starting mass is lost during its decomposition. Calcining template precursors that contain large fractions of labile species—for instance, highly hydrated salts—may result in porous templates with high specific porosity that are more open to diffusive flows during CVD. Such templates may also be desired if more exocellular space is desired in the perimorphic framework.
The term “recycled” is used herein to describe process materials being utilized in a given step of the production process that have previously been utilized for that step. Since practical losses of process materials (e.g. process liquid losses from evaporation or filtration) during production of a perimorphic product may occur, virgin process materials may be used to replenish these losses, and a “recycled” process material may partially comprise virgin material.
“Process materials,” as defined herein, comprise potentially recyclable non-perimorphic materials utilized to generate perimorphic materials. Process materials may comprise process liquids, process gases, extractants, template precursor materials, and template materials.
A “stock solution,” as defined herein, comprises solvated cations and anions and a process liquid, the solvated ions being hosted by the process liquid (which may be referred to in this context as the “host”). A stock solution is formed in the Separation Stage. A precursor is derived from a stock solution through one or more precipitation, dissolution, or decomposition reactions.
A “process liquid,” as defined herein, is a feedstock of either liquid water (“process water”) or solvent (“process solvent”) utilized in the Precursor Stage and the Separation Stage. The process liquid may play a number of different roles in these stages. In the Precursor Stage, the formation of a template precursor is hosted by the process liquid, and the precursor may incorporate the process liquid into its crystal—for instance, a hydrous salt may be formed in process water and incorporate some of the process water into its crystal structure. In the Separation Stage, an extractant is hosted by the process liquid, and the solvated ions produced by reactions between the template, process liquid, and extractant are hosted by the process liquid. The process liquid may be involved in the production of the extractant and may itself react with the template during the Separation Stage.
A “residual liquid,” as defined herein, is a portion of process liquid, which may or may not host solvated ions, that remains unseparated from a solid (e.g. a precursor or a perimorphic product) upon separation of the solid from the main portion of the process liquid. Residual liquid may be contained within a perimorphic product or wetted to its surface. Residual liquid may comprise a very small fraction of the overall process liquid. A solid's retention of residual liquid may require further separation if a dry solid is desired.
An “extractant,” as defined herein, comprises an acid hosted by a process liquid, the two phases together comprising an “extractant solution.” The extractant may be present in the extractant solution in very dilute concentrations. In some cases, the extractant may be produced from (and within) the process liquid. For example, a carbonic acid (H2CO3) extractant may be produced from (and within) a process water, according to the reaction H2O(l)+CO2(aq)→H2CO3(aq).
“Endomorphic extraction,” as defined herein, comprises the selective removal of a portion of an endomorph from a perimorphic composite. Endomorphic extraction comprises a reaction between an endomorph and an extractant solution that produces solvated ions that are exfiltrated from the surrounding perimorph, resulting in concurrent displacement of the endomorph, consumption of the extractant from the extractant solution, and generation of a stock solution. Generally, the removal of substantially all of an endomorph's mass is desired. Occasionally the partial removal of an endomorph's mass nay be desired, or only partial removal of an endomorph's mass may be achievable.
“Perimorphic separation,” as defined herein, comprises the separation of a perimorphic product after endomorphic extraction from non-perimorphic, conserved process materials. Conserved, non-perimorphic phases may comprise process liquid, stock solution, and precipitates of the stock solution. Perimorphic separation may comprise many different industrial separation techniques, (e.g. filtration, centrifugation, froth flotation, solvent-based separations, etc.)
A “solventless precipitation,” as defined herein, comprises the precipitation of a template precursor in the Precursor Stage, wherein the precipitation is substantially driven by a solution destabilization mechanism that does not require the introduction of a miscible antisolvent into the process liquid. As a first example of a solventless precipitation technique, a stock solution may be spray-dried. As a second example of a solventless precipitation technique, a metastable metal bicarbonate stock solution may be depressurized to reduce CO2 solubility, causing CO2 gas to be released and a metal carbonate to be precipitated. We note that the term “solventless precipitation” does not imply the absolute absence of a miscible liquid or solvent during precipitation, but rather indicates that precipitation is not principally driven by mixing a miscible liquid into the stock solution. One scenario that could be envisioned is a miscible liquid mixed with the process liquid that remains at substantially the same concentration throughout the Liquid Cycle.
“Shuttling,” as defined herein, comprises an endomorphic extraction technique that may be used during the Separation Stage, wherein, concurrently: (i) an extractant is generated via reaction of a process gas with a process liquid; (ii) an endomorph is reacted with the extractant solution; (iii) the extractant is consumed; (iv) the solvated ions in the stock solution are exfiltrated from a perimorph; and (v) a precipitate is formed from the stock solution outside of the perimorph. For example, shuttling may comprise, concurrently: (i) forming H2CO3 extractant via dissolving CO2 into process water; (ii) reacting an MgO endomorph with the H2CO3 extractant solution; (iii consuming HCO3; (iv) forming Mg2+ and (HCO3)− ions that are exfiltrated from a perimorph; and (v) precipitating magnesium carbonate in the surrounding process water.
“MgCO3·xH2O” is herein used to describe a magnesium carbonate. It may comprise any hydrous or anhydrous magnesium carbonate, as well as basic magnesium carbonates such as hydromagnesite.
A “Template Cycle,” as defined herein, comprises a cyclical loop in which a template is constituted, utilized, and reconstituted.
A “Liquid Cycle,” as defined herein, comprises a cyclical loop in which a process liquid is utilized for liquid-phase extraction of the endomorph and liquid-phase formation of the precursor.
A “Gas Cycle,” as defined herein, comprises a cyclical loop in which a process gas is dissolved into a process liquid to create an extractant solution, then subsequently released and recaptured. The release may be associated with the formation of either a template precursor or a template.
The “yield” of a perimorphic material, or of a procedure used to make the perimorphic material, is defined herein as the perimorphic mass divided by the sum of the endomorphic and perimorphic masses. The yield can be used to understand how much template material is required to create a given amount of perimorphic material.
The substructure represented in the center of the diagram is somewhat more compact than the left-hand substructure, because its volume contains more perimorphic area. The substructure represented on the right of the diagram is the most compact-its volume, though similar to the volume of the other two substructures, contains the most perimorphic area. This diagram demonstrates that a perimorphic framework's compactness is imparted by the volume-specific surface area of the porous template—i.e. the total amount of internal and external surface area per unit of template volume, where the template volume includes the template's positive and negative spaces.
The “General Method” is the most basic form of the method. It comprises a method for synthesizing a perimorphic product wherein substantial portions of the template material and the process liquid are conserved and may be reused. As such, the General Method may be performed cyclically. All variants of the method disclosed in the present disclosure comprise some variant of the General Method.
The General Method comprises a series of steps that is herein presented, for ease of description, in 4 stages (i.e. the Precursor Stage, Template Stage, Replication Stage, and Separation Stage). Each stage is defined according to one or more steps, as described below:
In practice, each step within these stages may itself comprise multiple, subsidiary steps. Additionally, each of the steps may occur concurrently with steps from another stage, such that in practice different stages may overlap in chronology. This can especially be expected in variants employing one-pot techniques. As a hypothetical example of this, a stock solution might be continuously sprayed alongside an adsorbate material into a furnace. In this hypothetical furnace, precursor particles might be precipited from the stock solution, template particles might be formed by heating of the precursor particles, and perimorphic material might be adsorbed to the template particles continuously and concurrently. This would correspond to steps assigned herein to the Precursor Stage, Template Stage, and Replication Stage, respectively.
Similarly, it is anticipated that in practice, many variants of the General Method may incorporate the steps described in the 4 Stages in different sequences. Also, in some variants, a step assigned by definition to one of the four stages herein might instead occur in a different stage. Such variants are anticipated herein and do not deviate from the inventive method, which is only presented herein as a discrete sequence of 4 stages for the sake of describing the overall cycle.
Ancillary processing steps (e.g. rinsing, drying, blending, condensing, spraying, agitating, etc.) may also be incorporated into the method at each stage. As a hypothetical example of this, a Replication Stage might involve coating a template material with a perimorphic material via a liquid-phase adsorption procedure, then filtering, rinsing and drying the resulting PC material. The incorporation of these processing steps in many variants will be obvious to those skilled in the art and, as such, they are not enumerated herein.
The inputs and outputs of the General Method are illustrated in
Variants of the General Method
The following discussion enumerates a number of ways in which the General Method may be variously implemented. The omission of variants from this discussion should not be interpreted as limiting, since an exhaustive list of ways in which the General Method may be implemented is not practical.
The General Method is intended to offer a means for cyclical production of perimorphic products while conserving process materials. In each cycle of the General Method, some portion of the process materials utilized are conserved and reused. In some variants, substantially all of the process materials utilized may be conserved and reused. In other variants, a portion of the process materials may be lost. One hypothetical example of this would be evaporative losses of process liquids from open tanks or wet filters.
In some variants of the General Method, process steps may correspond to batch processes. In other variants, process steps may correspond to continuous processes.
In some variants of the General Method, the solventless precipitation may comprise at least one of the following techniques: heating or cooling the stock solution to change the solubility of a solute in the stock solution; volatilizing a dissolved gas within the stock solution; depressurizing the stock solution; atomization of the stock solution; spray-drying the stock solution or spray pyrolysis.
In some variants of the General Method, a precursor structure may comprise at least one of the following: an elongated, thin, equiaxed, or hierarchical-equiaxed superstructure; an elongated superstructure with a length-to-diameter ratio greater than 200:1; an elongated superstructure with a length-to-diameter ratio between 50:1 and 200:1; a spheroidal or spherical superstructure; a hollow superstructure; a fragmentary superstructure comprising fragments of some other parent superstructure; a curved, fragmentary superstructure comprising fragments of a hollow superstructure.
In some variants of the General Method, a precursor structure may be precipitated around one or more other sacrificial structures, which may be present as inclusions in the precursor structure after its precipitation. In some variants, these inclusions in the precursor structure may be subsequently removed, resulting in voids.
In some variants of the General Method, a precursor structure may measure less than 1 μm along its major axis. In some variants, the precursor may measure between 1 μm and 100 μm along its major axis. In some variants, the precursor may measure between 100 μm and 1,000 μm along its major axis.
In some variants of the General Method, the precursor material may comprise at least one of the following: a hydrate; a metal bicarbonate or carbonate; a Group I or Group II metal bicarbonate or carbonate; a mixture of salts. In some variants, the precursor may comprise MgCO3·xH2O in the form of at least one of: hexahydrate, lansfordite, nesquehonite, hydromagnesite, dypingite, magnesite, nanocrystalline or non-crystalline MgCO3·xH2O.
In some variants of the General Method, the stock solution may comprise at least one of the following: metal cations and oxyanions; an aqueous metal bicarbonate solution; a Group I or Group II metal bicarbonate; an organic salt; Mg(HCO3)2. In some variants, the stock solution may comprise at least one of a dissolved gas, acid, and base. In some variants, the stock solution may be metastable.
In some variants of the General Method, the process liquid conserved in the Precursor Stage may comprise a distillate. In some variants, the distillate may be formed by condensing the process liquid vapor formed during spray-drying or spray-pyrolysis. In some variants, a process liquid conserved in the Precursor Stage may host solvated ions, the process liquid and ions together comprising a mother liquor.
In some variants of the General Method, the treatment performed on a precursor material in the Template Stage may comprise at least one of the following: decomposing the precursor; partially or locally decomposing the precursor; decomposing the precursor surface; thermal decomposition; and oxidizing an organic phase present within a precursor structure. In some variants, the treatment may comprise flash-drying, spray-drying, spray pyrolysis, vacuum drying, rapid heating, slow heating, sublimation. In some variants, a vapor released during the treatment may be conserved. In some variants, the vapor released may comprise at least one of CO2 and H2O. In some variants, treatment may comprise at least one of: coarsening the grain structure of the precursor or a decomposition product of the precursor; exposing to a reactive vapor; exposing to water vapor; sintering; sintering with the assistance of dopants.
In some variants of the General Method, a template material may comprise at least one of the following: a metal carbonate, a metal oxide, a Group I or II metal oxide, a transition metal, and MgO. In some variants, a template structure may comprise at least one of the following: macropores, mesopores, hierarchical porosity, subunits larger than 100 nm, subunits between 20 nm and 100 nm, and subunits between 1 nm and 20 nm.
In some variants of the General Method, a template structure may comprise at least one of the following: an elongated, thin, equiaxed, or hierarchical-equiaxed superstructure; an elongated superstructure with a length-to-diameter ratio greater than 200:1; an elongated superstructure with a length-to-diameter ratio between 50:1 and 200:1; a spheroidal or spherical superstructure; a hollow superstructure; a fragmentary superstructure comprising fragments of some other parent superstructure; and a curved, fragmentary superstructure comprising fragments of a hollow superstructure.
In some variants of the General Method, adsorbing the perimorphic material to the templating surface may comprise at least one of the following: a coating technique, physical vapor deposition, and chemical vapor deposition. In some variants, the coating technique may comprise coating a liquid or solid organic coating onto the templating surface, then forming a derivative carbon coating from the parent coating. In some variants, deposition may comprise pyrolytic decomposition of a vapor-phase organic compound at a temperature between 350° C. and 950° C. In some variants, a perimorphic carbon may be annealed after being adsorbed to the templating surface.
In some variants of the General Method, endomorphic extraction may utilize an extractant solution comprising a weak acid as an extractant. In some variants, an extractant solution may be formed by dissolving a process gas in process water. In some variants, the extractant solution may be an aqueous solution of H2CO3 formed by dissolving liquid or gaseous CO2 in process water. In some variants, endomorphic extraction may comprise a shuttling technique. In some variants, endomorphic extraction may be performed under conditions of elevated pressure or temperature.
In some variants of the General Method, the perimorphic separation may comprise at least one of decantation, hydrocyclones, settling, sedimentation, flotation, froth flotation, centrifugal separation, filtration, and liquid-liquid extraction. In some variants, the perimorphic separation may separate the perimorphic product from substantially all of the process liquid. In some variants, the perimorphic product may retain a residual portion of the process liquid. In some variants, the perimorphic product may be naturally buoyant due to its retention of internal gas. In some variants, the perimorphic product's internal gas may be expanded by reducing pressure of the surrounding process liquid, increasing the buoyancy of the perimorphic product and causing flotation. In some variants, a portion of the perimorphic product's internal gas may be exfiltrated by reducing pressure of the surrounding process liquid, followed by re-pressurizing the surround process liquid, such that hydrostatic pressure causes the process liquid to infiltrate the perimorphic product.
In some variants of the General Method, the perimorphic framework may comprise at least one of a carbonaceous material, a pyrolytic carbon, an anthracitic network of carbon, an spx network of carbon, and a helicoidal network of carbon.
In some variants of the General Method, under 532 nm excitation, the carbonaceous perimorphic framework may comprise at least one of a Raman spectral ID/IG ratio of between 4.0 and 1.5; a Raman spectral ID/IG ratio between 1.5 and 1.0; a Raman spectral ID/IG ratio between 1.0 and 0.1; a Raman spectral ITr/IG ratio between 0.0 and 0.1; a Raman spectral ITr/IG ratio between 0.1 and 0.5; a Raman spectral ITr/IG ratio between 0.5 and 1.0; a Raman spectral I2D/IG ratio between 0 and 0.15; a Raman spectral I2D/IG ratio between 0.15 and 0.3; and a Raman spectral I2D/IG ratio between 0.30 and 2.0.
In some variants of the General Method, under 532 nm excitation, the carbonaceous perimorphic framework may comprise at least one of an unfitted Raman spectral D peak positioned between 1345 and 1375 cm−1; an unfitted Raman spectral D peak positioned between 1332 and 1345 cm−1; an unfitted Raman spectral D peak positioned between 1300 and 1332 cm−1; an unfitted Raman spectral G peak positioned between 1520 cm−1 and 1585 cm−1; an unfitted Raman spectral G peak positioned between 1585 cm−1 and 1600 cm−1; and an unfitted Raman spectral G peak positioned between 1600 cm−1 and 1615 cm−1.
In some variants of the General Method, the perimorphic product may comprise a perimorphic framework. In some variants, the perimorphic framework may comprise at least one of a native morphology, a non-native morphology, internal gas, a hydrophobic surface, a hydrophilic surface, mesopores, one or more macropores, hierarchical porosity.
In some variants of the General Method, the perimorphic framework may measure less than 1 μm along its major axis. In some variants, the perimorphic framework may measure between 1 μm and 100 m along its major axis. In some variants, the perimorphic framework may measure between 100 μm and 1,000 μm along its major axis. In some variants, the perimorphic framework may comprise an elongated, thin, equiaxed, or hierarchical-equiaxed superstructure. In some variants, an elongated perimorphic framework may comprise a length-to-diameter ratio between 50:1 and 200:1. In some variants, the perimorphic framework's equiaxed superstructure may be spheroidal or spherical. In some variants, the perimorphic framework's equiaxed superstructure may be hollow. In some variants, the perimorphic framework may comprise fragments of a hollow shell. In some variants, the perimorphic framework may comprise a noncellular space.
In some variants of the General method, the perimorphic framework may comprise a BET surface area of 1,500 to 3,000 m2/g. In some variants, the perimorphic framework may comprise a BET surface area of 10 to 1,500 m2/g.
In some variants of the General Method, the perimorphic product may be subjected to further treatment after perimorphic separation. In some variants, the further treatment after perimorphic separation may comprise at least one of flash-drying, spray-drying, spray-pyrolysis, decomposition, chemical reaction, annealing, and chemical functionalization.
In some variants of the General Method, the Liquid Cycle may also incorporate the recapture and conservation of process liquid released (possibly in vapor phase) during the Template Stage, although this is not reflected as an output in
In some variants of the General Method, a Gas Cycle may be incorporated into the method. The inputs and outputs of the General Method with a Gas Cycle are illustrated in
The Preferred Method, described below, comprises variants of the General Method in which a MgCO3·xH2O template precursor material is derived from an aqueous Mg(HCO3)2 stock solution and a portion of the CO2 process gas is conserved via a Gas Cycle. The inputs and outputs of the Preferred Method are shown in
Certain variants of the Preferred Method may employ pressure modulations in order to form concentrated stock solutions and improve precipitation processes. Concentrated stock solutions may be associated with many benefits, including superior precipitation kinetics, reduced process water volumes, smaller vessels, and improved energy efficiency. Two exemplary ways that this can be done are illustrated in
In the first frame of
Another way that a concentrated stock solution may be obtained is by performing the endomorphic extraction in a pressurized reactor. A schematic showing this is illustrated in
In the course of describing procedures to generate the exemplary materials described in the subsequent sections, certain furnace schemes have been detailed. These schemes may be used for the exemplary Template Stage procedures detailed in Section V and for the exemplary Replication Stage procedures detailed in Section VI.
Scheme A: In Scheme A, a Thermcraft tube furnace modified to be a rotary furnace may be employed with a quartz tube (
For exemplary procedures performed using Scheme A, a material sample may be placed inside the belly, such that it agitated within the reactor. Loose fitting ceramic blocks located outside the belly section on each side of the furnace's heating zone allowing for gas flow and powder containment. Packed glass wool may be used to affix the position of the ceramic blocks while acting as a gas permeable layer. The ends of the tube may be fitted with two stainless-steel flanges to allow for gas to flow for the system.
Scheme B: An MTI rotary tube furnace with a quartz tube (
Scheme C: A Lindberg Blue-M tube furnace with a quartz tube may be used. The quartz tube may be 150 mm in OD. The furnace has a clam shell design with a cylindrical heated chamber having dimensions of 190 mm diameter and 890 mm heated length. The furnace has a wattage of 11,200 W with a maximum operating temperature of 1200° C. The tube may be substantially level. For exemplary procedures performed using Scheme C, a sample may be placed within a ceramic boat. This may then be placed inside the quartz tube within the heating zone prior to the initialization of heating. Loose fitting ceramic blocks located outside the furnace's heating zone allow for gas flow. The ends of the tube may be fitted with two aluminum flanges to allow for gas flow through the system.
Scheme D: A Vulcan 3-550 Muffle furnace may be used. The furnace has a rectangular heated chamber having dimensions of 190 mm×240 mm×228 mm. The furnace has a wattage of 1440 W with a maximum operating temperature of 1100° C. For exemplary procedures performed using Scheme D, a material sample may be placed within a ceramic boat. This may then be placed inside the muffle furnace prior to the initialization of heating.
Scheme E: A TA Instruments Q600 TGA/DSC may be used. For exemplary procedures performed using Scheme E, a 90 μL alumina pan may be used to hold a material sample. Gas flow may be 100 sccm of a specified gas unless otherwise noted. The heating rate may be mentioned in the exemplary procedures where Scheme E is used.
A number of analytical techniques were utilized to characterize the procedures and materials presented herein. These are detailed below.
Solution concentrations were measured using electrolytic conductivity (“conductivity”). The conductivity is a measured response of a solution's electrical conductance. The electrical response of a solution may be correlated to the concentration of ions dissolved in the solution, and as ions in solution are precipitated, the conductivity value decreases. An analog to this measurement is total dissolved solids (“TDS”) which relates the conductivity measurement to a referenced ion concentration (typically potassium chloride), dependent on the salt compound dissolved.
Thermogravimetric analysis (TGA) was used to analyze the thermal stability and composition of materials. All TGA characterization was performed on a TA Instruments Q600 TGA/DSC. A 90 μL alumina pan was used to hold the sample during TGA analysis. All analytical TGA procedures were performed at 20° C. per min unless otherwise mentioned. Either air or Ar (Ar) was used as the carrier gas during analytical TGA procedures unless otherwise mentioned.
Raman spectroscopy was performed using a ThermoFisher DXR Raman microscope equipped with a 532 nm excitation laser. For each sample analyzed, 16 point spectra were generated using measurements taken over a 4×4 point rectangular grid with point-to-point intervals of 5 μm. The 16 point spectra were then averaged to create an average spectrum. The Raman peak intensity ratios and Raman peak positions reported for each sample all derive from the sample's average spectrum. No profile fitting software was utilized, so the reported peak intensity ratios and peak positions relate to the unfitted peaks pertaining to the overall Raman profile.
Gas adsorption measurements were made using a Micromeritics Tristar II Plus. Nitrogen adsorption was measured at a temperature of 77 K across a range of pressure (p) values, where
Increments of pressure ranged from
Micromeritics MicroActive software was used to calculate the BET specific surface area derived from the BET monolayer capacity assuming the cross-sectional area σm(N2, 77 K)=0.162 nm2. Samples were preconditioned by degassing with continuously flowing dry nitrogen gas at 100° C. prior to analysis.
The pore size distribution (PSD) and cumulative volume of pores is another technique that may be performed from gas adsorption data to lend insight into the sintering behavior of particles. The data was collected by a Micromeritics Tristar II Plus measuring nitrogen adsorption and desorption at 77 K between pressures of
with increments ranging from
Samples were preconditioned by degassing with continuously flowing dry nitrogen gas at 100° C. prior to analysis.
Micromeritics MicroActive software was used to calculate adsorption-desorption PSD and cumulative volume of pores by applying the Barrett, Joyner and Halenda (BJH) method. This method provides a comparative assessment of mesopore size distributions for gas adsorption data. For all BJH data, the Faas correction and Harkins and Jura thickness curve may be applied. The cumulative volume of pores, VPORE (cm3/g), may be measured for both adsorption and desorption portions of the isotherm.
There are a number of exemplary materials described in the present disclosure. To aid in identification and tracking of these exemplary materials, a material naming system has been adopted and is described below. All names of exemplary materials are bolded; N2 herein describes an exemplary material, while N2 refers to nitrogen gas.
Exemplary types of template precursor materials are denoted Sx, where S designates the first one or two letters of the template precursor material (i.e. N for nesquehonite, L for lansfordite, Li for lithium carbonate, C for magnesium citrate, A for amorphous/non-crystalline MgCO3·xH2O, H for hydromagnesite, M for magnesite, E for epsomite, and Ca for calcium carbonate) and where x designates different types of the precursor compound (e.g. H1 and H2 designate two different types of hydromagnesite precursors).
Exemplary types of template materials are named in the format SxTy. The Sx name component designates the precursor type that was utilized to create the template type SxTy, and the Ty name component designates a specific treatment that was utilized to create the template type SxTy. For example, N1T1 and N1T2 indicate two different template types formed from two different treatments on the precursor type N1. We note that while the full SxTy name denotes a specific template type, the Ty name component by itself is only specific with respect to a given Sx precursor type. For example, the treatments utilized to make the template types N1T1 and N2T1 were different, despite these template types sharing the same T1 name component.
Exemplary types of PC materials are named in the format SxTyPz, where the SxTy name component designates the template type and the Pz name component designates a specific type of carbon perimorph. For example, M3T1P1 and M3T1P2 indicate two different PC materials formed from the same M3T1 template material. The Pz name component within the SxTyPZ name is unique—i.e. each Pz name component specifies a unique type of perimorph, irrespective of the SxTy template type utilized to make the perimorph.
Exemplary types of perimorphic frameworks (i.e. the porous perimorphic product resulting from endomorphic extraction) are named in the format Pz, where the Pz name component is not prefaced with an SxTy template type. The Pz name component utilized to name a framework type matches the Pz name component of the SxTyPz PC material type from which the framework type was derived.
The exemplary types of template precursor materials, template materials, perimorpic composite materials, and perimorphic materials in this disclosure are enumerated in
This Section details the generation of exemplary template precursor materials at small scales using exemplary procedures. As such, these procedures comprise partial implementations of the General Method. It should be therefore understood that these procedures must be coupled with other procedures in a full implementation of the General Method. Additionally, it should be understood that these procedures are merely demonstrative of analogous, larger-scale procedures that would be used for industrial-scale manufacturing.
Various techniques may be utilized in the precipitation of precursor materials. For example, the stock solution may be heated to evaporate the process liquid, causing the stock solution to become supersaturated and to precipitate a precursor material. This may be combined with techniques to control the shape and size of the precipitated template precursor particles. For example, the stock solution may be spray-dried to create discrete spheres or hollow spheres. Other techniques may be utilized that will be obvious to those skilled in the art.
Example N1: In an exemplary Precursor Stage procedure, an elongated nesquehonite (MgCO3·3H2O) template precursor material may be derived from a stock solution of aqueous Mg(HCO3)2.
To demonstrate this derivation at a small scale, a representative stock solution may first be prepared. This stock solution represents the stock solution that might be generated during the Separation Stage of a full implementation of the General Method. For this example, a representative aqueous Mg(HCO3)2 stock solution may be generated using water, CO2 gas, and MgO.
First, a 0.24 mol kg−1 Mg mixture comprising deionized water and Akrochem Elastomag 170, a commercial magnesium oxide (MgO) product, may be made. This mixture may be carbonated in a circulation tank with a sparge tube bubbling CO2 to generate carbonic acid. The CO2 bubbling may be discontinued after the MgO is completely dissolved to form the stock solution. The stock solution may be approximately 14.5° C.
Next, air bubbling may be initiated through a sparge tube through the stock solution in the circulation tank at an approximate flow rate of 12 scfmair. This bubbling may cause precipitation of nesquehonite particles and an associated emission of CO2 process gas. Bubbling and circulation may be continued until the conductivity of the solution stabilizes. At this point, the aqueous mixture of nesquehonite particles may be filtered, separating the particles from the aqueous Mg(HCO3)2 filtrate. This filtrate comprises a mother liquor and substantially all of the process water. In a full implementation of the General Method, the separated process water may be conserved for reuse, as shown in
Nesquehonite template precursor particles of the type generated by this procedure may be identified herein as N1 and may be seen in the SEM micrograph in
Aside from the presence of some minor debris, the crystals have smooth, thin surfaces. The elongated morphology of these crystals may be valuable. In applications requiring interlocking particles, such as filtration membranes, an elongated morphology may be useful. In applications requiring the assembly of a percolative network, such as for electron transport, elongated particles may achieve percolation with fewer particles than more equiaxed particle morphologies. In applications requiring mechanical reinforcement, elongated particles may provide superior tensile properties.
Example H1: In another exemplary Precursor Stage procedure, a hierarchical-equiaxed hydromagnesite (Mg5(CO3)4(OH)2·4H2O) template precursor material may be derived from a stock solution of aqueous Mg(HCO3)2.
To demonstrate this derivation at a small scale, a representative stock solution may first be prepared. This stock solution represents the stock solution that might be generated in the Separation Stage of a full implementation of the General Method. For this example, an aqueous Mg(HCO3)2 stock solution with an approximate molality of 0.14 mol kg−1 Mg (aq) may first be prepared as the representative stock solution.
Next, the stock solution may be placed in a 1 L Buchi rotary evaporator vessel, which may then be rotated at 280 RPM in a 100° C. water bath. Crystallization may be allowed to proceed until most of the Mg ions have been precipitated as hydromagnesite precursor particles. Associated with this precipitation, CO2 process gas may be emitted. In a full implementation of the General Method, the CO2 process gas released during precipitation may be conserved using conventional techniques.
The resulting hydromagnesite mixture may then be filtered to separate the solids from the aqueous Mg(HCO3)2 filtrate. This filtrate comprises a mother liquor and substantially all of the process water. In a full implementation of the General Method, the separated mother liquor may be conserved for reuse.
Hydromagnesite template precursor particles of the type generated by this procedure are identified herein as H1 and may be seen in the representative SEM micrographs in
Example H2: In another exemplary Precursor Stage procedure, an elongated, hierarchical hydromagnesite (Mg5(CO3)4(OH)2·4H2O) template precursor material may be derived from a stock solution of aqueous Mg(HCO3)2.
To demonstrate this derivation at a small scale, nesquehonite may first be precipitated from a representative aqueous Mg(HCO3)2 stock solution. This stock solution represents the stock solution that might be generated during the Separation Stage of a full implementation of the General Method. For this example, a representative stock solution and an aqueous mixture of precipitated nesquehonite may be obtained using the procedure described in Example N1. Associated with this nesquehonite precipitation, CO2 process gas may be emitted. In a full implementation of the General Method, the released CO2 process gas may be conserved using conventional techniques.
Next, the nesquehonite mixture may be heated to 100° C. and maintained at that temperature until recrystallization into hydromagnesite is complete. In this exemplary procedure, the process water may be completely evaporated, separating it from the solid residue of elongated hydromagnesite particles. In a full implementation of the General Method, the separated process water may be conserved using conventional techniques.
Hydromagnesite template precursor particles of the type generated by this procedure are identified herein as H2 and may be seen in the representative SEM micrographs in
Example H3: In another exemplary Precursor Stage procedure, a plate-like hydromagnesite (Mg5(CO3)4(OH)2·4H2O) template precursor material may be derived from a stock solution of aqueous Mg(HCO3)2.
To demonstrate this derivation at a small scale, a hierarchical hydromagnesite may first be derived from a representative aqueous Mg(HCO3)2 stock solution. This stock solution represents the stock solution that might be generated during the Separation Stage of a full implementation of the General Method. For this example, a representative stock solution and precipitated hydromagnesite particles may be obtained using the procedure described in Example H2. Associated with the precipitation, CO2 process gas may be emitted. In a full implementation of the General Method, the released CO2 process gas may be conserved using conventional techniques. Additionally, separated process water may be conserved in a full implementation of the General Method.
Next, the hierarchical hydromagnesite particles may be mechanically broken. This might be accomplished in a number of ways using known milling techniques. For the purpose of demonstration, the particles may be slurried in process water. The mixture may then be agitated using high-shear techniques to break the delicate, hierarchical hydromagnesite particles into their constituent, individualized plates. The plate-like hydromagnesite particles may then be filtered from the process water. In a full implementation of the General Method, the separated process water may be conserved for reuse.
Hydromagnesite template precursor particles of the type generated by this procedure are identified herein as H3 and may be seen in the representative SEM micrograph
Example L1: In another exemplary Precursor Stage procedure, an equiaxed lansfordite (MgCO3·5H2O) template precursor material may be derived from a stock solution of aqueous Mg(HCO3)2.
To demonstrate this derivation at a small scale, a representative stock solution may first be prepared. This stock solution represents the stock solution that might be generated in the Separation Stage of a full implementation of the General Method. For this example, a representative aqueous Mg(HCO3)2 stock solution with a concentration of approximately 0.25 mol kg−1 Mg (aq) may be prepared and chilled to 2° C.
The chilled stock solution may then be subjected to N2 bubbling at a flow rate of 4 scfhair. The resulting precipitation may cause CO2 process gas to be emitted. In a full implementation of the General Method, CO2 process gas released during precipitation may be conserved using conventional techniques.
After 67 minutes, N2 bubbling may be discontinued. The crystals formed may be allowed to stir for an additional 50 minutes after discontinuation of N2 bubbling, and the mixture may then be filtered to separate the solids from the mother liquor. The solids may be rinsed with 5° C. deionized water. In a full implementation of the General Method, the separated mother liquor may be conserved for reuse.
Lansfordite template precursor particles of the type generated by this procedure are identified herein as L1 and may be seen in the representative SEM micrograph
Raman spectroscopy may be used to characterize the chemical composition of the template precursor materials. Applying this Raman spectroscopy method results in a match of peak positions consistent with lansfordite at 1083 cm−1, as seen in
Example L2: In another exemplary Precursor Stage procedure, an equiaxed lansfordite (MgCO3·5H2O) template precursor material may derived from a stock solution of aqueous Mg(HCO3)2.
To demonstrate this derivation at a small scale, a representative stock solution may first be prepared. This stock solution represents the stock solution that might be generated in the Separation Stage of a full implementation of the General Method. For this example, a representative stock solution may be obtained as follows. First, an aqueous mixture of precipitated nesquehonite may be obtained using the procedure described in Example N1. The concentration of this mixture may be adjusted to 0.62 mol kg−1 Mg. The mixture may then be added to a high-pressure baffled reactor outfitted with a gas inducing impeller. The system may be stirred at 700 RPM and cooled to 5° C. while injecting CO2 process gas into the reactor's headspace up to a pressure of 850 psi, or until all solids have been dissolved, resulting In the representative, pressurized stock solution.
Upon depressurizing the stock solution to atmospheric pressure, the stirring rate may be reduced to 500 RPM and the solution may be maintained at 12° C. while air is flowed through the headspace. The resulting precipitation of lansfordite particles may cause CO2 process gas to be emitted. In a full implementation of the General Method, CO2 process gas released during precipitation may be conserved using conventional techniques.
After 228 minutes, the mixture of lansfordite particles may be discharged from the reactor and then filtered to separate the lansfordite solids from the mother liquor. In a full implementation of the General Method, the separated mother liquor may be conserved for reuse. For analytical purposes, the lansfordite solids may be rinsed with deionized water, re-suspended in ethanol, filtered again, and dried in a vacuum oven up to 29 inHg at room temperature.
Lansfordite template precursor particles of the type generated by this procedure are identified herein as L2. Raman spectral analysis confirms that the product of this reaction matches that of lansfordite, as seen in
Compared to other equiaxed MgCO3·xH2O-type precursors (e.g. magnesite), lansfordite may be significantly more industrially scalable and less costly. The prismatic, equiaxed morphology may be desirable for applications in which perimorphic products must be integrated with liquids and viscosity effects must be minimized. Additionally, due to lansfordite's relatively high state of hydration, more template precursor volume is generated for a given mass of Mg than is obtainable with less hydrated MgCO3·xH2O, and more template pore volume may be obtained upon decomposition of the precursor material. This can be used to create perimorphic frameworks with more exocellular space.
Example L3: In another exemplary Precursor Stage procedure, an equiaxed, partially dehydrated template lansfordite (MgCO3·5H2O) template precursor material may be derived from a stock solution of aqueous Mg(HCO3)2.
To demonstrate this derivation at small scale, an aqueous lansfordite mixture may first be derived from a representative aqueous Mg(HCO3)2 stock solution. This stock solution represents the stock solution that might be generated during the Separation Stage of a full implementation of the General Method. For this example, a representative stock solution and aqueous lansfordite mixture may be obtained using the procedure described in Example L2. As described in Example L2, the precipitation of lansfordite particles may cause CO2 process gas to be emitted. In a full implementation of the General Method, CO2 process gas released during precipitation may be conserved using conventional techniques.
The concentration of the lansfordite mixture may be adjusted to a 7 wt % concentration of solids. The mixture may then be spray-dried, causing a partial dehydration of the lansfordite material. To demonstrate this at small scale, a Sinoped LPG-5 spray dryer may be used for spray-drying. The lansfordite particles in the 7 wt % mixture may kept continuously suspended via stirring in a vessel. The mixture may be pumped from the vessel at a rate ranging between 116 mL/min and 162 mL/min into the spray dryer's BETE XAER250 air atomizing nozzle. Compressed air may also be delivered into the nozzle at a flow rate ranging between 1.2 scfmair, at 20 psig and 3.6 scfmair at 59 psig. The inlet temperature of the spray dryer may be set to 300° C., producing an outlet temperature ranging between 111° C. and 123° C.
The dry, partially dehydrated lansfordite particles may be collected by a cyclonic particle separator. In a full implementation of the General Method, the process water vapor generated by spray-drying may be conserved using conventional techniques.
The partially dehydrated lansfordite template precursor particles of the type generated by this procedure are identified herein as L3. Process liquids and gases may be recovered through typical industrial methods for reuse in Separation Stage.
The TGA mass loss of 67.1% for an L3 template precursor material generated according to the procedure described above confirms that partial dehydration occurred (the theoretical mass loss for lansfordite is 76.9%, as shown in
Example M1: In another exemplary Precursor Stage procedure, an equiaxed magnesite (MgCO3) template precursor material may be derived from a stock solution of aqueous Mg(HCO3)2.
To demonstrate this derivation at a small scale, a representative stock solution may first be prepared. This stock solution represents the stock solution that might be generated in the Separation Stage of a full implementation of the General Method. For this example, a representative aqueous Mg(HCO3)2 stock solution with a concentration of 0.25 mol kg−1 Mg (aq) may be prepared. This stock solution may then be slurried with additional MgO to provide more Mg ions. In a full implementation of the General Method, MgCO3·xH2O precipitated from a stock solution might be utilized to provide more Mg ions. However, for the purpose of this demonstration, the additional MgO may comprise a commercial MgO product (Elastomag 170) that has been calcined at 1050° C. for 1 hour. With this additional loading of Mg ions, the total Mg present in the stock solution-mixture may be 1.5 mol kg−1 Mg.
Next, this stock solution-mixture may be placed in a pressure vessel with magnetic stirring, a high-pressure gas inlet, and a purging needle valve. CO2 may be flowed for 2 minutes to purge the vessel of air, after which it may be fully sealed and pressurized with CO2 to 725 psi at 14.4° C. The vessel may be heated on a heating stir plate. Under magnetic stirring and heating, after 291 minutes, the vessel may reach 193.7° C. and 975 psi. Inside the vessel, magnesite is precipitated During this thermal treatment, and CO2 process gas may be emitted into the vessel's headspace. The vessel may then be depressurized and allowed to cool over the course of 30 minutes, releasing steam and CO2 continuously. In a full implementation of the General Method, the CO2 process gas released during precipitation and subsequent depressurization may be conserved using conventional techniques.
The mixture of magnesite particles may then be discharged from the vessel and then filtered to separate the magnesite solids from the mother liquor. In a full implementation of the General Method, the separated mother liquor may be conserved for reuse in the Separation Stage. The magnesite may be dried at 100° C.
The magnesite template precursor material of the type generated by this procedure are identified herein as M1. The particles display an equiaxed rhombohedral morphology and are shown in the SEM micrograph in
Example M2: In another exemplary Precursor Stage procedure, an equiaxed magnesite (MgCO3) template precursor material may be derived from a stock solution of aqueous Mg(HCO3)2.
To demonstrate this derivation at small scale, nesquehonite may be generated from a stock solution of aqueous Mg(HCO3)2 using the procedure described in Example N1. This precipitation may cause CO2 process gas to be emitted. In a full implementation of the General Method, the CO2 process gas released during precipitation may be conserved using conventional techniques. Likewise, the separated mother liquor may be conserved in a full implementation of the General Method.
In this exemplary procedure, the nesquehonite may then be combined with water to make a mixture with a concentration of 1.5 mol kg−1 Mg. The mixture may be placed in a pressure vessel with magnetic stirring, a high-pressure gas inlet, and a purging needle valve. The headspace of the pressure vessel may contain ambient pressure air, with no additional gas input. The pressure vessel may then be sealed.
The mixture may be magnetically stirred in the vessel for 10 minutes. Then, the vessel may be heated to 175° C. over 68 minutes. The reaction temperature may fluctuate During this thermal treatment, reaching a maximum temperature of 180° C. and a maximum pressure of 1190 psi, at which condition any CO2 liberated from nesquehonite in the reaction may be rendered supercritical. The pressure vessel may be then be allowed to cool for 199 minutes.
The resulting mixture of magnesite particles may be discharged from the vessel and then filtered to separate the magnesite solids from the mother liquor. In a full implementation of the General Method, the separated mother liquor may be conserved for reuse in the Separation Stage. The magnesite may be dried at 100° C.
The magnesite template precursor material of the type generated by this procedure are identified herein as M2. The particles display an equiaxed rhombohedral morphology and are shown in the SEM micrograph in
Example A1: In another exemplary Precursor Stage procedure, a hollow non-crystalline MgCO3·xH2O template precursor material may be derived from a stock solution of aqueous Mg(HCO3)2.
To demonstrate this derivation at a small scale, a representative stock solution may first be prepared. This stock solution represents the stock solution that might be generated in the Separation Stage of a full implementation of the General Method. For this example, a representative aqueous Mg(HCO3)2 stock solution with a concentration of 0.43 mol kg−1 Mg (aq) may be prepared. This may be done by mixtureing a commercial MgCO3·xH2O product (“Light Magnesium Carbonate” supplied by Akrochem Corporation) in water at a solids concentration equivalent to 0.43 mol kg−1 Mg. This mixture may be carbonated using pressurized CO2 gas in a circulated pressure vessel. The system may be pressurized by injecting CO2 gas into the vessel to a total pressure of 555 psi. This may be maintained for 2 hours and 13 minutes at 34° C. or until all solids are dissolved. At this point, the vessel may be depressurized and stored under atmospheric pressure at 4° C.
The chilled stock solution may then be spray-dried. To demonstrate this, the stock solution may be pumped at a rate of 35 mL/min through a BETE XAER150 air atomizing nozzle of a Sinoped LPG-5 spray dryer. Compressed air may be delivered into the nozzle at a flow rate of 2.8 scfmair at 45 psig. The inlet temperature of the spray dryer may be set to 165° C., resulting in an outlet temperature of 110° C.
The particles resulting from spray-drying the stock solution may be collected by a cyclonic particle separator. In a full implementation of the General Method, both the process water vapor and the CO2 process gas emitted by spray-drying may be conserved using conventional techniques.
The type of MgCO3·xH2O template precursor material resulting from this process is identified herein as A1. SEM image analysis of A1 particles, as shown in the SEM micrographs in
Raman spectral analysis showed that the product of this reaction does have a Raman peak that may be associated with crystalline carbonate, located at 1106 cm−1. However, it does not match with any of the typical MgCO3·xH2O peaks (
Example A2: In another exemplary Precursor Stage procedure, a hollow, hierarchical-equiaxed, non-crystalline MgCO3·xH2O template precursor material may be derived from a stock solution of aqueous Mg(HCO3)2.
To demonstrate this derivation at a small scale, a representative stock solution may first be prepared. This stock solution represents the stock solution that might be generated in the Separation Stage of a full implementation of the General Method. For this example, a representative aqueous Mg(HCO3)2 stock solution with a concentration of 1.39 mol kg−1 Mg (aq) may be prepared. This may be done by mixtureing a commercial Mg(OH)2 product (“Versamag” supplied by Akrochem Corporation) in water at a solids concentration equivalent to 1.49 mol kg−1 Mg. This mixture may be carbonated using pressurized CO2 gas in a circulated pressure vessel. The system may be pressurized by injecting CO2 gas into the vessel to a total pressure between 700-800 psig. This may be maintained for 2 hours at 10° C. or until substantially all (i.e. >90%) solids are dissolved. At this point, the contents may be depressurized and stored under atmospheric pressure between 4-10° C.
The stock solution may then be spray-dried. To demonstrate this, the stock solution may be pumped at a rate of 2.7 mL/min through a 0.7 mm Buchi B-290 two fluid air atomizing nozzle in a Buchi B-191 spray drying system. Compressed air may be delivered into the nozzle at a flow rate of 0.6 scfmair at 88 psig. The inlet temperature of the spray dryer may be set to 130° C., resulting in an outlet temperature between 85-89° C. The aspirator may be set to 18 scfmair.
The particles resulting from spray-drying the stock solution may be collected by a cyclonic particle separator. In a full implementation of the General Method, both the process water vapor and the CO2 process gas emitted by spray-drying may be conserved using conventional techniques.
The type of MgCO3·xH2O template precursor material resulting from this process is identified herein as A2. SEM image analysis of A2 particles, as shown in the SEM micrographs in
Raman spectral analysis showed that the MgCO3·xH2O spheres have no distinct Raman peak that may be associated with crystalline carbonate. Additionally, TGA analysis of the template precursor fails to match with the common crystalline forms of MgCO3·xH2O with a mass loss of 68.4%, as seen in
Example A3: In another exemplary Precursor Stage procedure, a hollow, hierarchical-equiaxed, non-crystalline MgCO3·xH2O template precursor material may be derived from a stock solution of aqueous Mg(HCO3)2.
To demonstrate this derivation at a small scale, a representative stock solution may first be prepared. This stock solution represents the stock solution that might be generated in the Separation Stage of a full implementation of the General Method. For this example, a representative aqueous Mg(HCO3)2 stock solution with a concentration of 1.08 mol kg−1 Mg (aq) may be prepared. This may be done by mixtureing a commercial Mg(OH)2 product (“Versamag” supplied by Akrochem Corporation) in water at a solids concentration equivalent to 1.12 mol kg−1 Mg. This mixture may be carbonated using pressurized CO2 gas in a circulated pressure vessel. The system may be pressurized by injecting CO2 gas into the vessel to a total pressure between 700-800 psig. This may be maintained for 2 hours at 10° C. or until substantially all solids (i.e. >90%) are dissolved. At this point, the contents may be depressurized and stored under atmospheric pressure between 4-10° C.
The stock solution may then be spray-dried. To demonstrate this, the stock solution may be pumped at a rate of 2.7 mL/min through a 0.7 mm Buchi B-290 two fluid air atomizing nozzle in a Buchi B-191 spray drying system. Compressed air may be delivered into the nozzle at a flow rate of 0.6 scfmair at 88 psig. The inlet temperature of the spray dryer may be set to 90° C., resulting in an outlet temperature between 56-58° C. The aspirator may be set to 18 scfmair.
The particles resulting from spray-drying the stock solution may be collected by a cyclonic particle separator. In a full implementation of the General Method, both the process water vapor and the CO2 process gas emitted by spray-drying may be conserved using conventional techniques.
The type of MgCO3·xH2O template precursor material resulting from this process is identified herein as A3. SEM image analysis of A3 particles, as shown in the SEM micrographs in
The macropores are located throughout the shell, which can be seen in the carbon perimorphic frameworks grown on them.
Raman spectral analysis showed that the MgCO3·xH2O spheres have no distinct Raman peak that may be associated with crystalline carbonate. Additionally, TGA analysis of the template precursor fails to match with the common crystalline forms of MgCO3·xH2O with a mass loss of 72.9%, as seen in
Example C1: In another exemplary Precursor Stage procedure, a hollow, hierarchical-equiaxed magnesium citrate template precursor material may be derived from a stock solution of aqueous magnesium citrate.
To demonstrate this derivation at a small scale, a representative stock solution may first be prepared. This stock solution represents the stock solution that might be generated in the Separation Stage of a full implementation of the General Method. For this example, a representative aqueous magnesium citrate stock solution with a concentration of 0.52 mol kg−1 Mg (aq) may be prepared by reacting citric acid (supplied by Sigma Aldrich) with a 0.52 mol kg−1 aqueous mixture of Mg(OH)2 (Versamag, supplied by Akrochem).
The stock solution may then be spray-dried. To demonstrate this, the stock solution may be pumped at a rate of 3.75 mL/min through a Buchi B-290 two-fluid nozzle of a Buchi B-191 spray dryer. Compressed air may be delivered into the nozzle at a flow rate of 0.6 scfmair at 88 psig with the aspirator airflow set to 18 scfmair. The inlet temperature may be set to 220° C., resulting in an outlet temperature of 110° C.
The particles resulting from spray-drying the stock solution may be collected by a cyclonic particle separator. In a full implementation of the General Method, the process water vapor emitted by spray-drying may be conserved using conventional techniques.
The type of magnesium citrate template precursor material resulting from this process is identified herein as C1. SEM analysis of C1 particles, as shown in the SEM micrographs in
Raman spectral analysis confirms that the product of this reaction matches that of magnesium citrate, as shown in
Example E1: In another exemplary Precursor Stage procedure, an elongated template precursor material of epsomite (magnesium sulfate heptahydrate, MgSO4·7H2O) may be derived from an aqueous stock solution of magnesium sulfate.
To demonstrate this derivation at a small scale, a representative stock solution may first be prepared. This stock solution represents the stock solution that might be generated in the Separation Stage of a full implementation of the General Method. For this example, a representative aqueous magnesium sulfate stock solution with a concentration of 4.06 mol kg−1 Mg (aq) may be prepared by dissolving epsomite in water at room temperature. This may be done in a glass beaker magnetically stirred at 700 RPM.
Once dissolved, 410.86 g of acetone may be added dropwise by a separatory funnel, which may result in immediate crystal formation in the solution. While this represents an antisolvent precipitation, which is generally undesirable, solventless precipitation of epsomite could easily be accomplished by chilling or spray-drying the stock solution. Moreso than demonstrating an engineered precursor morphology or demonstrating scalable procedure, the purpose of the Example E1 procedure was moreso just to precipitate epsomite, so that the template materials and perimorphic materials derived from an epsomite precursor compound might be demonstrated and analyzed in subsequent sections of the current disclosure. In a full implementation of the General Method, the mother liquor separated after a solventless precipitation may be conserved for reuse in the Separation Stage.
After 22 minutes, the precipitation of the epsomite may be complete. The resulting mixture may be collected and filtered. The particles may be dried.
The type of epsomite template precursor material resulting from this process is identified herein as E1. The particles may be observed via optical microscope as elongated rods with hexagonal cross sections, as shown in
Raman spectral analysis confirms that the product of this reaction matches that of epsomite, as seen in
Example H4: In another exemplary Precursor Stage procedure, a Li-doped hydromagnesite (Mg5(CO3)4(OH)2·4H2O) template precursor material may be derived from an aqueous stock solution of Mg(HCO3)2 that also contains a small concentration of aqueous Li2CO3.
To demonstrate this derivation at a small scale, a representative stock solution may first be prepared. This stock solution represents the stock solution that might be generated in the Separation Stage of a full implementation of the General Method. For this example, a representative aqueous Mg(HCO3)2 stock solution may be prepared, with the additional step of adding lithium carbonate (Li2CO3). This may be done as follows. First, an MgO powder (Akrochem Elastomag 170 calcined at 1050° C. for 1 hour) may be slurried into water at a solids concentration of 0.23 mol kg−1 Mg (s). This may be done in a glass beaker with magnetic stirring. To this mixture, Li2CO3 (Sigma Aldrich) may be added at a solids concentration of 2.71·10−3 mol kg−1 Li (s). The mixture may be carbonated with a sparge tube bubbling CO2 gas to generate aqueous H2CO3. The CO2 flow may be discontinued after the MgO and Li2CO3 are completely dissolved. The Mg(HCO3)2 stock solution may then be filtered to remove any residual undissolved impurities.
Next, the stock solution may be heated to 100° C. in an uncovered glass beaker with magnetic stirring. This condition may be maintained for 2 hours, during which hydromagnesite particles may be precipitated. After 2 hours, the resulting mixture may be filtered, and the solid hydromangesite may be dried in a forced air circulation at 100° C.
The type of Li-doped hydromagnesite template precursor material resulting from this process is identified herein as H4. The particles are shown in the SEM micrographs of
Example H5: In an exemplary Precursor Stage procedure, a Li-doped hydromagnesite (Mg5(CO3)4(OH)2·4H2O) template precursor material may be derived from an aqueous stock solution of Mg(HCO3)2 that also contains a moderate concentration of Li2CO3.
To demonstrate this derivation at a small scale, a representative stock solution may first be prepared. This stock solution represents the stock solution that might be generated in the Separation Stage of a full implementation of the General Method. For this example, a representative aqueous Mg(HCO3)2 stock solution may be prepared, with the additional step of adding lithium carbonate (Li2CO3). This may be done as follows.
First, an MgO powder (Akrochem Elastomag 170 calcined at 1050° C. for 1 hour) may be slurried into water at a solids concentration of 0.23 mol kg−1 Mg (s). This may be done in a glass beaker with magnetic stirring. To this mixture, Li2CO3 (Sigma Aldrich) may be added at a solids concentration of 2.74·10−2 mol kg−1 Li (s). The mixture may be carbonated with a sparge tube bubbling CO2 gas to generate aqueous H2CO3. The CO2 flow may be discontinued after the MgO and Li2CO3 are completely dissolved. The Mg(HCO3)2 stock solution may then be filtered to remove any residual undissolved impurities.
Next, the stock solution may be heated to 100° C. in an uncovered glass beaker with magnetic stirring. This condition may be maintained for 1 hour, during which hydromagnesite particles may be precipitated. After 1 hour, the resulting mixture may be filtered, and the solid hydromangesite may be dried in a forced air circulation at 100° C.
The type of Li-doped hydromagnesite template precursor material resulting from this process is identified herein as H5. The particles are shown in the SEM micrographs of
Example M3: In another exemplary Precursor Stage procedure, an equiaxed magnesite (MgCO3) template precursor material may be derived from a stock solution of aqueous Mg(HCO3)2.
To demonstrate this derivation at a small scale, a representative stock solution may first be prepared. This stock solution represents the stock solution that might be generated in the Separation Stage of a full implementation of the General Method. For this example, the stock solution may be generated in a high-pressure reactor. First, a commercial hydromagnesite product (Akrochem Light Magnesium Carbonate) may be slurried in water at a solids concentration of 0.74 mol kg−1 Mg (s). This mixture may be placed in a circulated pressure vessel. The sealed vessel may then be heated to 145° C., at which temperature ˜800 psi of gaseous CO2 may be introduced into the system. This reaction may continue to recirculate at 145° C. for a duration of 139 minutes, reaching a maximum pressure of 900 psi. During this thermal treatment, the hydromagnesite may be dissolved, forming aqueous Mg(HCO3)2, and magnesite may be precipitated from the Mg(HCO3)2. At this point, the vessel may be depressurized, releasing CO2 process gas. In a full implementation of the General Method, the CO2 process gas released during precipitation and subsequent depressurization may be conserved using conventional techniques.
The resulting mixture of magnesite particles may be discharged from the vessel and then filtered to separate the magnesite solids from the mother liquor. In a full implementation of the General Method, the separated mother liquor may be conserved for reuse in the Separation Stage. The magnesite may be dried at 100° C.
The type of magnesite template precursor material resulting from this process is identified herein as M3. The equiaxed magnesite particles may be seen in the SEM micrograph of
Example M4: In another exemplary Precursor Stage procedure, an equiaxed magnesite (MgCO3) template precursor material may be derived from a stock solution of aqueous, Na-rich Mg(HCO3)2.
To demonstrate this derivation at a small scale, a representative stock solution may first be prepared. This stock solution represents the stock solution that might be generated in the Separation Stage of a full implementation of the General Method. For this example, the stock solution may be generated in a high-pressure reactor. First, a commercial hydromagnesite product (Akrochem Light Magnesium Carbonate) may be slurried in water at a solids concentration of 0.74 mol kg−1 Mg. To this mixture, a commercial NaHCO3 product (Arm & Hammer) may be added at a concentration of 2.17·10−3 mol kg−1 Na. This mixture may be placed in a circulated pressure vessel. The sealed vessel may then be heated to 145° C. upon which ˜800 psi of gaseous CO2 may be introduced into the system. This reaction may continue to recirculate at 145° C. for a duration of 135 minutes, reaching a maximum pressure of 840 psi. During this thermal treatment, the hydromagnesite may be dissolved, forming aqueous Mg(HCO3)2, and magnesite may be precipitated from the aqueous Mg(HCO3)2. At this point, the vessel may be depressurized, releasing CO2 process gas. In a full implementation of the General Method, the CO2 process gas released during precipitation and subsequent depressurization may be conserved using conventional techniques.
The resulting mixture of magnesite particles may be discharged from the vessel and then filtered to separate the magnesite solids from the mother liquor. In a full implementation of the General Method, the separated mother liquor may be conserved for reuse in the Separation Stage. The magnesite may be dried at 100° C.
The type of magnesite template precursor material resulting from this process is identified herein as M4. The equiaxed magnesite particles may be seen in the SEM micrograph of
Example M5: In another exemplary Precursor Stage procedure, an equiaxed magnesite (MgCO3) template precursor material may be derived from a stock solution of aqueous, Na-rich Mg(HCO3)2.
To demonstrate this derivation at a small scale, a representative stock solution may first be prepared. This stock solution represents the stock solution that might be generated in the Separation Stage of a full implementation of the General Method. For this example, the stock solution may be generated in a high-pressure reactor. First, a commercial hydromagnesite product (Akrochem Light Magnesium Carbonate) may be slurried in water at a solids concentration of 0.74 mol kg−1 Mg. To this mixture, a commercial NaHCO3 product (Arm & Hammer) may be added at a concentration of 0.19 mol kg−1 Na. This mixture may be placed in a circulated pressure vessel. The sealed vessel may then be heated to 145° C. upon which ˜800 psi of gaseous CO2 may be introduced into the system. This reaction may continue to recirculate at 145° C. for a duration of 137 minutes, reaching a maximum pressure of 850 psi. During this thermal treatment, the hydromagnesite may be dissolved, forming aqueous Mg(HCO3)2, and magnesite may be precipitated from the aqueous Mg(HCO3)2. At this point, the vessel may be depressurized, releasing CO2 process gas. In a full implementation of the General Method, the CO2 process gas released during precipitation and subsequent depressurization may be conserved using conventional techniques.
The resulting mixture of magnesite particles may be discharged from the vessel and then filtered to separate the magnesite solids from the mother liquor. In a full implementation of the General Method, the separated mother liquor may be conserved for reuse in the Separation Stage. The magnesite may be dried at 100° C.
The type of magnesite template precursor material resulting from this process is identified herein as M5. The equiaxed magnesite particles may be seen in the SEM micrograph of
Comparing M3, M4, and M5, there are no appreciable morphological differences that can be readily identified based on SEM analysis.
Example N2: In another exemplary Precursor Stage procedure, an elongated nesquehonite (MgCO3·3H2O) template precursor material may be derived from an aqueous stock solution of Mg(HCO3)2.
To demonstrate this derivation at small scale, lansfordite may first be generated from a stock solution of aqueous Mg(HCO3)2 using the procedure described in Example L2. This precipitation may cause CO2 process gas to be emitted. In a full implementation of the General Method, the CO2 process gas released during precipitation may be conserved using conventional techniques. Likewise, the separated mother liquor may be conserved in a full implementation of the General Method.
Next, water may be heated to 35° C. in a glass beaker. Once the water has reached temperature, the lansfordite may be added to produce a mixture with a concentration of 0.74 mol kg−1 Mg. The mixture may be magnetically stirred at 600 RPM and maintained at 35° C. for 100 minutes. During this thermal treatment, the lansfordite may be dissolved and nesquehonite may be precipitated. The mixture may then be filtered to separate the mother liquor from the lansfordite. In a full implementation of the General Method, the separated mother liquor may be conserved.
The type of nesquehonite template precursor material resulting from this process is identified herein as N2. Optical micrographs are shown in
Example N3: In another exemplary Precursor Stage procedure, an elongated nesquehonite (MgCO3·3H2O) template precursor material may be derived from an aqueous stock solution of Mg(HCO3)2.
To demonstrate this derivation at small scale, lansfordite may first be generated from a stock solution of aqueous Mg(HCO3)2 using the procedure described in Example L2. This precipitation may cause CO2 process gas to be emitted. In a full implementation of the General Method, the CO2 process gas released during precipitation may be conserved using conventional techniques. Likewise, the separated mother liquor may be conserved in a full implementation of the General Method.
Next, a 10.84 mM aqueous solution of SDS (TCI Chemical) may be heated to 35° C. in a glass beaker. Once the water has reached temperature, the lansfordite may be added to produce a mixture with a concentration of 0.74 mol kg−1 Mg. The mixture may be magnetically stirred at 600 RPM and maintained at 35° C. for 100 minutes. During this thermal treatment, the lansfordite may be dissolved and nesquehonite may be precipitated. The mixture may then be filtered to separate the mother liquor from the lansfordite. In a full implementation of the General Method, the separated mother liquor may be conserved.
The type of nesquehonite template precursor material resulting from this process is identified herein as N3. Optical micrographs are shown in
Example Li1: In another exemplary Precursor Stage procedure, a hollow, hierarchical-equiaxed Li2CO3 template precursor material may be derived from a stock solution of aqueous Li2CO3.
To demonstrate this derivation at a small scale, a representative stock solution may first be prepared. This stock solution represents the stock solution that might be generated in the Separation Stage of a full implementation of the General Method. For this example, a representative aqueous Li2CO3 stock solution may be prepared as follows. First, a commercial Li2CO3 product (supplied by FMC) may be slurried in water at a concentration of 0.54 mol kg−1 Li. This mixture may be carbonated in an overhead stirred reactor fitted with a gas dispersing blade and a sparge tube. CO2 gas may be flowed into the mixture through the sparge tube at a rate of 9 sfchair, for 175 minutes or until the solids are completely dissolved. At this point, the solution may be diluted with water to adjust the concentration to 0.27 mol kg−1 Li (aq).
This representative stock solution may then be spray-dried. To demonstrate this, the stock solution may be pumped at a rate of 7 mL/min through a Buchi B-290 two-fluid nozzle of a Buchi B-191 spray dryer. Compressed air may be delivered into the nozzle at a flow rate of 0.6 scfmair, at 88 psig with the aspirator airflow set to 18 scfmair. The inlet temperature may be set to 170° C., resulting in an outlet temperature of 100° C.
The particles resulting from spray-drying the stock solution may be collected by a cyclonic particle separator. In a full implementation of the General Method, the CO2 process gas and process water vapor emitted by spray-drying may be conserved using conventional techniques.
The type of lithium carbonate template precursor material resulting from this process is identified herein as Li1. The particles are hollow, hierarchical-equiaxed structures, as seen in the SEM micrographs of
This Section details the generation of exemplary template materials at small scales using exemplary procedures. As such, these procedures comprise partial implementations of the General Method. It should be therefore understood that these procedures must be coupled with other procedures in a full implementation of the General Method. Additionally, it should be understood that these procedures are merely demonstrative of analogous, larger-scale procedures that would be used for industrial-scale manufacturing.
A number of exemplary procedures for making template materials are described in this section. In some exemplary procedures, template precursor materials may be treated to form template materials in a separate and distinct Template Stage procedure, and the resulting template materials may then be utilized in a separate and distinct Replication Stage procedure. In other instances, the Template Stage and the Replication Stage procedures may both be performed in the same reactor. Some of these exemplary Template Stage procedures utilize template precursor materials previously named and described in Section V. Additionally, new template precursor materials have been utilized.
Example N1T1: In an exemplary Template Stage procedure, a nesquehonite template precursor material may be thermally treated to form a porous MgO template material.
To demonstrate this, N1-type nesquehonite particles may first be generated using the procedure described in Example N1. This material represents the template precursor material that might be generated in the Precursor Stage of a full implementation of the General Method.
Next, the template precursor material may be thermally treated. This may be performed in a TGA instrument under an inert gas flow of Ar as described in Scheme E in Section III. The sample of N-type nesquehonite particles may be heated under Ar gas from room temperature to a final temperature of 1,000° C. at a rate of 10° C./min. During this thermal treatment, CO2 gas may be released. In a full implementation of the General Method, the CO2 process gas released during the decomposition of the template precursor material may be conserved using conventional techniques. Upon reaching 1,000° C., the sample may be cooled back down to room temperature.
The type of porous MgO template material resulting from this process is identified herein as N1T1. The template particles retain the precursor particles' elongated superstructure, as shown in the SEM micrographs of
Example H1T1: In another exemplary Template Stage procedure, a hydromagnesite template template precursor material may be thermally treated to form a porous MgO template material.
To demonstrate this, H1-type hydromagnesite particles may first be generated using the procedure described in Example H1. This material represents the template precursor material that might be generated in the Precursor Stage of a full implementation of the General Method.
Next, the template precursor material may be thermally treated. This may be performed in a TGA instrument under an inert gas flow of Ar as described in Scheme E in Section III. The sample of H1-type hydromagnesite particles may be heated under Ar gas from room temperature to a final temperature of 1,000° C. at a rate of 10° C./min. During this thermal treatment, CO2 gas may be released. In a full implementation of the General Method, the CO2 process gas released during the decomposition of the template precursor material may be conserved using conventional techniques. Upon reaching 1,000° C., the sample may be allowed to cool to room temperature.
The type of porous MgO template material resulting from this process is identified herein as H1T1. The template particles retain the precursor particles' hierarchical-equiaxed, rosette superstructure, as shown in the SEM micrographs of
Example H2T1: In another exemplary Template Stage procedure, a hydromagnesite template template precursor material may be thermally treated to form a porous MgO template material.
To demonstrate this, H2-type hydromagnesite particles may first be generated using the procedure described in Example H2. This material represents the template precursor material that might be generated in the Precursor Stage of a full implementation of the General Method.
Next, the template precursor material may be thermally treated. This may be performed according to Scheme B in a tube furnace as detailed in Section III. The sample of H2-type hydromagnesite particles may be placed in a ceramic boat and introduced into a tube furnace at room temperature. The furnace may then be heated under Ar flow of 2000 sccm to 1050° C. at a heating rate of 20° C./min. During this thermal treatment, CO2 gas may be released. In a full implementation of the General Method, the CO2 process gas released during the decomposition of the template precursor material may be conserved using conventional techniques. The sample may then be maintained at 1050° C. for two hours, after which the furnace may be allowed to cool to room temperature.
The type of porous MgO template material resulting from this process is identified herein as H2T1. The template particles retain the precursor particles' elongated, rosette superstructure, as shown in the SEM micrographs of
Some of the subunits observed in
During the thermal treatment in Example H2T1, the porous MgO template material derived from decomposition of the template precursor may undergo grain growth and sintering due to atomic diffusion. The distance over which diffusion may occur may be a function of the temperature. Hence, modulating the temperature and duration of the Template Stage treatment may be useful for fine engineering of a template's substructure (and accordingly of a perimorphic framework's substructure).
During coarsening, the porous substructure of the template materials may also be densified. This may affect the fractional composition of positive and negative template space. Taken to an extreme, densification of the porous substructure may continue until the negative space—i.e. the template's pore structure—is eliminated. As particles sinter to one another, higher-order porosity may be obtained via the pores between these formerly discrete particles. This technique has been utilized by workers to create template structures comprising macroscopic, porous networks of sintered metal oxide particles. Macroscopic, monolithic template structures like this can be formed in Template Stage and recycled using the General Method.
Example H1T2: In another exemplary Template Stage procedure, a hydromagnesite template template precursor material may be thermally treated to form an MgO template material.
To demonstrate this, H1-type hydromagnesite particles may first be generated using the procedure described in Example H1. This material represents the template precursor material that might be generated in the Precursor Stage of a full implementation of the General Method.
Next, the template precursor material may be thermally treated. This may be performed in a TGA according to Scheme E as detailed in Section III. The sample of H-type hydromagnesite particles may be heated under Ar gas from room temperature to a final temperature of 1200° C. at a heating rate of 10° C./min, during which CO2 gas may be released. In a full implementation of the General Method, the CO2 process gas released during the decomposition of the template precursor material may be conserved using conventional techniques. The sample may be held at 1200° C. for 10 minutes, then allowed to cool.
The type of MgO template material resulting from this process is identified herein as H1T2. The template particles resulting from this procedure are shown in the SEM micrographs of
Example N1T2: In another exemplary Template Stage procedure, a nesquehonite template precursor material may be thermally treated to form an MgO template material.
To demonstrate this, N1-type nesquehonite particles may first be generated using the procedure described in Example N1. This material represents the template precursor material that might be generated in the Precursor Stage of a full implementation of the General Method.
Next, the template precursor material may be thermally treated. This may be performed in a TGA according to Scheme E under an Ar flow from room temperature to a final temperature of 1200° C. at a heating rate of 10° C./min. During this thermal treatment, CO2 gas may be released. In a full implementation of the General Method, the CO2 process gas released during the decomposition of the template precursor material may be conserved using conventional techniques. The sample may be held at 1200° C. for 10 minutes, then allowed to cool.
The type of MgO template material resulting from this process is identified herein as N1T2. The template particles have lost the porous substructure evolved during thermal decomposition due to progressive sintering at high temperature.
Example N1T3: In another exemplary Template Stage procedure, a nesquehonite template precursor material may be thermally treated to form a porous MgO template material.
To demonstrate this, N1-type nesquehonite particles may first be generated using the procedure described in Example N1. This material represents the template precursor material that might be generated in the Precursor Stage of a full implementation of the General Method.
Next, the template precursor material may be thermally treated. This may be performed in a tube furnace according to Scheme B, as detailed in Section III. The sample may be heated from room temperature to 460° C. under Ar gas flow of 1271 sccm. At this point, acetylene (C2H2) gas may be introduced into the system to begin depositing carbon the templating surface. During this Replication Stage procedure, the template, which may not have completed its thermal decomposition, may continue decomposing in the high temperature environment and CO2 gas may be released. In a full implementation of the General Method, the CO2 process gas released during the decomposition of the template precursor material may be conserved using conventional techniques. This condition may be maintained for 3 hours. Acetylene flow may be terminated and the furnace may then be allowed to cool to room temperature under sustained Ar flow.
The type of porous MgO template material resulting from this process is identified herein as N1T3.
Example M1T1: In another exemplary Template Stage procedure, a magnesite template precursor material may be thermally treated to form a porous MgO template material.
To demonstrate this, M1-type magnesite particles may first be generated using the procedure described in Example M1. This material represents the template precursor material that might be generated in the Precursor Stage of a full implementation of the General Method.
Next, the template precursor material may be thermally treated. This may be performed in a TGA according to Scheme E as detailed in Section III. The sample may be heated from room temperature to 1050° C. at a rate of 50° C./min under Ar flow. During this thermal treatment, CO2 gas may be released. In a full implementation of the General Method, the CO2 process gas released during the decomposition of the template precursor material may be conserved using conventional techniques. The sample may be held at 1050° C. for 1 minute, then allowed to cool.
The type of porous MgO template material resulting from this process is identified herein as M1T1. The template particles retain the precursor particles' equiaxed superstructure, as shown in the SEM micrographs of
Example M1T2 In another exemplary Template Stage procedure, a magnesite template precursor material may be thermally treated to form a porous MgO template material.
To demonstrate this, M1-type magnesite particles may first be generated using the procedure described in Example M1. This material represents the template precursor material that might be generated in the Precursor Stage of a full implementation of the General Method.
Next, the template precursor material may be thermally treated. This may be performed in a tube furnace according to Scheme B, as detailed in Section III. The sample may be heated from room temperature to a final temperature of 1050° C. at a heating rate of 20° C./min and under an Ar flow of 2360 sccm. During this thermal treatment, CO2 gas may be released. In a full implementation of the General Method, the CO2 process gas released during the decomposition of the template precursor material may be conserved using conventional techniques. The sample may then be held at 1050° C. for 4 hours, then allowed to cool.
The type of MgO template material resulting from this process is identified herein as M1T2. The template particles retain the precursor particles' equiaxed superstructure, as shown in the SEM micrographs of
Example M1T3: In another exemplary Template Stage procedure, a magnesite template precursor material may be thermally treated to form a porous MgO template material.
To demonstrate this, M1-type magnesite particles may first be generated using the procedure described in Example M1. This material represents the template precursor material that might be generated in the Precursor Stage of a full implementation of the General Method.
Next, the template precursor material may be thermally treated. This may be performed in a TGA according to Scheme E, as detailed in Section III. The sample may be heated from room temperature to a final temperature of 1200° C. at a rate of 50° C./min under flowing Ar. During this thermal treatment, CO2 gas may be released. In a full implementation of the General Method, the CO2 process gas released during the decomposition of the template precursor material may be conserved using conventional techniques. The sample may then be held at 1200° C. for 1 minute, then allowed to cool.
The type of MgO template material resulting from this process is identified herein as M1T3. The template particles retain the precursor particles' equiaxed superstructure, as shown in the SEM micrographs of
Example M1T4: In another exemplary Template Stage procedure, a magnesite template precursor material may be thermally treated to form a porous MgO template material.
To demonstrate this, M1-type magnesite particles may first be generated using the procedure described in Example M1. This material represents the template precursor material that might be generated in the Precursor Stage of a full implementation of the General Method.
Next, the template precursor material may be thermally treated. This may be performed in a tube furnace according to Scheme B, as detailed in Section III. The sample may be heated from room temperature to a final temperature of 1200° C. at a heating rate of 20° C./min under an Ar flow of 2000 sccm. During this thermal treatment, CO2 gas may be released. In a full implementation of the General Method, the CO2 process gas released during the decomposition of the template precursor material may be conserved using conventional techniques. The sample may then be held at 1200° C. for 4 hours, then allowed to cool.
The type of MgO template material resulting from this process is identified herein as M1T4. The template particles retain the precursor particles' equiaxed superstructure, as shown in the SEM micrographs of
Example E1T1: In another exemplary Template Stage procedure, an epsomite template precursor material may be thermally treated to form a dehydrated, basic MgSO4 template material.
To demonstrate this, epsomite particles may first be generated. The epsomite particles used in this exemplary procedure are generated as described in Example E1. This material represents the template precursor material that might be generated in the Precursor Stage of a full implementation of the General Method.
Next, the template precursor material may be thermally treated. This may be performed in a forced air circulation oven. The sample may be heated from room temperature to a final temperature of 215° C. During this thermal treatment, the hydrous epsomite particles may be dehydrated. The sample may be held at 215° C. for 2 hours, then allowed to cool.
The resulting porous, dehydrated MgSO4 sample is shown in
If the dehydrated MgSO4 material is used in a high-temperature Replication Stage procedure, the MgSO4 may initially experience further thermal effects and sintering, and this may be considered as a part of the thermal treatment used to generate the template material. Such a procedure may be performed according to Scheme B in a tube furnace, as detailed in Section III. This portion of the thermal treatment may comprise heating a sample of the dehydrated MgSO4 material from room temperature to 580° C. under Ar gas flowing at 1102 sccm. During this thermal treatment, the MgSO4 may continue coarsening, and a portion may decompose to MgO.
The type of MgSO4 template material resulting from this process is identified herein as E1T1. Next, propylene (C3H6) gas may be introduced into the furnace, commencing surface replication. To the extent that the MgSO4 template material is still coarsening, the Template Stage and Replication Stage may overlap. At some point, the pyrolytic formation of the carbon perimorphic material over the E1T1 template material may stabilize the latter, preventing further coarsening and representing the true completion of the Template Stage. CVD may be continued for 2 hours, then the furnace may then be allowed to cool under sustained Ar flow.
After the furnace has cooled to room temperature, the PC material (E1T1P16) is collected. This PC material is shown in the SEM micrograph of
Example H4T1: In another exemplary Template Stage procedure, a Li-doped hydromagnesite precursor material may be thermally treated to form a porous MgO template material.
To demonstrate this, H4-type hydromagnesite particles may first be generated using the procedure described in Example H4. This material represents the template precursor material that might be generated in the Precursor Stage of a full implementation of the General Method.
Next, the template precursor material may be thermally treated. This may be performed in a tube furnace according to Scheme B under an Ar flow of 2000 sccm, as detailed in Section III. The sample may be heated from room temperature to a temperature of 1050° C. at a heating rate of 20° C./min. During this thermal treatment, CO2 gas may be released. In a full implementation of the General Method, the CO2 process gas released during the decomposition of the template precursor material may be conserved using conventional techniques. The sample may be held at 1050° C. for 20 minutes, after which the furnace may be allowed to cool.
The type of porous MgO template material resulting from this process is identified herein as H4T1. The template particles retain the precursor particles' plate-like superstructure, as shown in the SEM micrograph of
At 80 nm to 100 nm in diameter, the subunits of the template particles in H4T1 are considerably larger than the 50 nm to 60 nm subunits shown in the SEM micrograph of
Example H5T1: In another exemplary Template Stage procedure, a Li-doped hydromagnesite precursor material may be thermally treated to form a porous MgO template material.
To demonstrate this, H5-type hydromagnesite particles may first be generated using the procedure described in Example H5. This material represents the template precursor material that might be generated in the Precursor Stage of a full implementation of the General Method.
Next, the template precursor material may be thermally treated. This may be performed in a tube furnace according to Scheme B under an Ar flow of 2000 sccm, as detailed in Section III. The sample may be heated from room temperature to a temperature of 1050° C. at a heating rate of 20° C./min. During this thermal treatment, CO2 gas may be released. In a full implementation of the General Method, the CO2 process gas released during the decomposition of the template precursor material may be conserved using conventional techniques. The sample may be held at 1050° C. for 20 minutes, after which the furnace may be allowed to cool.
The type of porous MgO template material resulting from this process is identified herein as H5T1. The template particles retain the precursor particles' plate-like superstructure, but with much larger interstitial gaps between the subunits, as shown in the SEM micrograph of
Compared to H4T1-type template particles (
Example H6T1: In another exemplary Template Stage procedure, a hydromagnesite template template precursor material may be thermally treated to form a porous MgO template material.
To demonstrate this, a commercial hydromagnesite product (“Light Magnesium Carbonate” supplied by Akrochem Corporation) comprised predominately of plate-like particles may be employed. This commercial product was selected as it may provide similar chemical and morphological properties to that of a hydromagnesite template precursor; for this reason, this precursor material is described herein as H6. It represents the template precursor material that might be generated in the Precursor Stage of a full implementation of the General Method.
Next, the template precursor material may be thermally treated. This may be performed in a muffle furnace according to Scheme D, as detailed in Section III. The sample may be placed in a ceramic boat within the muffle furnace. The sample may be heated from room temperature to 750° C. at a heating rate of 5° C./min. During this thermal treatment, CO2 gas may be released. In a full implementation of the General Method, the CO2 process gas released during the decomposition of the template precursor material may be conserved using conventional techniques. The sample may then be held at 750° C. for 1 hour, then allowed to cool to room temperature.
The type of porous MgO template material resulting from this process is identified herein as H6T1. The template particles retain the precursor particles' plate-like superstructure, as shown in the SEM micrographs of
Examples M3T1, M4T1, M5T1: In another set of exemplary Template Stage procedures, magnesite template precursor materials may be thermally treated to form porous MgO template materials.
To demonstrate this, M3-type magnesite particles may first be generated using the procedure described in Example M3, M4, and M5. This material represents the template precursor material that might be generated in the Precursor Stage of a full implementation of the General Method.
Next, the template precursor material may be thermally treated. This may be performed in a muffle furnace according to Scheme D, as detailed in Section III. The sample (M3, M4 or M5) may be placed in a ceramic boat within the muffle furnace. The sample may be heated from room temperature to 580° C. at a heating rate of 5° C./min. The sample may then be maintained at to 580° C. for 1 hour. During this thermal treatment, CO2 gas may be released. In a full implementation of the General Method, the CO2 process gas released during the decomposition of the template precursor material may be conserved using conventional techniques. Next, the sample may be heated from 580° C. to 1050° C. at a heating rate of 5° C./min, with the sample maintained at this temperature for 3 hours. Then, it may be allowed to cool to room temperature.
The types of porous MgO template material resulting from this process are identified herein as M3T1, M4T1, and M5T1 (corresponding to variants based on M3, M4, and M5 template precursor materials). The M3T1, M4T1, and M5T1 template materials may be compared to demonstrate the use of dopants to increase coarsening effects during a thermal treatment. It is instructive to look at carbon perimorphic frameworks formed on these templates, since the frameworks in their native morphology are replicas of the templating surfaces (and negative replicas of the templating bulk). Additionally, carbon frameworks are also partially electron-transparent, allowing visualization of the templates' internal substructure.
The PC materials made using the M3T1, M4T1, and M5T1 template materials are identified herein as M3T1P2, M4T1P19 and M5T1P20, respectively (these exemplary Replication Stage procedures are described in Section VI). The endomorphic MgO in these PC materials may then be extracted with an aqueous H2CO3 extractant solution, leaving behind the carbon perimorphic products P1, P19 and P20. These perimorphic materials may be examined in order to determine the templates' substructure.
The P1, P19 and P20 perimorphic materials are shown in the SEM micrographs of
The PC materials (M5T1P20) made from M5T1 are shown in
Similar to the observations made for Li-doped magnesite template precursors, this shows that Na-doping may aid in coarsening the template and reducing the compactness of the perimorphic frameworks. Other dopants may have similar effects.
Examples M3T2, M4T2, and M5T2: In another set of exemplary Template Stage procedures, magnesite template precursor materials may be thermally treated to form porous MgO template materials.
To demonstrate this, M3-type, M4-type, and M5-type magnesite particles may first be generated using the procedure described in Examples M3, M4, and M5. This material represents the template precursor material that might be generated in the Precursor Stage of a full implementation of the General Method.
Next, the template precursor material may be thermally treated. This may be performed in a muffle furnace according to Scheme D, as detailed in Section III. The sample (M3, M4 or M5) may be placed in a ceramic boat within the muffle furnace. The sample may be heated from room temperature to 580° C. at a heating rate of 5° C./min. The sample may then be maintained at to 580° C. for 1 hour. During this thermal treatment, CO2 gas may be released. In a full implementation of the General Method, the CO2 process gas released during the decomposition of the template precursor material may be conserved using conventional techniques. Next, the sample may be heated from 580° C. to 900° C. at a heating rate of 5° C./min, with the sample maintained at this temperature for 1 hour. Then, it may be allowed to cool to room temperature.
The types of porous MgO template material resulting from this process are identified herein as M3T2, M4T2, and M5T2 (corresponding to variants based on M3, M4, and M5 template precursor materials).
Results for BJH are limited to a pore size range of 1.70 nm and 300 nm for this N2 gas adsorption method. Using the calculated BJH cumulative pore volume it may be possible to determine the porosity of the template particles. The porosity may be defined as the ratio of specific pore volume to the specific template volume and can be thought of as the percentage of total space occupied by pores with respect to the total particle. The BJH desorption cumulative pore volume (VPORE) may be used as a measure of the specific pore volume of the template particles. The specific MgO volume (VMgO) may be the specific volume of the MgO component of the porous MgO template—i.e. the reciprocal of the theoretical density of MgO. The specific template volume (VTEM) may be the sum of specific pore volume and specific MgO volume. Using the formula shown below, the porosity of the template particles may be determined:
After the 900° C. thermal treatment, the doped samples have 1.5% (M4T2) and 58% (M5T2) lower porosity than the undoped template material (M3T2). This demonstrates, as in previous exemplary procedures, that the level of dopant in the template material can be used to influence coarsening and densification effects. Taken in tandem with the results of Li-doping that have been described, this demonstrates the ability to tune a perimorphic framework's compactness, the size and morphology of its cellular subunits, and its ratio of endocellular vs. exocellular space.
Example M3T3: In another exemplary Template Stage procedure, magnesite template precursor materials may be thermally treated to form porous MgO template materials.
To demonstrate this, M3-type magnesite particles may first be generated using the procedure described in Example M3. This material represents the template precursor material that might be generated in the Precursor Stage of a full implementation of the General Method.
Next, the template precursor material may be thermally treated. This may be performed in a muffle furnace according to Scheme D, as detailed in Section III. The template precursor sample may be placed in a ceramic boat within the muffle furnace. The sample may be heated from room temperature to 580° C. at a heating rate of 5° C./min. The sample may then be maintained at to 580° C. for 13.5 hr. During this thermal treatment, CO2 gas may be released. In a full implementation of the General Method, the CO2 process gas released during the decomposition of the template precursor material may be conserved using conventional techniques. Next, the sample may be heated from 580° C. to 1050° C. at a heating rate of 5° C./min, with the sample maintained at this temperature for 1 hr. Then, it may be allowed to cool to room temperature.
The type of porous MgO template material resulting from this process are identified herein as M3T3.
Example N2T1: In another exemplary Template Stage procedure, a nesquehonite template precursor material may be thermally treated to form a porous MgO template material.
To demonstrate this, N2-type nesquehonite particles may first be generated using the procedure described in Example N2. This material represents the template precursor material that might be generated in the Precursor Stage of a full implementation of the General Method.
Next, the template precursor material may be thermally treated using steam as a coarsening aid. This may be performed in a rotary tube furnace according to Scheme A, as detailed in Section III. The quartz tube may be rotated at 1 rpm. Under dry Ar flow, an N2 sample may be heated from room temperature to 450° C. at a heating rate of 5° C./min in the furnace. Once the furnace reaches 450° C., Ar flow through a bubbler may be started at a flow rate of 2360 sccm. The chamber of the bubbler may be maintained at slight positive pressure of 0.23 psig and an external temperature of 100° C. may be maintained to saturate the bubbler headspace with water vapor. The furnace may be maintained at 450° C. for 1 hour, after which it may then be heated at a heating rate of 5° C./min to 500° C. After 1 hour at 500° C., the furnace may be heated at a heating rate of 5° C./min to the final temperature of 1000° C. and held at 1000° C. for 1 hour. During this thermal treatment, CO2 gas may be released. In a full implementation of the General Method, the CO2 process gas released during the decomposition of the template precursor material may be conserved using conventional techniques. At this point, dry Ar flow may be resumed and the sample may be cooled to room temperature under flowing, dry Ar.
The type of porous MgO template material resulting from this process is identified herein as N2T1. The template particles retain the precursor particles' elongated superstructure. This can be observed in SEM micrographs of an exemplary PC material (N2T1P21) made via surface replication on the N2T1 template particles. The N2T1P21 PC material, comprising a thin, electron-transparent carbon perimorphic phase and an N2T1 endomorphic phase, is shown in the SEM micrographs of
In
Example N2T2: In another exemplary Template Stage procedure, a nesquehonite template precursor material may be thermally treated to form a porous MgO template material.
To demonstrate this, N2-type nesquehonite particles may first be generated using the procedure described in Example N2. This material represents the template precursor material that might be generated in the Precursor Stage of a full implementation of the General Method.
Next, the template precursor material may be thermally treated. This may be performed in a rotary tube furnace according to Scheme A, as detailed in Section III. The quartz tube may be rotated at 1 rpm. Under dry Ar flow, an N2-type sample may be heated from room temperature to 450° C. at a heating rate of 5° C./min in the furnace. Once the furnace reaches 450° C., dry Ar flow through a bubbler may be started at a flow rate of 2360 sccm. The furnace may be maintained at 450° C. for 1 hour, after which it may then be heated at a heating rate of 5° C./min to 500° C. After 1 hour at 500° C., the furnace may be heated at a heating rate of 5° C./min to the final temperature of 1000° C. and held at 1000° C. for 1 hour. During this thermal treatment, CO2 gas may be released. In a full implementation of the General Method, the CO2 process gas released during the decomposition of the template precursor material may be conserved using conventional techniques. At this point, dry Ar flow may be resumed and the sample may be cooled to room temperature under flowing, dry Ar.
The type of porous MgO template material resulting from this process is identified herein as N2T2. N2 gas adsorption may be performed on these templates, applying methods described previously. As seen in
Examples N2T3, N2T4, N2T5, and N2T6: In another set of exemplary Template Stage procedures, nesquehonite template precursor materials may be thermally treated to form porous MgO template materials.
To demonstrate this, N2-type nesquehonite particles may first be generated using the procedure described in Example N2. This material represents the template precursor material that might be generated in the Precursor Stage of a full implementation of the General Method.
Next, the template precursor material may be thermally treated in several ways summarized in
The type of porous MgO template materials resulting from the processes are identified herein as N2T3, N2T4, N2T5, and N2T6. These variants were performed to test how thermal treatment parameters affect the resulting template morphology.
During thermal treatment of hydrated MgCO3·xH2O template precursor materials, H2O and CO2 are the two primary gases released. The thermogravimetric mass loss profiles for N2-type template precursor material in Ar is shown in
The N2T3-type template material is shown in the SEM micrographs of
The N2T4-type template material is shown in the SEM micrographs of
The N2T5-type template material is shown in the SEM micrographs of
The template particles' internal macropores are inherited by the PC particles produced in the Replication Stage and the perimorphic frameworks produced in the Separation Stage. These internal macropores can be clearly observed in Cui's mesoporous graphene fibers. Elimination of these macropores in the template material results in their absence in the perimorphic material, as shown in the SEM micrographs of
A PC material (N2T6P22) made on N2T6-type template material is shown in the SEM micrographs of
Example L2T1: In another exemplary Template Stage procedure, a lansfordite template precursor material may be thermally treated to form a porous MgO template material.
To demonstrate this, L2-type lansfordite particles may first be generated using the procedure described in Example L2. This material represents the template precursor material that might be generated in the Precursor Stage of a full implementation of the General Method.
Next, the template precursor material may be thermally treated. This may be performed in a tube furnace according to Scheme B, as described in Section III. An L2-type sample may be placed in the tube furnace. Under an Ar flow of 1220 sccm, the furnace may be heated from room temperature to 640° C. at a heating rate of 20° C./min and maintained at 640° C. for 2 hours. During this thermal treatment, CO2 gas may be released. In a full implementation of the General Method, the CO2 process gas released during the decomposition of the template precursor material may be conserved using conventional techniques. The furnace may then be allowed to cool to room temperature under sustained Ar flow.
The type of porous MgO template material resulting from this process is identified herein as L2T1. The morphology of L2T1-type template particles can be discerned from the native morphology of carbon perimorphic frameworks synthesized on them. Such frameworks are shown in the SEM micrographs of
Example L3T1: In another exemplary Template Stage procedure, a partially dehydrated lansfordite template precursor material may be thermally treated to form a porous MgO template material.
To demonstrate this, L3-type partially dehydrated lansfordite particles may first be generated using the procedure described in Example L3. This material represents the template precursor material that might be generated in the Precursor Stage of a full implementation of the General Method.
Next, the template precursor material may be thermally treated. This may be performed in a tube furnace according to Scheme B, as described in Section III. The L3-type template precursor material may be placed in the tube furnace. While under an Ar flow of 1220 sccm, the furnace may be heated from room temperature to 640° C. at 20° C./min and maintained at 640° C. for 2 hours. During this thermal treatment, CO2 gas may be released. In a full implementation of the General Method, the CO2 process gas released during the decomposition of the template precursor material may be conserved using conventional techniques. After this thermal treatment, the furnace may be allowed to cool to room temperature under sustained Ar flow. The furnace may then be allowed to cool to room temperature under sustained Ar flow.
The type of porous MgO template material resulting from this process is identified herein as L3T1. The morphology of the L3T1-type template particles can be discerned from the native morphology of carbon perimorphic frameworks synthesized on them. In
Example L3T2: In another exemplary Template Stage procedure, a partially dehydrated lansfordite template precursor material may be thermally treated to form a porous MgO template material.
To demonstrate this, L3-type partially dehydrated lansfordite particles may first be generated using the procedure described in Example L3. This material represents the template precursor material that might be generated in the Precursor Stage of a full implementation of the General Method.
Next, the template precursor material may be thermally treated. This may be performed in a tube furnace according to Scheme B, as described in Section III, with a few modifications. While under a CO2 flow of 815 sccm, the furnace may be heated to 540° C. and maintained at that temperature. An L3-type sample may be staged inside the quartz tube but outside of the heating zone prior to the thermal treatment. Then, the template precursor material may be rapidly introduced into the preheated zone by a pushing mechanism and maintained at 540° C. for 30 minutes. During this thermal treatment, CO2 gas may be released. In a full implementation of the General Method, the CO2 process gas released during the decomposition of the template precursor material may be conserved using conventional techniques. Finally, the treated template may be removed from the heat zone and allowed to cool to room temperature under sustained CO2 flow.
The type of porous MgO template material resulting from this process is identified herein as L3T2. A C@MgO PC material made by forming a thin carbon perimorph on the L3T2-type template material is shown in the SEM micrographs of
Example A1T1: In another exemplary Template Stage procedure, a spray-dried MgCO3·xH2O template precursor material comprising hollow, spherical particles may be thermally treated to form a porous MgO template material.
To demonstrate this, A1-type spray-dried MgCO3·xH2O particles may first be generated using the procedure described in Example A1. This material represents the template precursor material that might be generated in the Precursor Stage of a full implementation of the General Method.
Next, the template precursor material may be thermally treated. This may be performed in a tube furnace according to Scheme A with a tube rotation speed of 1 RPM, as detailed in Section III. The sample may be placed in the tube furnace. While under an Ar flow of 1271 sccm, the furnace may be heated from room temperature to 100° C. at a heating rate of 20° C./min and maintained at 100° C. for 1 hour. The furnace may then be heated to 500° C. at a heating rate of 20° C./min and maintained at 500° C. for 1 hour. Finally, the furnace may be heated to 640° C. at a heating rate of 20° C./min and maintained at 640° C. for 3 hours. During this thermal treatment, CO2 gas may be released. In a full implementation of the General Method, the CO2 process gas released during the decomposition of the template precursor material may be conserved using conventional techniques. The furnace may then be allowed to cool to room temperature under sustained Ar flow.
The type of porous MgO template material resulting from this process is identified herein as A1T1. The template particles retain the precursor particles' hollow, hierarchical-equiaxed superstructure, with some particles comprising fragments of the shells, as shown in the SEM micrographs of
Example A3T1: In another exemplary Template Stage procedure, a spray-dried MgCO3·xH2O template precursor material comprising hollow, spherical particles may be thermally treated to form a porous MgO template material.
To demonstrate this, A3-type spray-dried MgCO3·xH2O particles may first be generated using the procedure described in Example A3. This material represents the template precursor material that might be generated in the Precursor Stage of a full implementation of the General Method.
Next, the template precursor material may be thermally treated. This may be performed in a tube furnace according to Scheme A, as detailed in Section III. The sample may be placed in a ceramic boat in the tube furnace. While under an N2 flow of 2408 sccm, the furnace may be heated from room temperature to 200° C. at a heating rate of 20° C./min and maintained at 200° C. for 1 minute. The furnace may then be heated to 500° C. at a heating rate of 5° C./min and maintained at 500° C. for 1 minute. Finally, the furnace may be heated to 900° C. at a heating rate of 20° C./min and maintained at 900° C. for 15 minutes. During this thermal treatment, CO2 gas may be released. In a full implementation of the General Method, the CO2 process gas released during the decomposition of the template precursor material may be conserved using conventional techniques. The furnace may then be allowed to cool to room temperature under sustained N2 flow.
The type of porous MgO template material resulting from this process is identified herein as A3T1. The template particles retain the precursor particles' hollow, hierarchical-equiaxed superstructure, as shown in
Example C1T1: In another exemplary Template Stage procedure, a spray-dried template precursor material comprising hollow, hierarchical-equiaxed particles may be thermally treated to form a porous MgO template material.
To demonstrate this, C1-type spray-dried particles may first be generated using the procedure described in Example A1. This material represents the template precursor material that might be generated in the Precursor Stage of a full implementation of the General Method.
Next, the template precursor material may be thermally treated. This may be performed in a muffle furnace according to Scheme D, as detailed in Section III. The C1-type template precursor material may be placed in a ceramic boat within the muffle furnace. The sample may then be heated from room temperature to 650° C. at a heating rate of 5° C./min. The sample may be maintained at 650° C. for 3 hours. During this thermal treatment, CO2 gas may be released. In a full implementation of the General Method, the CO2 process gas released during the decomposition of the template precursor material may be conserved using conventional techniques. Then, the furnace may then be allowed to cool to room temperature.
The type of porous MgO template material resulting from this process is identified herein as C1T1.
Example Ca1T1: In another exemplary Template Stage procedure, a precipitated CaCO3 template precursor material (Albafil), herein described as Ca1, may be thermally treated to form a porous MgO template material.
The precipitated Car-type particles represent the template precursor material that might be generated in the Precursor Stage of a full implementation of the General Method.
Next, the template precursor material may be thermally treated. This may be performed in a tube furnace according to Scheme C, as detailed in Section III. The Car-type sample may be placed in a ceramic boat within the tube furnace. The furnace may be heated to 1050° C. under flowing Ar at 1102 sccm. During this thermal treatment, CO2 gas may be released. In a full implementation of the General Method, the CO2 process gas released during the decomposition of the template precursor material may be conserved using conventional techniques. When the furnace reaches 1050° C., methane (CH4) gas may be introduced into the system to begin forming a carbon perimorph on the templating surfaces. While this surface replication step may be thought of as part of the Replication Stage, the template material may continue coarsening concurrently until stabilized by the carbon perimorph. The system may be maintained at 1050° C. for 15 minutes under flowing CH4 and Ar, then CH4 flow may be discontinued and the furnace may be allowed to cool to room temperature under sustained Ar flow.
The type of calcium oxide (CaO) template material resulting from this process is identified herein as Ca1T1, and the PC material made using the Ca1T1 template material is identified herein as Ca1T1P17. It is instructive to look at the P17-type carbon perimorphic material after extraction of the endomorphic Ca1T1 template material, since the frameworks in their native morphology are replicas of the templating surfaces (and negative replicas of the templating bulk). Additionally, carbon frameworks are also partially electron-transparent, allowing visualization of the templates' internal substructure.
Example Li1T1: In another exemplary Template Stage procedure, spray dried lithium carbonate template precursor material comprising hollow, hierarchical-equiaxed particles may be thermally treated to form a porous Li2CO3 template material.
To demonstrate this, Li1-type spray-dried particles may first be generated using the procedure described in Example Li1. This material represents the template precursor material that might be generated in the Precursor Stage of a full implementation of the General Method.
Next, the template precursor material may be thermally treated. This may be performed in a tube furnace according to Scheme C, as detailed in Section III. The Li1-type sample may be placed in a ceramic boat within the tube furnace. The furnace may be heated to 580° C. under flowing Ar at 1271 sccm. At this point, C3H6 gas may be introduced into the system to begin forming a carbon perimorph on the templating surfaces. While this surface replication step may be thought of as part of the Replication Stage, the template material may continue coarsening concurrently until stabilized by the carbon perimorph. The system may be maintained at 580° C. for 870 minutes under flowing C3H6 and Ar, then C3H6 flow may be discontinued and the furnace may be allowed to cool to room temperature under sustained Ar flow.
The type of Li2CO3 template material resulting from this process is identified herein as Li1T1, and the PC material made using the Li1T1 template material is identified herein as Li1T1P18. It is instructive to look at the P18-type carbon perimorphic material after extraction of the endomorphic Li1T1 template material, since the frameworks in their native morphology are replicas of the templating surfaces (and negative replicas of the templating bulk). Additionally, carbon frameworks are also partially electron-transparent, allowing visualization of the templates' internal substructure.
In order to demonstrate the general applicability of the Replication Stage to a variety of template materials, a number of exemplary Replication Stage procedures are presented below. For purposes of demonstration, each exemplary procedure comprises CVD growth of carbon perimorphs on selected template particles. However, it should be noted that other procedures, and perimorphs of alternative compositions, will be obvious to those knowledgeable in the art.
In some exemplary Replication Stage procedures, template materials may been formed from template precursor materials in a separate and distinct Template Stage that occurred in a different reactor. In other instances, the Template Stage procedure and the Replication Stage procedure may both be performed in the same reactor.
Some of the exemplary Replication Stage procedures presented in this section utilize previously named and described template materials. Additionally, new template materials are described in some of the exemplary Template Stage procedures, and for this reason, we describe the synthesis of these new templates in this section.
In some procedures, the Replication Stage immediately follows the Template Stage. In this scenario the Rn is described as ‘N/A’ in
The Separation Stage comprises endomorphic extraction and perimorphic separation. In some variants, this may occur in an integrated, one-pot technique. In other variants, the Separation Stage may occur in two or more separate and distinct stages. For example, endomorphic extraction may involve mixtureing the PC material in the conserved process liquid and dissolving the endomorphic material inside the perimorphic material. Then, the perimorphic material may be separated from the stock solution. The stock solution may then be precipitated at atmospheric pressure. The precipitate may then be slurried into the process water at a higher solids concentration. By modulating temperature or pressure, the solids in this concentrated mixture may then be re-dissolved at higher concentrations to create a concentrated stock solution that may be utilized in Precursor Stage.
Example VIIa: In an exemplary endomorphic extraction procedure, MgO endomorphs may be extracted from carbon perimorphs may be obtained by dissolving the MgO in an extractant solution comprising aqueous H2CO3.
First, about 12.5 grams of a C@MgO PC powder comprising approximately 94.75% MgO endomorph and 5.25% carbon perimorph (grown via CVD) by weight may be slurried into 2.5 L of deionized water in a 3 L round-bottom flask. This water represents a conserved process water that would be derived from the Precursor Stage in a full implementation of the General Method. A gas line fitted with a 0.5 μm diffusion stone (to decrease CO2 bubble size and increase reaction efficiency) may be fed into the bottom of the flask, and the water may be stirred with a magnetic stir plate. CO2 gas may be continuously bubbled into the tank over 141 minutes at a flow rate of 4 scfhair. This CO2 represents a conserved process gas that would be derived from the Precursor Stage or Template Stage in a full implementation of the General Method. Dissolution of the CO2 and reaction with the process water generates an aqueous H2CO3 extractant solution. The reaction of the aqueous H2CO3 extractant solution with the endomorphic MgO results in endomorphic extraction and the generation of a new aqueous Mg(HCO3)2 stock solution outside of the carbon perimorphic frameworks.
Perimorphic separation of the carbon perimorphic frameworks from the aqueous Mg(HCO3)2 stock solution may be obtained by filtering the mixture. The carbon perimorphic frameworks may be rinsed and dried, and an ash test may be performed. The carbon perimorphic frameworks may contain approximately 9.49% MgO, representing a 99.5% removal efficiency of the MgO template material. The remainder of unextracted MgO may be hermetically encapsulated within certain carbon frameworks. Higher extraction efficiencies may be gained with higher-energy agitation techniques, which may facilitate the breaching of sealed perimorphic walls.
In a full implementation of the General Method, the separated aqueous Mg(HCO3)2 stock solution may then be conserved to be utilized in the Precursor Stage.
Example VIIb: In another exemplary endomorphic extraction procedure, MgO endomorphs may be extracted from carbon perimorphs via a shuttling technique.
First, 500 mL of water may be magnetically stirred in a 1 L glass beaker at 700 RPM. This water represents a conserved process water that would be derived from the Precursor Stage in a full implementation of the General Method. Next, CO2 process gas may be continuously bubbled at 3-5 scfhair through the process water from a dip tube, forming an aqueous H2CO3 extractant solution. This CO2 represents a conserved process gas that would be derived from the Precursor Stage or Template Stage in a full implementation of the General Method. Approximately 10 g of a C@MgO PC material (yield 3.5%) comprising elongated particles may gradually be introduced into the extractant solution. Upon complete integration of the C@MgO PC material into the solution, the mixture may appear black and possess a pH of 9. The beaker may be covered to maintain a CO2-rich atmosphere.
After 24 hours of reaction, the conductivity of the mixture may be 19.7 mS/cm measured at 19.6° C., have a pH of 8, and the mixture may appear gray. This mixture comprises the perimorphic product and a new aqueous Mg(HCO3)2 stock solution that might be used in the Precursor Stage in a full implementation of the General Method. The solids may then be separated from the stock solution using conventional techniques.
The solids from this mixture may be seen in the optical micrograph
In
The mechanism for this may be the preferential adsorption and nucleation of CO2 nanobubbles in the hydrophobic carbon framework, increasing the internal CO2 pressure within the framework and therefore the solubility of Mg(HCO3)2 within the framework. This creates a concentration gradient that drives the solvated ions into the surrounding process water, where they precipitate due to the lower external CO2 pressure. Hence, shuttling reduces the volume of process water needed for endomorphic extraction, as well as the required vessel size.
Example VIIc: Endomorphic extraction of certain metal oxide or metal carbonate compounds may be facilitated by rendering the CO2 supercritical. In an exemplary procedure, 3.007 g of MgO (Elastomag 170 calcined at 1050° C. for 1 hour) may be slurried with 100.00 g DI water, resulting in a solution conductivity of 340 S/cm at 12.4° C. This translates to a mixture concentration of 30 g/L MgO. The mixture may be poured into a 1 L pressure vessel with magnetic stirring and a heating mantle. Approximately 600 g of dry ice (solid CO2) may be added to the reactor, and the reactor then sealed. After 101 minutes of heating, the minimum conditions for supercritical CO2 conditions may be surpassed at 31.4° C. and 1,125 psi. After a total 144 minutes, the reactor conditions may reach 36.2° C. and 1200 psi. The reactor may then be actively chilled with a cooling coil for 74 minutes, after which its conditions may reach 18.3° C. and 675 psi. The pressure in the reactor may then be slowly released, and after 6 minutes the reactor may have equilibrated to atmospheric pressure with a temperature probe reading of −5.0° C., due to the pressure release. Approximately 23 minutes after the pressure release, a sample may be taken from the solution, which may have a conductivity of 30.2 mS/cm at 4.5° C. The solution may be clear, with no signs of particles or precipitation. This higher concentration solution may then be utilized for crystallization of the template precursor material.
Example VIId: In another exemplary Separation Stage procedure, endomorphic extraction of a water-soluble endomorphic template material may be obtained via simple dissolution in water. This may be demonstrated by mixtureing a C@MgSO4 PC material, as shown in the SEM micrograph of
Perimorphic Separations
Perimorphic products may be separated using a number of conventional techniques. In one technique, a liquid-liquid separation may be utilized. This may be demonstrated by taking the mixture produced by the shuttling process described above and blending it with an immiscible solvent, like hexane. The carbon perimorphic frameworks migrate into the solvent phase, while the nesquehonite remains in the aqueous phase. This results in phase separation and two distinct slurries, as shown in
Separation of the carbon perimorphic frameworks may also be obtained simply using flotation. In some carbon perimorphic frameworks, air bubbles may remain trapped in the exocellular pores during the liquid-phase endomorphic extraction. This may render the frameworks buoyant or quasi-buoyant upon endomorphic extraction. Furthermore, subjecting a mixture of these bubble-infused frameworks to a partial vacuum increases their buoyancy, as internal bubbles expand and extrude water from the porous framework. The progressive flotation and separation of carbon perimorphic frameworks under partial vacuum is shown in
A number of variants and improvements of these separation techniques may be readily envisioned. Flotation may be improved with the use of a solvent, as would be typical in a conventional froth flotation process. Frameworks made on template materials with greater particle porosity may retain more air and be more buoyant. Hollow spheres, in particular, may contain more trapped air and be more buoyant.
Concentrating Stock Solutions
In some cases, it may be desirable to create a concentrated stock solution after separating the perimorphic product. A mixture of precipitated particles, such as the nesquehonite precipitated in the shuttling procedure described above, may be re-dissolved under conditions that allow higher solution concentrations. For example, an aqueous mixture of precipitated MgCO3·xH2O particles may be subjected to higher CO2 pressure in order to make a concentrated stock solution, as illustrated in
Example VIIe: In one exemplary procedure, endomorphic MgO may be dissolved at higher concentrations under pressure. To demonstrate this, 15 g of MgO (Elastomag 170) template may be slurried with 750 g of deionized water, which may represent a conserved process water retained from the Precursor Stage. The water may be chilled to 5° C. The solids concentration of the mixture may be 20 g/L MgO, or approximately double the maximum concentration of MgO that can be dissolved into an aqueous H2CO3 extractant solution at atmospheric pressure. The mixture may have a solution pH of approximately 10.5 and a resulting solution conductivity of 146 S/cm. The mixture may be poured into a 1 L pressure vessel with magnetic stirring, a high-pressure gas inlet, and a purging needle valve. The reactor may be sealed and purged through a purging needle valve by opening the high-pressure gas inlet, allowing pressurized CO2 gas, representing conserved CO2 process gas recaptured in the Precursor Stage and Template Stage, to flow into the vessel for 2 minutes to displace any air. The purge valve may then be closed, and the reactor pressurized with CO2 to 125 psi. After 65 minutes, the conductivity may be approximately 15.6 mS/cm measured at 16.3° C. and a pH of 8.5. This conductivity represents an Mg(HCO3)2 solution concentration equivalent to 10 g/L of dissolved MgO, which at atmospheric pressure may require an order of magnitude longer reaction time to achieve. At the 290 minute mark, the conductivity may be approximately 27.8 mS/cm measured at 19.5° C. and a pH of 7.5. The conductivity value and pH measurements at 290 minutes may signify Mg(HCO3)2 solubilities greater than the approximately 10 g/L MgO possible at atmospheric pressure.
Increased CO2 pressure may likewise be used to create concentrated stock solutions from MgCO3·xH2O solutes, such as those produced via shuttling procedures. These concentrated stock solutions may be produced via multistep Separation Stage procedures, in which stock solutions are used to precipitate solids that are re-dissolved under conditions allowing higher solubility. Alternatively, endomorphic extractions utilizing aqueous H2CO3 extractant solutions might be performed under increased CO2 pressure, such that higher concentrations are obtained without precipitation and re-dissolution under increased CO2 pressure.
In the Preferred Method, carbon perimorphic frameworks are synthesized using MgO templates derived from MgCO3·xH2O precursors. While coarsening may reduce the fine structuring of these MgO templates, a typical MgO template comprises a porous substructure of conjoined nanocrystals. This creates a labyrinthine framework with endocellular and exocellular labyrinths. This labyrinthine structure is not specific to carbon frameworks formed on these templates-any framework will have the same native morphology. However, carbon frameworks with thin, conformal perimorphic walls can be utilized to study these architectures due to their ability to create fine, electron-translucent replicas of the template.
As an example,
While the subunits are uniformly equiaxed, the superstructures of frameworks derived from porous MgO templates have diverse geometries to the variety of precursors from which MgO templates can be derived. For example, frameworks generated on MgO templates made from nesquehonite template precursor impart elongated, fibroidal superstructures, as shown by the labyrinthine framework in
In addition to these diverse architectures, fragmentation and deformation of the frameworks may result from mechanical agitation, such as the multilayer stack of thin pseudomorphs shown in
If the template precursor, or some decomposition product of a template precursor, is coarsened during the Template Stage, the resulting perimorphic frameworks become less compact. In one experiment, MgO templates were sintered for 2 hours at 1,000° C. prior to a Replication Stage. The resulting MgO templates were quasi-polyhedral and generally larger than 100 nm in diameter.
While these less compact, less regular frameworks are not as ordered as more compact frameworks with more regular geometries, they may still be assembled into multicellular clusters with attractive functional properties like those described in U.S. Patent Application 62/448,129. One benefit of less compact frameworks with flexible perimorphic walls is their increased pseudoelasticity—i.e. a collapsed framework that is natively coarse, despite being densified by its collapse, may retain the ability to expand back to its native dimensions without covalent failure.
Elasticity is shown in the sequences of optical images in
Raman spectroscopy is commonly used to characterize carbons and is a critical tool used to characterize the lattice structure of the exemplary carbon perimorphic materials in this disclosure. The details regarding the equipment and techniques used for Raman analysis are detailed in Section III.
Three main spectral features are typically associated with sp2-bonded carbon: the “G band” (typically at or around 1585 cm−1), the “2D band” (typically between 2500 and 2800 cm−1), and the “D band” (typically between 1200 and 1400 cm−1). The G band is associated with sp2-hybridized carbon. The D band is associated with radial breathing mode phonons in polycyclic sp2-hybridized carbon and is activated by defects. Therefore, the D band is associated with disorder and the peak intensity ratio of the D and G bands iprovides a measure of disorder. Another feature associated with disorder is an interband region located between the D and G bands. The presence of broad peaks within this interband range increases the height of the trough between the D and G bands, and this height may therefore be used as a measure of disorder, where higher troughs are associated with greater disorder. For this reason, the present disclosure utilizes the height of this trough to characterize disorder. The trough height is defined herein as the local minimum intensity value occurring between the wavenumber associated with the D peak and the wavenumber associated with the G peak. The intensity value at this wavenumber is then compared to the G peak intensity to characterize disorder.
In order to avoid resorting to subjective profile-fitting judgments, the present disclosure analyzes the unfitted Raman spectra of the carbon perimorphic materials presented herein. All references to peak positions and peak intensities therefore relate to unfitted peak positions and are derived without profile fitting. Additionally, all peak positions and peak intensities reported are measured under 532 nm excitation. The intensities of the G, 2D, D, and trough are designated herein as IG, I2D, ID, and ITr, respectively.
The ID/IG peak intensity ratios for the carbon perimorphs in the PC materials range between 0.78-1.27, indicating that these samples comprise disordered carbons. This disorder is corroborated by the generally high ITr/IG peak intensity ratios, which range between 0.17-0.64 as shown in
For crystalline sp2-hybridized carbons like graphite, the G band is expected to be centered around ˜1580 cm−1. It has also been shown that the G band can be red-shifted for carbons under compressive strain and blue-shifted under tensile strain. For sp2 carbons, the D band, if present, should be centered around ˜1350 cm−1 (for 532 nm laser). Red-shifting of the D band position, as seen in some of the samples, is indicative of sp3 defect states present within the disordered sp2 carbons.
For the samples described herein, the G band peak positions range between 1581-1609 cm−1, and the D band peak positions range between 1324-1358 cm−1, as shown in
This section is organized according to the following outline:
The term “graphenic,” as used herein, describes a two-dimensional, polycyclic structure of sp2-hybridized or sp3-hybridized atoms. While graphene denotes a form of carbon, we utilize the term “graphenic” herein to describe a variety of graphene polymorphs (including known or theorized polymorphs such as graphene, amorphous graphene, phagraphene, haeckelites, etc.), as well as to describe other two-dimensional graphene analogues (e.g. atomic monolayers of BN, BCxN, etc.) Hence, the term “graphenic” is intended to encompass any hypothetical polymorph meeting the basic criteria of two-dimensionality, polycyclic organization and sp2 or sp3 hybridization.
“Two-dimensional” herein describes a molecular-scale structure comprising a single layer of atoms. A two-dimensional structure may be embedded or immersed in a higher-dimensional space to form a larger-scale structure that, at this larger scale, might be described as a three-dimensional. For instance, a graphenic lattice of subnanoscopic thickness might curve through three-dimensional space to form the atomically thin wall of a nanoscopically three-dimensional cell. This cell would still be described two-dimensional at the molecular scale.
A “ring” is defined herein as a covalent chain of atoms that together comprise a closed, polyatomic polygon of fewer than 10 atomic vertices. Each of the cyclic structures in a polycyclic arrangement comprise a ring. Each of the atoms comprising a given ring may be described as an atomic member belonging to that ring, and the ring may be described accordingly (i.e. a “6-member” ring describes a hexagonal ring formed by 6 atomic members).
An “sp2 ring” is herein defined as a ring comprising all sp2-hybridized atomic members.
An “spx ring” is herein defined as a ring comprising atomic members that do not all share the same orbital hybridization.
A “chiral ring” is defined herein as an spx ring in which the covalent chain of atomic members comprises one or more chiral segments, wherein the two atomic termini of these chiral segments are sp3-hybridized atoms connected to each other via sp3-sp3 bonds. Chiral rings occur at tectonic zone transitions.
A “chiral column” is defined herein as a series of z-adjacent chiral rings connected to one another via one or more z-directional chains of sp3-sp3 bonds. A chiral column tends to form over a base-layer chiral ring and represents the lateral terminus of a diamondlike seam. A chiral column may contain one or more spx helices.
An “spx helix” is defined herein as a type of helical, one-dimensional chain constructed from both sp2-hybridized and sp3-hybridized atomic members. The axis of an spx helix is z-oriented.
An “spx double helix” is defined herein as the structure formed by two spx helices sharing the same chirality and the same axis.
An “sp2 helix” is defined herein as a type of helical, one-dimensional chain constructed from only sp2-hybridized atomic members. The axis of an spx helix is z-oriented.
An “sp2 double helix” is defined herein as the structure formed by two sp2 helices sharing the same chirality and the same axis.
“Adjacent rings” herein describes two rings that have at least two common atomic members, and thus share at least one common side. In organic chemistry these rings might comprise fused or bridged rings, but not spirocyclic rings. Two adjacent rings may be described as “ring-adjacent.”
“Ring-connected” herein describes a structure that is connected via a “ring pathway,” or path of adjacent rings. We may speak of ring-connectedness according to two usages. In the first usage, we may say that one part of a structure is ring-connected to some other part of the structure. This means that there is a ring pathway that connects the two referenced parts. For example, a ring R1 within a graphenic structure is ring-connected to another ring R2 within the structure if there exists a path of adjacent rings starting at R1 and ending at R2. In the second usage, we may say that a referenced structure is itself ring-connected. This means that any part of the referenced structure can be reached from any other part via at least one ring pathway. We may also describe structures that are not ring-connected as ring-disconnected.
A “ring pathway” herein describes a pathway of adjacent rings that connects two referenced structures.
A “ring connection” herein describes a single ring that ring-connects two referenced structures.
“Sp2 ring-connected” herein describes a structure that is connected via an “sp2 ring pathway,” or pathway of adjacent sp2 rings. Like ring-connectedness, we may speak of sp2 ring-connectedness according to two usages. In the first usage, we may say that one part of a structure is sp2 ring-connected to some other part of the structure. This means that there is an sp2 ring pathway that connects the two referenced parts. In the second usage, we may say that a referenced structure is itself sp2 ring-connected. This means that any part of the referenced structure can be reached from any other part via at least one sp2 ring pathway. Since sp2 ring-connectedness is a specific case of ring-connectedness, it implies ring-connectedness, while ring-connectedness does not imply sp2 ring-connectedness. In certain cases we may describe certain ring-connected structures as “sp2 ring-disconnected,” meaning that while they are ring-connected, they are not ring-connected by an sp2 ring pathway.
An “edge atom” is defined as an atom that (i) belongs to a ring, and (ii) is not surrounded on all sides by rings. An edge atom always has multiple nearest neighbors that are also edge atoms, forming a chain.
An “edge” is defined as a chain of edge atoms. Starting from any given edge atom, it is possible to trace from this first atom a chain of nearest-neighbor edge atoms, wherein any given pair of nearest-neighbor edge atoms within the chain are co-members of exactly one ring. Some edges may form a closed circuit, where the first atom and last atom traced are nearest neighbors to each other.
An “edge segment” is defined as a chain of nearest-neighbor edge atoms contained within a larger edge.
An “interior atom” is defined herein as an atom that (i) belongs to a ring, and (ii) is surrounded on all sides by rings.
A “graphenic structure” is defined herein as a polycyclic, ring-connected group of two or more rings. Every ring in a graphenic structure is ring-connected to every other ring, although not necessarily sp2 ring-connected. Each atom belonging to a graphenic structure may be classified as either an interior atom or an edge atom.
A “graphenic region” or “region” is herein defined as a subsidiary portion of some larger graphenic structure that itself fulfills all the requirements of a graphenic structure.
“Ring disorder” is herein defined as the presence of non-hexagonal rings in a graphenic structure.
Ring-disordered graphenic structures include amorphous, haeckelite, pentagonal, or other molecular tilings. The presence of non-hexagonal rings creates regions of nonzero Gaussian curvature in ring-disordered graphenic structures. If inserted into a hexagonally tiled lattice, a 5-member ring incudes positive Gaussian curvature, while a 7-member ring induces negative Gaussian curvature. For example, a fullerene comprises a curved graphenic structure formed by 20 hexagons and 12 pentagons.
“Ring order” is herein defined as a substantially hexagonal molecular tiling. Ring-ordered graphenic structures may be flexed or wrinkled due to their low bending stiffness.
A “system” is herein defined as some polyatomic physical structure comprising a group of atoms cohered via either chemical bonds or van der Waals interactions. A system may contain any number of graphenic structures, including none. It is a general term for describing some physical structure under consideration.
A “graphenic system” is herein defined as a system consisting of one or more distinct graphenic structures. A graphenic structure belonging to a graphenic system may be described as a “graphenic member” or “member” of the graphenic system. A graphenic system does not include any elements other than its graphenic members.
A “graphenic singleton” or “singleton” is herein defined as a graphenic system comprising a single, distinct graphenic structure.
A “graphenic assembly” or “assembly” is herein defined as a graphenic system comprising two or more distinct graphenic structures.
A “van der Waals assembly,” or “vdW assembly,” is herein defined as a multilayer graphenic assembly in which the graphenic structures are cohered principally or substantially by intermolecular forces. The graphenic structures in a vdW assembly may also be cohered via other mechanisms.
A “double screw dislocation” is herein defined as a dislocation formed by two screw dislocations sharing the same chirality and the same dislocation line. A double screw dislocation in a graphenic system forms a graphenic double helicoid. The braid-like geometry of double helicoids may physically interlock the two helicoids.
A “multilayer” graphenic system is herein defined as a graphenic system comprising more than one layer in vdW contact, on average. A multilayer graphenic system may possess monolayer regions.
Analytically, we may define a multilayer graphenic system as one possessing an average BET surface area no more than 2,300 m2/g, as measured by N2 adsorption.
A “Y-dislocation” is herein defined as a ring-connected, Y-shaped graphenic region formed by a layer's bifurcation into a laterally adjacent bilayer. The two “branches” of the Y-shaped region comprise z-adjacent spx rings, which together comprise a diamondlike seam situated at the interface between the laterally adjacent layer and bilayer. The characteristic Y-shaped geometry is associated with a cross-sectional plane of the layers and the diamondlike seam.
A “diamondlike seam” or “seam” is herein defined as a two-dimensional sheet of z-adjacent spx rings forming a z-oriented interface between xy-oriented layers to either side. A cubic diamondlike seam comprises chair conformations, while a hexagonal diamondlike seam comprises chair, boat, and potentially other conformations. A diamondlike seam may terminate in chiral columns.
A “bond line” is a linear arrangement of 2 or more side-by-side bonds possessing a generally parallel (but not necessarily a perfectly parallel) orientation.
A “graphenic network” herein describes a structure with a two-dimensional molecular-scale geometry that is at some larger scale three-dimensionally crosslinked. As a function of a graphenic network's crosslinking and network geometry, it cannot be broken without breaking some portion of its two-dimensional molecular structure. Graphenic networks comprise the broadest category of networks constructed from graphenic structures, as shown by this category's position at the apex of the classification chart in
A “layered” network is herein defined as a multilayer graphenic network comprising z-adjacent layers with either graphitic or nematic xy-alignment. Layered graphenic networks are shown as a subcategory of graphenic networks in the classification chart in
A “graphitic network” is herein defined as a type of layered graphenic network in which z-adjacent layers exhibit graphitic xy-alignment—i.e. they are substantially parallel. Graphitic networks may be characterized by an average <002> interlayer d-spacing of 3.45 Å or less, with no significant presence of interlayer spacings larger than 3.50 Å. Graphitic networks are shown as a subcategory of layered graphenic networks in the classification chart in
An “anthracitic network” is herein defined as a type of layered graphenic network comprising two-dimensional molecular structures crosslinked via certain characteristic structural dislocations, described herein as “anthracitic dislocations,” which include Y-dislocations, screw dislocations, and mixed dislocations having characteristics of both Y-dislocations and screw dislocations. Z-adjacent layers in anthracitic networks exhibit nematic alignment. Anthracitic networks may be characterized by a significant presence of <002> interlayer d-spacings larger than 3.50 Å. Anthracitic networks are shown as a subcategory of graphenic networks in the classification chart in
“Nematic alignment” is herein used to describe a molecular-scale, general xy-alignment between z-adjacent layers in a multilayer graphenic system. This term is typically used to denote a type of consistent but imperfect xy-alignment observed between liquid crystal layers, and we find it useful herein for describing the imperfect xy-alignment of z-adjacent layers in anthracitic networks. Nematic alignment may be characterized by a significant presence of <002> interlayer d-spacings larger than 3.50 Å.
An “spx network” is herein defined as a type of synthetic anthracitic network comprising a single, continuous graphenic structure, wherein the network is laterally and vertically crosslinked via diamondlike seams and mixed dislocations (e.g. chiral columns). In the context of maturation processes, an spx network may be described as an “spx precursor.”
Carbon spx networks can be further classified based on the extent of their internal grafting, which can be determined by the prevalence of its sp2-hybridized edge states prior to maturation. With respect to the extent of this grafting, a carbon spx network can be described as:
A “helicoidal network” is herein defined as a type of synthetic anthracitic network comprising screw dislocations. These screw dislocations may be formed via the maturation of chiral columns present in spx networks. Hence, an spx network may be described as an “spx precursor” of a helicoidal network. The derivation of helicoidal networks from spx precursors is indicated by the dashed arrow labeled “maturation” in the classification chart in
“Maturation” is herein defined as a structural transformation that accompanies the sp3-to-sp2 rehybridization of sp3-hybridized states in an spx precursor. Maturation of an spx precursor ultimately forms a helicoidal network; the extent of maturation is determined by the degree to which the sp3-to-sp2 rehybridization is completed. Maturation is progressive, so networks in intermediate states comprising both spx and helicoidal network features may be formed. Additionally, maturation may be localized; for instance, heating certain locations of the network, such as by laser, might cause localized maturation of the affected area.
A “highly mature” carbon helicoidal network is defined herein as a carbon helicoidal network having an average Du peak position that is at least 1340 cm−1 and is at least 8 cm−1 higher than that of its spx precursor.
An “x-carbon” is herein defined as a category of synthetic anthracitic networks constructed from graphene and comprising one of the following:
A “z-carbon” is herein defined as a category of synthetic anthracitic networks constructed from graphene and comprising one of the following:
A “helicoidal singleton” is herein defined as a singleton-type helicoidal network, wherein the helicoidal network comprises a single, ring-connected graphenic structure, and wherein the network is laterally and vertically crosslinked by screw dislocations.
A “helicoidal assembly” is herein defined as an assembly-type helicoidal network, wherein the helicoidal network comprises an assembly of multiple, helicoidal graphenic structures that are physically interlocked with one another via braid-like double helicoids (i.e. double screw dislocations).
An “spx preform” is a macroscopic assembly of distinct, spx precursors, referred to in this context as “spx microforms.” Various forming techniques may be used to impart a desired shape to an spx preform, such as an elongated, flat, or equiaxed shape.
A “macroform” is herein defined as a macroscopic, cohesive structure.
A “singleton-to-singleton” maturation is herein defined as a maturation process in which an spx precursor is matured to form a helicoidal singleton.
“A singleton-to-assembly” maturation is herein defined as a maturation process in which an spx precursor is disintegrated into a helicoidal assembly.
“Disintegration” is herein defined as the division of a singleton-type graphenic network into two or more distinct, ring-disconnected graphenic structures.
A “primordial domain” is defined herein as a graphenic domain nucleated and grown over a substrate prior to any tectonic encounters. When primordial domains are grown over a common surface toward one another, their edges may have a tectonic encounter.
A “primordial region” is defined herein as a region of a graphenic network generally coinciding with the network's primordial domains. We generally refer to a primordial region when describing some region of a graphenic system that was originally a primordial domain.
A “tectonic encounter” is a state of lateral near-contact between two edge segments during growth of a two-dimensional lattice. A tectonic encounter creates a tectonic interface between the two participating edge segments. The numerous tectonic encounters that may occur during the nucleation and growth of a graphenic system may be described as “tectonic activity.”
A “tectonic interface” is defined herein as the edge-to-edge interface formed by a tectonic encounter between two graphenic structures or regions.
A “zigzag-zigzag interface” is herein defined as a tectonic interface in which both of the edge segments are in the zigzag configuration.
A “zigzag-armchair interface” is herein defined as a tectonic interface in which one of the edge segments is in the zigzag configuration, while the other is in the armchair configuration.
An “offset zone” is herein defined as an interfacial zone within a tectonic interface in which one of the two participating edge segments are vertically offset—i.e. one of the edge segments is located above the other.
A “level zone” is herein defined as an interfacial zone within a tectonic interface in which the two participating edge segments are substantially level with each other and sufficiently aligned such that a bond line of two or more laterally adjacent sp2-sp2 bonds may be formed across the interface, resulting in one or more sp2 ring-connections.
A “crossover point” is herein defined as a location in a tectonic interface where the two participating edge segments crisscross, and where their alignment is inadequate to form a bond line of two or more laterally adjacent sp2-sp2 bonds. This may be because the 2p2 orbitals of the opposing sp2 edge atoms are too misaligned for π bonds to form.
“Sp2 grafting” is herein defined as the formation of a sp2-sp2 bond line between two edge atoms. Sp2 grafting creates sp2 ring-connections that may cause distinct graphenic structures to become ring-connected and coalesce into a larger graphenic structure. Sp2 grafting across a tectonic interface is favored in level zones.
“Sp3 grafting” is herein defined as the formation of sp3-sp3 bonds between two edge atoms. This may involve the sp2-to-sp3 rehybridization of sp2 edge atoms. Sp3 grafting creates spx rings that may cause distinct graphenic structures to become ring-connected and coalesce into a larger graphenic structure. Sp3 grafting across a tectonic interface is favored in offset zones.
A “base” or “base layer” is herein defined as the first graphenic layer formed by grafting across the tectonic interfaces between primordial domains during pyrolytic growth.
“Mesoscale” is used herein to describe a hierarchical level or feature (e.g. crosslinking, porosity) pertaining to a relatively larger size-scale than the molecular features. For example, a perimorphic framework's mesoscale crosslinking is a function of its crosslinking over size-scales more relevant to a discussion of its particle morphology than to a discussion of its molecular bonding structure.
A “micropore” is herein defined as a pore with a diameter of less than 2 nm, following IUPAC convention. A “microporous” structure or phase is characterized by the presence of micropores.
A “mesopore” is herein defined as a pore with a diameter between 2 nm and 50 nm, following IUPAC convention. A “mesoporous” structure or phase is characterized by the presence of mesopores.
A “macropore” is herein defined as a pore with a diameter of greater than 50 nm, following IUPAC convention. A “macroporous” structure or phase is characterized by the presence of macropores.
An “ambient superconductor” is herein defined as a material or article capable of entering a superconducting state at a temperature above 0° C. and an external pressure between 0 and 2 atm. “Ambient superconductivity” is herein defined as a superconducting state at a temperature above 0° C. and an external pressure between 0 and 2 atm.
Pyrolysis involves the decomposition of a gas, liquid, or solid carbonaceous material and may be used to form graphenic structures. In some pyrolysis procedures, this decomposition occurs over a substrate surface. The substrate may comprise the simple, flat surface of a foil or the more complex surfaces of particles. The graphenic systems synthesized on particles may inherit some of the particles' morphological attributes. In the '918 and '760 Applications, we define a number of terms related to template-directed synthesis. These terms are defined below.
A “template,” as defined herein, is a potentially sacrificial structure that imparts a desired morphology to another material formed in or on it. Of relevance for surface replication techniques are the template's surface (i.e. the “templating surface”), which is positively replicated, and its bulk phase (i.e. the “templating bulk”), which is negatively replicated. The template may also perform other roles, such as catalyzing the formation of the perimorphic material. A “templated” structure is one that replicates some feature of the template.
A “perimorph” or “perimorphic” material is a material formed in or on a solid-state or “hard” template material.
“Surface replication,” as defined herein, comprises a templating technique in which a template's surface is used to direct the formation of a thin, perimorphic wall of adsorbed material, the wall substantially encapsulating and replicating the templating surface upon which it is formed. Subsequently, upon being displaced, the templating bulk is replicated, in negative, by an endocellular space within the perimorphic wall. Surface replication creates a perimorphic framework with a templated pore-and-wall architecture.
A “perimorphic framework” (or “framework”), as defined herein, is the nanostructured perimorph formed during surface replication. A perimorphic framework comprises a nanostructured “perimorphic wall” (or “wall”) that may range from less than 1 nm to 100 nm in thickness but is preferably between 0.6 nm and 5 nm. Insomuch as it substantially encapsulates and replicates the templating surface, the perimorphic wall can be described as “conformal.” Perimorphic frameworks may be made with diverse architectures, ranging from simple, hollow architectures formed on nonporous templates to labyrinthine architectures formed on porous templates. They may also comprise different chemical compositions. A typical framework may be constructed from carbon and may be referred to as a “carbon perimorphic framework.”
An “endomorph,” as defined herein, comprises a template as it exists within a substantially encapsulating perimorphic phase. Therefore, after the perimorphic phase has been formed around it, the template may be described as an endomorph, or as “endomorphic.”
A “perimorphic composite,” as defined herein, is a composite structure comprising an endomorph and a perimorph. A perimorphic composite material may be denoted x@y, where x is the perimorphic element or compound and y is the endomorphic element or compound. For example, a perimorphic composite comprising a carbon perimorph on an MgO endomorph might be denoted C@MgO.
Numerous template elements or compounds may be employed, including carbon, metal oxides, oxyanionic salts, boron nitride, metal halides, and more. In particular, magnesium oxide (MgO) templates are often employed in chemical vapor deposition (“CVD”) processes due to their stability at high temperatures. Many of these templates are described in the '918 Application and the '154 Application. All that is required for many surface replication procedures involving CVD is a surface and the nucleation of a lattice that can be grown via autocatalysis or as a free radical condensate.
In the free radical condensate theory of growth, a free radical condensate (i.e. “condensate” or “FRC”) is formed during pyrolytic decomposition of a reactive vapor. A carbon FRC is a charged, hydrogenated precursor to the graphenic structure that can rapidly rearrange its carbon skeleton without breaking covalent bonds; hence it can be envisioned as a kind of charged, covalent liquid. A carbon FRC grows in the presence of a reactive vapor via radical addition reactions at its edges. As the condensate releases molecular hydrogen, its concentration of radicals diminishes, its self-rearrangement ceases, and it becomes an uncharged carbon structure. A gradual release of molecular hydrogen provides the FRC more time to rearrange itself into an energy-minimizing configuration-typically one that eliminates high-energy edge defects. This has been shown to promote edgeless graphenic structures like fullerenes. A sudden loss of hydrogen, by contrast, does not provide sufficient time for these energy-minimizing rearrangements to occur, which promotes the formation of graphenic structures with more edges.
If grown over a common substrate surface, graphenic structures may come into lateral contact with one another. These tectonic encounters, and the underlying factors that determine how they are resolved, have been the subject of scant research. In one case we have found, researchers observing the growth of ring-ordered, crystalline graphenic structures on copper foil found that a tectonic encounter could be resolved in one of two ways, as illustrated in
In the first scenario, the edge of one of the graphenic structures is subducted by the edge of the other—an event described herein as a “subduction event.” A subduction event allows continued growth of the subducting region over the subducted region, as illustrated in
In the second scenario described by the researchers, the edge of one of the graphenic structures may graft to the edge of the other via sp2-sp2 bond formation between the opposing edge atoms. This sp2 grafting causes the two graphenic structures to coalesce to form a larger graphenic structure. The outcome of this event is illustrated in
The complexity of tectonics between graphenic structures is increased when the substrate surface becomes more topologically and topographically complex. It is further increased if we postulate edge disorder. We surmise herein that these factors are important in determining the outcomes of tectonic encounters. Lastly, it is increased if the tectonics occur in a substantially unconfined space, where steric effects of surrounding structures can be ignored. This may not be the case when pyrolysis occurs in certain microporous template particles, like Zeolite Y, where sp2 grafting between graphenic structures (as opposed to subduction) may be forced due to the z-directional confinement in these templates' micropores—i.e. a lack of overhead clearance.
To describe the local space around curved, two-dimensional graphenic structures, it is helpful to establish an intuitive orientation. On a curved surface, there exists some tangent plane at any given point that we can think of as an xy plane.
An example of a ring-disordered graphenic domain with nonzero curvature is modeled in
From the vertical perspective in
Analysis of exemplary systems may provide helpful clarification of these concepts. Unless stated otherwise, the models all depict sp2-hybridized or sp3-hybridized carbon atoms and do not show hydrogen atoms.
The side of RA labeled x in
In the system in
Next, we evaluate the atoms of the graphenic structure in
Atom 1 belongs to rings RA and RB, which do not completely surround it. Therefore, 1 meets the definition of an edge atom. Atoms 2 through 18 also meet this definition. Edge atoms are indicated by diagonal patterning in
For instance, starting from 1, we find that 2 is a nearest neighbor, an edge atom, and a co-member (along with 1) of exactly one ring (RA). Continuing this trace from 2 to 18, a circuit is formed that is closed by the bond between 18, the last atom in the chain, and 1, its nearest neighbor and the first atom in the chain. Together, these atoms represent the edge of the graphenic structure.
In
Next the atoms of the graphenic structure in
In
Therefore, the system in
In
Since the system in
In
In
Therefore, the graphenic system in
In
A number of different instruments were employed to characterize the materials synthesized in the present disclosure. The following discussion provides information on these instruments and context for how we analyzed the related data.
All Raman spectroscopic characterization was performed using a ThermoFisher DXR Raman microscope equipped with a 532 nm excitation laser and Omnic profile-fitting software. Specific laser powers were used and are specified where applicable.
Raman spectroscopy is commonly used to characterize the molecular structure of carbon, and a prolific literature exists on this subject. Two main spectral features are typically associated with optical excitation of sp2-hybridized carbon: the G band (typically exhibiting a peak intensity value at approximately 1580 cm−1 to 1585 cm−1 in graphitic sp2 carbon), and the D band (exhibiting a peak intensity value at approximately 1350 cm−1 under optical excitation). The “2D” band representing a second order of the D band is also observed in some graphitic carbons, and its peak intensity value is typically located at approximately 2700 cm−1. The G band is assigned to the vibrations of sp2-sp2 bonds. The D band is assigned to the radial breathing mode of sp2-hybridized carbon atoms arranged in rings, and for Raman observation this requires back-scattering of electrons at a defect site.
Researchers have described an amorphization trajectory in the spectra of graphitic carbon showing a progression in disorder from graphite to amorphous carbon that is helpful to understand the dynamics of the D band. In graphite, no D peak is present due to the absence of activating defects. In carbons comprising smaller graphenic domains, the density of edge states is increased, and as edge states increase the D peak is activated by backscattering at the edge defects. The D peak intensity increases toward a maximum, corresponding to a nanocrystalline graphite. Further amorphization in the form of ring disorder diminishes the intensity of the D peak. Lastly, the D peak disappears as further amorphization eliminates a polycyclic, sp2-hybridized structure altogether.
The Raman spectral peaks associated with sp3-hybridized carbon include a peak at 1306 cm−1 (associated with hexagonal diamond), a peak at 1325 cm−1 (associated with hexagonal diamond) and a peak at 1332 cm−1 (associated with cubic diamond). Cubic diamond comprises 100% chair conformations, whereas hexagonal diamond comprises both chair conformations and boat conformations, giving it a lower Raman frequency and lower thermodynamic stability.
Raman-active phonons are known to be strain-dependent. Because the presence of strain within a lattice causes the lattice's vibrational frequencies to shift, Raman spectroscopy can be utilized to understand the strain states within a lattice. However, strain can also shift spectral peaks from their normally identified positions to new positions, making identification more ambiguous. The primary indicator of strain in a ring-ordered graphene structure is the position of the G peak and 2D peak, both of which are sensitive to tension and compression. The G peak has been shown to shift to lower frequencies (i.e. a “red-shift”) when the sp2-sp2 bonds are stretched and to higher frequencies (i.e. a “blue-shift”) when they are compressed. In graphenic structures with non-uniform strain fields, multiple modes of the G band may be present.
In disordered carbons, several Raman spectral features have been observed in addition to the D peak. A broad Raman peak (sometimes referred to as D″) often fitted between 1500 cm−1 and 1550 cm−1 in amorphous sp2-hybridized carbons is generally observed to increase with ring disorder. It is herein attributed to low-correlation, red-shifted modes of the G band associated with stretched, weakened sp2-sp2 bonds, which proliferate as ring disorder and lattice distortion increase in sp2-hybridized graphenic structures. Ferrari & Robertson have shown that the G peak red-shifts into this range in Stage II of the amorphization trajectory. In graphene oxide, this red-shifted mode of the G peak may be found alongside the normal G peak, indicating the presence of weaker sp2 bonds alongside normal sp2 bonds within the lattice. This is in good agreement with the customary interpretation of graphene oxide as a non-uniform lattice with both ring-disordered and ring-ordered regions.
Another feature (referred to as the D′ peak) observed in disordered carbons is fitted at 1620 cm−1, where it may appear as a shoulder on the G peak. This feature is often observed to accompany the D peak in sp2-hybridized carbons with a high density of edge states, and its intensity relative to the D peak has been shown to increase in proximity to lattice edges.
Another feature observed in disordered carbons, sometimes referred to as the D* peak, is a broad band fitted between 1100 cm−1 and 1200 cm−1. A peak intensity value at 1175 cm−1 within this range has been attributed to the sp2-sp3 bonds formed between sp2 and sp3 atoms at the transitions between sp2 and sp3 networks found within soot. It has also been attributed to hexagonal diamond. The assignment of this peak to sp3 carbon in nanodiamond and diamondlike materials by some researchers has been disputed by Ferrari & Robertson, who provided evidence that it should be assigned, along with a broad peak at ˜1240 cm−1, to trans-polyacetylene, a protonated aliphatic sp2 chain arguably present in those carbons.
In the present disclosure, Raman spectral analysis may involve reference to unfitted or fitted spectral features. “Unfitted” spectral features pertain to spectral features apparent prior to deconvolution via profile-fitting software. Unfitted features may therefore represent a convolution of multiple underlying features, but their positions are not subjective. “Fitted” spectral features pertain to the spectral features assigned by profile-fitting software. Imperfect profile fitting indicates the potential presence of other underlying features that have not been deconvoluted.
For clarity, features pertaining to the unfitted Raman profile are labeled with a subscript “u”—e.g. the “Gu” band. In the present disclosure, profile fitting is performed using OMNIC Peak Resolve software to deconvolute features contributing to the overall spectral profile. These fitted features are labeled with an “f”—e.g. the “Df” band. The software's Gaussian-Lorentzian lineshape setting was used by default, allowing a fitted band to adopt a Gaussian and Lorentzian character, with the fractional Gaussian character being determined by the software in order to optimize the fit. Other profile-fitting methods may change the locations, intensities, and trends of fitted peaks.
An additional unfitted feature defined within the present disclosure is the trough (“Tr”), a region of lower Raman intensity values located between the Du and Gu bands in the overall spectral profile. The Tru intensity is defined as the minimum intensity value occurring between the Du peak and the Gu peak. The trough intensity value indicates underlying spectral dynamics such as red-shifting of the G band corresponding to ring disorder and lattice distortion and can be analyzed without resorting to subjective profile-fitting judgments, making it a practically useful feature.
Averaged Raman spectra, where utilized herein, represent the average of multipoint spectral measurements made of the sample over a rectangular grid. The distinct point spectra are normalized and then averaged to create a composite spectrum.
X-Ray Diffraction of the carbon powders was performed by EAG Laboratories. XRD data was collected by a coupled Theta:2-Theta scan on a Rigaku Ultima-III diffractometer equipped with copper x-ray tube with Ni beta filter, parafocusing optics, computer-controlled slits, and a D/teX Ultra 1D strip detector. Profile fitting software was used to determine the peak positions and widths.
Thermogravimetric (TGA) analysis of the carbon powders was performed on a TA Instruments Q600 TGA/DSC. Thermal oxidation studies were performed by heating the powder samples in air.
Transmission Electron Microscope (TEM) imaging was performed on an FEI Tecnai F20 operated at 200 kV. A 300 mesh Copper Grid with lacey carbon was used. All samples were prepared in ethanol and allowed to dry at room temperature.
Gas adsorption data may be collected by a Micromeritics Tristar II Plus, measuring nitrogen adsorption at 77 K between pressures of
with increments ranging from
Micromeritics MicroActive software may be used to calculate the BET specific surface area, derived from the BET monolayer capacity assuming the cross-sectional area of 0.162 nm2. All samples were preconditioned by degassing with continuously flowing dry nitrogen gas at 100° C. prior to analysis except samples F2 and F3 which were degassed at 200° C. prior to analysis.
The pore size distribution (PSD) and cumulative volume of pores is another technique that may be performed from gas adsorption data to lend insight into the sintering behavior of particles. The data was collected by a Micromeritics Tristar II Plus, measuring nitrogen adsorption and desorption at 77 K between pressures of
with increments ranging from
Samples were preconditioned by degassing with continuously flowing dry nitrogen gas at 100° C. prior to analysis.
Micromeritics MicroActive software may be used to calculate adsorption-desorption PSD and cumulative volume of pores by applying the Barrett, Joyner and Halenda (BJH) method. This method provides a comparative assessment of mesopore size distributions for gas adsorption data. For all BJH data, the Faas correction and Harkins and Jura thickness curve may be applied. The cumulative volume of pores may be measured for both adsorption and desorption portions of the isotherm.
The following discussion summarizes the procedures used to complete each study (i.e. Study A through Study G). We generally endeavor to label samples according to the Study with which they are most associated—i.e. Sample A1 is the first sample associated with Study A. Within a single experiment, multiple samples may be evaluated, and multiple procedures may have been performed to create the samples. The procedures and samples are labeled the same—e.g. “Sample B2” is made via “Procedure B2”.
The present disclosure employs exemplary procedure. Other procedures, including those employing pyrolysis of alternative solid- or liquid-state carbonaceous precursor materials, the use of alternative substrates or catalysts, or other basic parameters, might be used as substitutes for those described herein without deviating from the inventive concept. In order to establish the versatility of the method, the mechanics of synthesis, and certain observable trends that might be exploited, a number of exemplary x-carbon synthesis procedures have been performed.
Procedures—Study A
For Procedures A1, A2, and A3, a rotary tube furnace may be employed with a quartz tube. The quartz tube may be a 60 mm OD quartz tube containing a middle 12″ section of 100 mm OD tube (the “belly”) positioned within the furnace's heating zone as shown in
For Procedures A4 and A5, a tube furnace may be employed with a quartz tube. The quartz tube may be a 60 mm OD tube. The furnace may be kept level (i.e. not tilted). Ceramic blocks may be inserted on each side of the furnace's heating zone (with the powder sample being placed between the blocks and inside the heating zone). The powder sample may be placed in open ceramic boats inside the tube. The tube may be fitted with two stainless steel flanges. Gas may be flowed in through a gas inlet on one flange and out through a gas outlet in the other flange.
Using the furnace configurations described above, five carbon samples may be synthesized utilizing the following procedures:
Procedure A1: A 500 g sample of “Elastomag 170” (a commercial magnesia powder supplied by Akrochem) magnesium oxide template precursor powder may be loaded into the quartz tube inside the tube furnace's heating zone. The rotary tube furnace may be set to a non-rotating mode. While under 500 sccm flow of argon (Ar) gas, the furnace may be heated from room temperature to a temperature setting of 1,050° C. over 50 minutes. Under sustained Ar gas flow, the furnace may then be allowed to cool to 750° C. over the next 30 minutes. During this period, the MgO template precursor morphology may be changed due to calcination into the desired template morphology. This condition may be held for an additional 30 minutes, after which a 250 sccm flow of propylene (C3H6) gas may be initiated, while holding the Ar flow unchanged, and this condition may be held for 60 minutes. The C3H6 flow may then be discontinued and the furnace allowed to cool to room temperature under sustained Ar flow. At this point, the C@MgO perimorphic composite powder as synthesized may be analyzed via Raman spectroscopy or thermogravimetric analysis (TGA). The MgO template may then be selectively extracted from the C@MgO perimorphic composite powder by acid-etching with hydrochloric acid (HCl) under magnetic stirring conditions, resulting in a mixture of carbon in an aqueous MgCl2 solution. The carbon may then be filtered from the solution, rinsed with deionized water three times, and dried to form a carbon powder. A carbon powder made via such a procedure is herein referred to as “Sample A1.”
Procedure A2: A 500 g sample of Elastomag 170 (a commercial magnesia powder supplied by Akrochem) magnesium oxide (MgO) template precursor powder may be loaded into the quartz tube inside the tube furnace's heating zone. The rotary tube furnace may be set to a non-rotating mode. While under 500 sccm flow of Ar gas, the furnace may be heated from room temperature to a temperature setting of 1,050° C. over 50 minutes, and then held at this condition for 30 minutes. During this period, the MgO template precursor morphology may be changed due to calcination into the template morphology desired. Next, a 500 sccm flow of methane (CH4) gas may be initiated while holding Ar flow unchanged, and this condition may be held for 30 minutes. The CH4 flow may then be discontinued and the furnace allowed to cool to room temperature under sustained Ar flow. At this point in the procedure, the C@MgO perimorphic composite powder as synthesized may be analyzed via Raman spectroscopy or thermogravimetric analysis (TGA). The MgO template may then be selectively extracted from the C@MgO perimorphic composite powder by acid-etching with HCl under magnetic stirring conditions, resulting in a mixture of carbon in an aqueous MgCl2 solution. The carbon may then be filtered from the solution, rinsed with deionized water three times, and dried to form a carbon powder. A carbon powder made via such a procedure is herein referred to as “Sample A2.”
Procedure A3: An MgO powder may be generated by calcining Light Magnesium Carbonate (a commercial hydromagnesite powder supplied by Akrochem) for 2 hours at a temperature of 1,050° C. for 2 hours. A 300 g sample of the pre-calcined powder may be loaded into the quartz tube inside the tube furnace's heating zone. The rotary tube furnace may be set to rotate at 2.5 RPM. While under 500 sccm flow of Ar gas, the furnace may be heated from room temperature to a temperature setting of 650° C. over 30 minutes, and then held at this condition for 30 minutes. Next, a 270 sccm flow of C3H6 gas may be initiated while holding Ar flow unchanged, and this condition may be held for 60 minutes. The C3H6 flow may then be discontinued and the furnace allowed to cool to room temperature under sustained Ar flow.
At this point in the procedure, the C@MgO perimorphic composite powder as synthesized may be analyzed via Raman spectroscopy or thermogravimetric analysis (TGA). The MgO template may then be selectively extracted from the C@MgO perimorphic composite powder by acid-etching with HCl under magnetic stirring conditions, resulting in a mixture of carbon in an aqueous MgCl2 solution. The carbon may then be filtered from the solution, rinsed with deionized water three times, and dried to form a carbon powder. A carbon powder made via such a procedure is herein referred to as “Sample A3.”
Procedures—Study B
For Procedures B1-B3, an MgO powder may be generated by calcining a template precursor powder comprising rhombohedral magnesite (MgCO3) crystals. The precursor powder may be calcined in a Vulcan 3-550 Muffle Furnace at a temperature of 580° C. for an hour followed by 1,050° C. for 3 hours with heating ramp rates of 5° C./min.
For Procedure B4, an MgO powder may be generated by calcining a template precursor powder comprising light magnesium carbonate crystals. The precursor powder may be calcined in a Vulcan 3-550 Muffle Furnace at a temperature of 750° C. for an hour with a heating ramp rate of 5° C./min.
For Procedures B1-B3, an MTI rotary tube furnace may be employed with a quartz tube. The quartz tube may be a 60 mm OD quartz tube containing a middle 12″ section of 100 mm OD tube (the “belly”) positioned within the furnace's heating zone as shown in
For Procedure B4, a tube furnace may be employed with a quartz tube. The quartz tube may be a 60 mm OD tube. The furnace may be kept level (i.e. not tilted). Ceramic blocks may be inserted on each side of the furnace's heating zone (with the powder sample being placed between the blocks and inside the heating zone). The powder sample may be placed in open ceramic boats inside the tube. The tube may be fitted with two stainless steel flanges. Gas may be flowed in through a gas inlet on one flange and out through a gas outlet in the other flange.
Procedure B1: The CVD procedure may be performed for 16 hours at a temperature of 640° C. under flowing gas conditions. The flowing gas may comprise 1,220 sccm CO2 and 127 sccm C3H6. The quartz tube may be rotated at 1 rpm. After cooling the resulting C@MgO powder to room temperature under flowing CO2, the MgO template may be selectively extracted from the C@MgO perimorphic composite powder by acid-etching with HCl under magnetic stirring conditions, resulting in a mixture of carbon in an aqueous MgCl2 solution. The carbon may then be filtered from the solution, rinsed with deionized water three times, and dried to form a carbon powder. A carbon powder made via such a procedure is herein referred to as “Sample B1.”
Procedure B2: The CVD procedure may be performed for 20 hours at a temperature of 580° C. under flowing gas conditions. The flowing gas may comprise 1,220 sccm CO2 and 127 sccm C3H6. The quartz tube may be rotated at 1 rpm. After cooling the resulting C@MgO powder to room temperature under flowing CO2, the MgO template may be selectively extracted from the C@MgO perimorphic composite powder by acid-etching with HCl under magnetic stirring conditions, resulting in a mixture of carbon in an aqueous MgCl2 solution. The carbon may then be filtered from the solution, rinsed with deionized water three times, and dried to form a carbon powder. A carbon powder made via such a procedure is herein referred to as “Sample B2.”
Procedure B3: The CVD procedure may be performed for 32.5 hours at a temperature of 540° C. under flowing gas conditions. The flowing gas may comprise 1,220 sccm CO2 and 127 sccm C3H6. The quartz tube may be rotated at 1 rpm. After cooling the resulting C@MgO powder to room temperature under flowing CO2, the MgO template may be selectively extracted from the C@MgO perimorphic composite powder by acid-etching with HCl under magnetic stirring conditions, resulting in a mixture of carbon in an aqueous MgCl2 solution. The carbon may then be filtered from the solution, rinsed with deionized water three times, and dried to form a carbon powder. A carbon powder made via such a procedure is herein referred to as “Sample B3.”
Procedure B4: The CVD procedure may be performed for 1 hour at a temperature of 580° C. under flowing gas conditions. The flowing gas may comprise 1,138 sccm CO2 and 276 sccm C2H2. After cooling the resulting C@MgO powder to room temperature under flowing CO2, the MgO template may be selectively extracted from the C@MgO perimorphic composite powder by acid-etching with HCl under magnetic stirring conditions, resulting in a mixture of carbon in an aqueous MgCl2 solution. The carbon may then be filtered from the solution, rinsed with deionized water three times, and dried to form a carbon powder. A carbon powder made via such a procedure is herein referred to as “Sample B4.”
Procedures—Study C
For Procedures C1 and C2, an MgO powder may be generated by treating a template precursor powder comprising sodium doped elongated nesquehonite template precursor crystals. The sodium doped nesquehonite template precursor may be precipitated from a solution stock of magnesium bicarbonate solution. First, in a 57 liter pressure vessel a mixture of concentration 0.62 mol kg−1 Mg comprised of magnesium hydroxide (Akrochem Versamag) and deionized water may be prepared. This mixture may be recirculated while carbonated with CO2 up to 60 psig to form a solution stock of magnesium bicarbonate (Mg(HCO3)2). After approximately 22 hours, the solution may be filtered to remove undissolved solids. The resulting solution stock may have a concentration of 0.29 mol kg−1 Mg. Then, sodium bicarbonate (NaHCO3) may be added to the solution stock to bring the concentration of sodium in the system to 1.7·10−3 mol kg−1 Na. Additional CO2 may be added to the vessel for 20 minutes to digest any unwanted precipitant. The system may be heated up to 34° C. and depressurized to allow for crystallization over 25.5 hours. The mixture generated from crystallization of sodium doped elongated nesquehonite template precursor crystals may then be filtered, rinsed with deionized water and acetone, and dried in a 45° C. in a forced air recirculation oven. The template precursor may be used as is in the CVD Replication step and conversion to MgO occurs in-situ during the heating ramp stage.
For Procedures C1 and C2, a tube furnace may be employed with a quartz tube. The quartz tube may be a 60 mm OD tube. The furnace may be kept level (i.e. not tilted). Ceramic blocks may be inserted on each side of the furnace's heating zone (with the powder sample being placed between the blocks and inside the heating zone). The powder sample may be placed in open ceramic boats inside the tube. The tube may be fitted with two stainless steel flanges. Gas may be flowed in through a gas inlet on one flange and out through a gas outlet in the other flange.
Procedure C1: A 1.6 g sample of sodium doped elongated nesquehonite template precursor may be loaded into the quartz tube inside a tube furnace's heating zone. While under 1271 sccm flow of Ar gas, the furnace may be heated from room temperature to a temperature setting of 460° C. over 20 minutes, and then held at this condition for 15 minutes to equilibrate. Next, a 42 sccm flow of C2H2 gas may be initiated while holding Ar flow unchanged, and this condition may be held for 3 hours. The C2H2 flow may then be discontinued and the furnace allowed to cool to room temperature under sustained Ar flow with the resulting C@MgO powder may be collected. The MgO template may be selectively extracted from the C@MgO perimorphic composite powder by acid-etching with HCl under magnetic stirring conditions, resulting in a mixture of carbon in an aqueous MgCl2 solution. The carbon may then be filtered from the solution, rinsed with deionized water three times, and dried to form a carbon powder. A carbon powder made via such a procedure is herein referred to as “Sample C1.”
Procedure C2: A 1.9 g sample of sodium doped elongated nesquehonite template precursor may be loaded into the quartz tube inside a tube furnace's heating zone. While under 1,271 sccm flow of Ar gas, the furnace may be heated from room temperature to a temperature setting of 400° C. over 20 minutes, and then held at this condition for 15 minutes to equilibrate. Next, a 105 sccm flow of C2H2 gas may be initiated while holding Ar flow unchanged, and this condition may be held for 3 hours. The C2H2 flow may then be discontinued and the furnace allowed to cool to room temperature under sustained Ar flow with the resulting C@MgO powder may be collected. The MgO template may be selectively extracted from the C@MgO perimorphic composite powder by acid-etching with HCl under magnetic stirring conditions, resulting in a mixture of carbon in an aqueous MgCl2 solution. The carbon may then be filtered from the solution, rinsed with deionized water three times, and dried to form a carbon powder. A carbon powder made via such a procedure is herein referred to as “Sample C2.”
Procedures—Study D
For Procedures D1 and D2, an MgO powder may be generated by calcining a template precursor powder comprising light magnesium carbonate crystals. The precursor powder may be calcined in a Vulcan 3-550 Muffle Furnace at a temperature of 750° C. for an hour with a heating ramp rate of 5° C./min.
For Procedures D1 and D2, a tube furnace may be employed with a quartz tube. The quartz tube may be a 60 mm OD tube. The furnace may be kept level (i.e. not tilted). Ceramic blocks may be inserted on each side of the furnace's heating zone (with the powder sample being placed between the blocks and inside the heating zone). The powder sample may be placed in open ceramic boats inside the tube. The tube may be fitted with two stainless steel flanges. Gas may be flowed in through a gas inlet on one flange and out through a gas outlet in the other flange.
Procedure D1: A 0.9 g sample of a magnesium oxide template precursor may be loaded into the quartz tube inside a tube furnace's heating zone. While under 1,271 sccm flow of Ar gas, the furnace may be heated from room temperature to a temperature setting of 700° C. over 30 minutes, and then held at this condition for 15 minutes to equilibrate. Next, a 20 sccm flow of C3H6 gas may be initiated while holding Ar flow unchanged, and this condition may be held for 30 minutes. The C3H6 flow may then be discontinued and the furnace allowed to cool to room temperature under sustained Ar flow with the resulting C@MgO powder may be collected. The MgO template may be selectively extracted from the C@MgO perimorphic composite powder by acid-etching with HCl under magnetic stirring conditions, resulting in a mixture of carbon in an aqueous MgCl2 solution. The carbon may then be filtered from the solution, rinsed with deionized water three times, and dried to form a carbon powder. A carbon powder made via such a procedure is herein referred to as “Sample D1.”
Procedure D2: A 0.9 g sample of a magnesium oxide template precursor may be loaded into the quartz tube inside a tube furnace's heating zone. While under 1,271 sccm flow of argon (Ar) gas, the furnace may be heated from room temperature to a temperature setting of 700° C. over 30 minutes, and then held at this condition for 15 minutes to equilibrate. Next, a combination of 20 sccm flow of propylene (C3H6) gas along with 60 sccm of hydrogen (H2) gas may be initiated while holding Ar flow unchanged, and this condition may be held for 30 minutes. The C3H6 flow may then be discontinued and the furnace allowed to cool to 150° C. under sustained Ar and H2 flow. The H2 flow may be discontinued below 150° C. and the furnace was allowed to cool to room temperature and the resulting C@MgO powder may be collected. The MgO template may be selectively extracted from the C@MgO perimorphic composite powder by acid-etching with HCl under magnetic stirring conditions, resulting in a mixture of carbon in an aqueous MgCl2 solution. The carbon may then be filtered from the solution, rinsed with deionized water three times, and dried to form a carbon powder. A carbon powder made via such a procedure is herein referred to as “Sample D2.”
Procedures—Study E
For Procedures E1 and E2 an MgO powder may be generated by calcining Light Magnesium Carbonate (a commercial hydromagnesite powder supplied by Akrochem) in a rotating kiln in 2 stages in an air atmosphere as shown in
For Procedures E1A and E2A a tube furnace may be employed with a quartz tube. An MTI rotary tube furnace with a 60 mm OD quartz tube may be employed for CVD. The furnace may be kept level (i.e. not tilted). Ceramic blocks may be inserted on each side of the furnace's heating zone (with the powder sample being placed between the blocks and inside the heating zone). Glass wool may be used to fix the position of the ceramic blocks. The tube may be fitted with two stainless steel flanges. Gas may be flowed in through a gas inlet on one flange and out through a gas outlet in the other flange. Powder samples may be placed in ceramic boats, and the boats may be placed in the heating zone prior to initiating the procedure. For Procedures E2 and E4 a similar setup may be employed with minor modifications to allow rapid heating and/or cooling of the samples. These modifications will be described in their respective exemplary procedures.
Procedure E1: A 50 mm OD quartz tube, serving as a boat, containing 62 grams of this pre-calcined MgO powder may be loaded into the tube. After initiating a 2,000 sccm flow of Ar gas, the furnace may be heated from room temperature to a temperature setting of 700° C. over 20 minutes and held at this condition for 15 minutes. Next, a 1,274 sccm flow of C3H6 gas may be initiated while maintaining Ar flow, and this condition may be held for 30 minutes. The C3H6 flow may then be discontinued and the furnace allowed to cool to room temperature under sustained Ar flow. The C@MgO perimorphic composite powder may be collected.
The MgO template may be selectively extracted from the C@MgO perimorphic composite powder by acid-etching with HCl under magnetic stirring conditions, resulting in a mixture of carbon in an aqueous MgCl2 solution. The carbon may then be filtered from the solution, rinsed with deionized water three times, and dried to form a carbon powder. A carbon powder made via such a procedure is herein referred to as “Sample E1.”
Procedure E1A: This procedure involves rapidly heating and cooling a perimorphic composite material from room temperature to the desired temperature setting. In a ceramic boat, a 3.0 g quantity of the perimorphic composite powder described in Procedure E1 may be loaded and placed in a quartz tube outside the heated zone of the furnace. After initiating a 4,000 sccm flow of Ar gas, the furnace may be heated from room temperature to a temperature setting of 900° C. over 45 minutes and held at this condition for 15 minutes. Until the temperature setting has been achieved the sample may be kept outside the heat zone. Once the desired temperature has been attained the boat is pushed in with the introduction of minimal additional air and left in the heat zone for 30 minutes followed by moving it back outside the heat zone in the quartz tube. This may serve to expose the sample to the desired temperature only for a short period of time. The furnace may be allowed to cool to room temperature under sustained Ar flow. The C@MgO perimorphic composite powder may be collected at room temperature.
The MgO template may be selectively extracted from the C@MgO perimorphic composite powder by acid-etching with HCl under magnetic stirring conditions, resulting in a mixture of carbon in an aqueous MgCl2 solution. The carbon may then be filtered from the solution, rinsed with deionized water three times, and dried to form a carbon powder. A carbon powder made via such a procedure is herein referred to as “Sample E1A.”
Procedure E2: A 50 mm OD quartz tube, serving as a boat, containing 74 grams of this pre-calcined MgO powder may be loaded into the tube. After initiating a 2,000 sccm flow of Ar gas, the furnace may be heated from room temperature to a temperature setting of 580° C. over 20 minutes and held at this condition for 15 minutes. Next, a 1,274 sccm flow of C3H6 gas may be initiated while maintaining Ar flow, and this condition may be held for 3 hours. The C3H6 flow may then be discontinued and the furnace allowed to cool to room temperature under sustained Ar flow. The C@MgO perimorphic composite powder may be collected.
The MgO template may be selectively extracted from the C@MgO perimorphic composite powder by acid-etching with HCl under magnetic stirring conditions, resulting in a mixture of carbon in an aqueous MgCl2 solution. The carbon may then be filtered from the solution, rinsed with deionized water three times, and dried to form a carbon powder. A carbon powder made via such a procedure is herein referred to as “Sample E2.”
Procedure E2A: This procedure involves gradually heating and rapidly cooling a perimorphic composite material from room temperature to the desired temperature setting and back to room temperature again. In a ceramic boat, a 3.0 g quantity of the perimorphic composite powder described in Procedure E3 may be loaded and placed in a quartz tube in the heated zone of the furnace. After initiating a 4,000 sccm flow of Ar gas, the furnace may be heated from room temperature to a temperature setting of 1,050° C. over 50 minutes and held at this condition for 15 minutes. The furnace may be held at this temperature for an hour. The furnace may then be allowed to start to cool under sustained Ar flow and the ceramic boat may be pulled out of the heat zone as soon as the heaters power off. The C@MgO perimorphic composite powder post may be collected once at room temperature.
The MgO template may be selectively extracted from the C@MgO perimorphic composite powder by acid-etching with HCl under magnetic stirring conditions, resulting in a mixture of carbon in an aqueous MgCl2 solution. The carbon may then be filtered from the solution, rinsed with deionized water three times, and dried to form a carbon powder. A carbon powder made via such a procedure is herein referred to as “Sample E2A.”
Procedures—Study F
For Procedures F1, an MgO powder may be generated by calcining a template precursor powder comprising light magnesium carbonate crystals. The precursor powder may be calcined in a Vulcan 3-550 Muffle Furnace at a temperature of 750° C. for an hour with a heating ramp rate of 5° C./min.
For Procedure F1, a Thermcraft tube furnace modified to be a rotary furnace may be employed with a quartz tube. The quartz tube may be a 60 mm OD quartz tube containing an expanded middle 577 mm section of 130 mm OD tube (the “belly”) positioned within the furnace's heating zone. Quartz baffles inside the belly may facilitate agitation of the powder. The furnace may be kept level (i.e. not tilted). The template sample may be placed inside the belly in the heating zone, with ceramic blocks inserted outside the belly on each side of the furnace's heating zone. Glass wool may be used to fix the position of the ceramic blocks. The template sample may be placed in the tube without the use of ceramic boats such that it allowed to rotate freely within the belly. The tube may be fitted with two stainless steel flanges. Gas may be flowed in through a gas inlet on one flange and out through a gas outlet in the other flange.
For Procedures F2, F3, F4, F5, F6 and F7 a tube furnace may be employed with a quartz tube. An MTI rotary tube furnace with a 60 mm OD quartz tube may be employed for CVD. The furnace may be kept level (i.e. not tilted). Ceramic blocks may be inserted on each side of the furnace's heating zone (with the powder sample being placed between the blocks and inside the heating zone). Glass wool may be used to fix the position of the ceramic blocks. The tube may be fitted with two stainless steel flanges. Gas may be flowed in through a gas inlet on one flange and out through a gas outlet in the other flange. Powder samples may be placed in ceramic boats, and the boats may be placed in the heating zone prior to initiating the procedure.
Procedure F1 and F2: A 150 g quantity of a magnesium oxide template powder maybe loaded into the belly of the quartz tube. After initiating a 1,379 sccm flow of CO2 gas and a tube rotation speed of 1 RPM, the furnace may be heated from room temperature to a temperature setting of 580° C. at a ramp-rate of 20° C./min and held at this condition for 15 minutes. Next, a 276 sccm flow of C2H2 gas may be initiated while maintaining CO2 flow, and this condition may be held for 180 minutes. The C2H2 flow may then be discontinued and the furnace allowed to cool to room temperature under sustained CO2 flow. The powder may be collected. The C@MgO perimorphic composite powder may be further processed to create a carbon powder. The MgO template may be selectively extracted from the C@MgO perimorphic composite powder by acid-etching with HCl under magnetic stirring conditions, resulting in a mixture of carbon in an aqueous MgCl2 solution. The carbon may then be filtered from the solution, rinsed with deionized water three times, further rinsed with ethanol three times and dried to obtain a carbon powder herein referred to as “Sample F1”.
A 50 mg quantity of the Sample F1 carbon powder may be compacted in a 7 mm die set (Pike Technologies 161-1010) under 105 ksi hydraulic pressure. Under pressure the carbon may form a pellet herein referred to as “Sample F2” that may be stable enough to handle.
Procedure F3: Sample F2 may be placed in a ceramic boat and loaded into the quartz tube of a furnace. After initiating a 4,000 sccm flow of Ar gas, the furnace may be heated from room temperature to a temperature setting of 1,050° C. over 50 minutes and held at this condition for 30 minutes. The furnace may then be allowed to cool to room temperature under sustained Ar flow. The pellet may be collected once at room temperature and is herein referred to as “Sample F3”.
Procedure F4: A 100 mg quantity of Sample F1 powder may be placed in a ceramic boat and loaded into the quartz tube of a furnace. After initiating a 4,000 sccm flow of Ar gas, the furnace may be heated from room temperature to a temperature setting of 1050° C. over 50 minutes and held at this condition for 30 minutes. The furnace may then be allowed to cool to room temperature under sustained Ar flow. The powder may be collected once at room temperature.
A 50 mg quantity of this powder may then be compacted in a 7 mm die set (Pike Technologies 161-1010) under 105 ksi hydraulic pressure. Under pressure the perimorphic carbon frameworks do not form a pellet and remain a powder, herein referred to as Sample F4.
Procedure F5: A potassium carbonate (K2CO3) template precursor may be spray dried using an Sinoped LPG-5 spray dryer. A room temperature solution composed of 250.35 g of K2CO3 and 1, 667.2 g of deionized water (DI) was pumped at a rate of 23 mL/min into a rotary atomizer set to 24,000 RPM. The inlet temperature of the spray dryer was set to 195° C., which produced an outlet temperature of 139° C. The powder collected after spray drying was a K2CO3 template precursor.
A 100 g quantity of this K2CO3 template precursor powder may be loaded into a ceramic boat and placed in a quartz tube to generate a perimorphic composite powder using an MTI tube furnace. After initiating a 1,220 sccm flow of CO2 gas, the furnace may be heated from room temperature to a temperature setting of 640° C. at a ramp-rate of 20° C./min and held at this condition for 15 minutes. Next, a 162 sccm flow of C3H6 gas may be initiated while maintaining CO2 flow, and this condition may be held for 2 minutes. The C3H6 flow may then be discontinued and the furnace allowed to purge with Ar at a flow rate of 2,000 sccm for 30 minutes to clear all the CO2 present in the tube. The furnace may then be cooled to room temperature under sustained Ar flow. The powder may be collected. The C@K2CO3 perimorphic composite powder may be further processed to create a carbon powder. The K2CO3 template may be selectively extracted from the C@K2CO3 perimorphic composite powder by acid-etching with HCl under magnetic stirring conditions, resulting in a mixture of carbon in an aqueous KCl2 solution. The carbon may then be filtered from the solution, rinsed with deionized water three times to obtain an aqueous paste. This paste may be rinsed three times with ethanol to obtain an ethanol paste.
An ethanol paste of this carbon may be diluted with additional ethanol to create a very dilute mixture of 0.003 wt % carbon. This mixture may then be agitated with a high shear rotor stator homogenization processor, IKA T-25 digital Ultra-Turrax (UT), run at 12,000 RPM for 5 minutes. The mixture after agitation may be immediately poured over a glass frit vacuum filtration setup having a 47 mm diameter nylon filter (0.45 μm pore size) as the filtration medium. The vacuum filtration may be allowed to proceed undisturbed until all the liquid has been drained out. The vacuum is turned off and the filter with carbon may be dried in air in the vacuum filtration setup itself. Once dry, a flexible vdW assembly may release itself from the filter. This vdW assembly is herein referred to as “Sample F5”.
Procedure F6: Sample F5 may be placed in a ceramic boat and loaded into the quartz tube of a furnace. After initiating a 4,000 sccm flow of Ar gas, the furnace may be heated from room temperature to a temperature setting of 1,050° C. over 50 minutes and held at this condition for 30 minutes. The furnace may then be allowed to cool to room temperature under sustained Ar flow. The assembly may be collected once at room temperature and is herein referred to as “Sample F6”.
Procedure F7: Nesquehonite (MgCO3 3H2O) may be precipitated from lansfordite (MgCO3·5H2O) to produce elongated particles. A 45 g/L MgO equivalent magnesium bicarbonate (Mg(HCO3)2) solution may be prepared by high pressure dissolution of magnesium hydroxide (Akrochem Versamag) in carbonic acid at 720 psig. Lansfordite may be precipitated from this magnesium bicarbonate solution in a continuously stirred tank reactor (CSTR). The solution may be chilled to ˜14° C. and depressurized from 720 psig to 0 psig over 5 minutes while agitated at ˜700 RPM with a down pumping marine style impeller. Air may be continuously purged through the headspace at 4 SCFMair while chilled to ˜12° C. for 8 hrs. The solution may be allowed to stir at ˜350 RPM for an additional 18.5 hrs. The CSTR may then be heated to 34.5° C. while stirred at ˜720 RPM for 82 minutes. The solution may then be diluted with approximately 5 L of deionized water while continued heating to 43.8° C. for an additional 61 minutes. The contents of the CSTR may then by removed, filtered, and dried in a forced air circulation oven at 40° C. The resulting powder, identified herein as N2, are acicular crystals of nesquehonite.
An MgO powder may be generated by calcining N2 at 640° C. for 2 hours in an N2 gas flow of 2,000 sccm with a heating ramp-rate of 5° C./min in an MTI tube furnace with a 60 mm dia. quartz tube. A 2.4 g quantity of this MgO powder maybe loaded into a ceramic boat and placed in the quartz tube to generate C@MgO using an MTI tube furnace. After initiating a 815 sccm flow of CO2 gas, the furnace may be heated from room temperature to a temperature setting of 540° C. at a ramp-rate of 5° C./min and held at this condition for 15 minutes. Next, a 812 sccm flow of C2H2 gas may be initiated while maintaining CO2 flow, and this condition may be held for 2 minutes. The C2H2 flow may then be discontinued and the furnace allowed to purge with Ar at a flow rate of 1,698 sccm for 30 minutes to clear all the CO2 present in the tube. The furnace may then be heated to 900° C. at a ramp-rate of 20° C./min and held at this condition for 30 minutes. The furnace may then be cooled to room temperature under sustained Ar flow. The powder may be collected. The C@MgO perimorphic composite powder may be further processed to create a carbon powder. The MgO template may be selectively extracted from the C@MgO perimorphic composite powder by acid-etching with HCl under magnetic stirring conditions, resulting in a mixture of carbon in an aqueous MgCl2 solution. The carbon may then be filtered from the solution, rinsed with deionized water three times to obtain an aqueous paste. This paste may be rinsed three times with ethanol to obtain an ethanol paste.
An ethanol paste of this carbon may be diluted with additional ethanol to create a very dilute mixture of 0.003 wt % carbon. This mixture may then be agitated with a high shear rotor stator homogenization processor, IKA T-25 digital Ultra-Turrax (UT), run at 12,000 RPM for 5 minutes. The mixture after agitation may be immediately poured over a glass frit vacuum filtration setup having a 47 mm diameter nylon filter (0.45 μm pore size) as the filtration medium. The vacuum filtration may be allowed to proceed undisturbed until all the liquid has been drained out. The vacuum is turned off and the filter with carbon may be dried in air in the vacuum filtration setup itself. Once dry, a cohesive flexible buckypaper may release itself from the filter, herein referred to as “Sample F7.”
Procedures—Study G
Procedure G1: Magnesite (MgCO3) particles may be crystallized from a solution of magnesium bicarbonate to yield a powder of equiaxed template precursor particles.
An MTI rotary tube furnace may be employed with a quartz tube. The quartz tube may be a 60 mm OD quartz tube containing a middle 12″ section of 100 mm OD tube positioned within the furnace's heating zone. Quartz baffles inside the belly may facilitate agitation of the powder. The furnace may be kept level (i.e. not tilted). Ceramic blocks may be inserted on each side of the furnace's heating zone (with the powder sample being placed between the blocks and inside the heating zone). Glass wool may be used to fix the position of the ceramic blocks. The tube may be fitted with two stainless steel flanges. Gas may be flowed in through a gas inlet on one flange and out through a gas outlet in the other flange.
A 177 g quantity of the precipitated magnesite powder may be calcined to MgO at 640° C. for 10 min under Ar flow of 5 ft3/hr with heating ramp-rate of 20° C./min. The MgO powder already present in the quartz tube may be used to generate C@MgO using the furnace described. After initiating a 1,918 sccm flow of CO2 gas and a tube rotation speed of 1 RPM, the furnace may be heated from room temperature to a temperature setting of 640° C. at a ramp-rate of 20° C./min and held at this condition for 15 minutes. Next, a 127 sccm flow of C3H6 gas may be initiated while maintaining CO2 flow, and this condition may be held for 360 minutes. The C3H6 flow may then be discontinued and the furnace allowed to cool to room temperature under sustained CO2 flow.
The C@MgO perimorphic composite powder may be placed back in the tube in the same identical furnace/tube configuration for a second growth cycle. After initiating a 1,918 sccm flow of CO2 gas and a tube rotation speed of 1 RPM, the furnace may be heated from room temperature to a temperature setting of 640° C. at a ramp-rate of 20° C./min and held at this condition for 15 minutes. Next, a 127 sccm flow of C3H6 gas may be initiated while maintaining CO2 flow, and this condition may be held for 120 minutes. The C3H6 flow may then be discontinued and the furnace allowed to cool to room temperature under sustained CO2 flow.
The C@MgO perimorphic composite powder may be placed back in the tube in the same identical furnace/tube configuration for a third growth cycle. After initiating a 1,918 sccm flow of CO2 gas and a tube rotation speed of 1 RPM, the furnace may be heated from room temperature to a temperature setting of 640° C. at a ramp-rate of 20° C./min and held at this condition for 15 minutes. Next, a 127 sccm flow of C3H6 gas may be initiated while maintaining CO2 flow, and this condition may be held for 180 minutes. The C3H6 flow may then be discontinued and the furnace allowed to cool to room temperature under sustained CO2 flow.
The powder may be collected. The C@MgO perimorphic composite powder may be further processed to create a carbon powder. The MgO template may be selectively extracted from the C@MgO perimorphic composite powder by acid-etching with HCl under magnetic stirring conditions, resulting in a mixture of carbon in an aqueous MgCl2 solution. The carbon may then be filtered from the solution, rinsed with deionized water three times followed by a triple rinse with ethanol to obtain an ethanol paste. This paste may be dried to form a carbon powder.
This carbon powder may then be utilized for further CVD growth. An MTI rotary tube furnace may be employed with a quartz tube. The quartz tube may be a 60 mm OD quartz tube containing a middle 12″ section of 100 mm OD tube positioned within the furnace's heating zone. Quartz baffles inside the belly may facilitate agitation of the carbon powder. The furnace may be kept level (i.e. not tilted). Ceramic blocks may be inserted on each side of the furnace's heating zone (with the powder sample being placed between the blocks and inside the heating zone). Glass wool may be used to fix the position of the ceramic blocks. The tube may be fitted with two stainless steel flanges. Gas may be flowed in through a gas inlet on one flange and out through a gas outlet in the other flange. This assembly is shown in
After initiating a 1,918 sccm flow of CO2 gas and a tube rotation speed of 1 RPM, the furnace may be heated from room temperature to a temperature setting of 640° C. at a ramp-rate of 20° C./min and held at this condition for 15 minutes. Next, a 127 sccm flow of C3H6 gas may be initiated while maintaining CO2 flow, and this condition may be held for 180 minutes. The C3H6 flow may then be discontinued and the furnace allowed to cool to room temperature under sustained CO2 flow. The final mass of carbon powder collected, net of losses from migration into the glass wool, may be approximately 43.2 g. The carbon powder made via this procedure is herein referred to as “Sample G1.”
Procedures—Study H
Procedure H: An aqueous Mg(HCO3)2 solution may be produced by mixing 16 kg deionized water and 1.39 kg of a commercial-grade MgO powder (Versamag) in a pressure vessel equipped with an overhead stirring system and gas-inducing impeller. The mixture may be mixed at 700 RPM and cooled to 5° C. while being fed CO2 gas up to 850 psi for 2 hours. The resulting solution may be withdrawn from the pressure vessel at atmospheric pressure and fed at a rate of 56 mL/min into a BETE XA air atomizing nozzle comprising an FC7 Fluid Cap and AC1802 Air Cap. Compressed air for droplet atomization may be delivered into the nozzle at a flow rate of 5 SCFH air at 54 psi. The inlet temperature of the spray dryer may be set to 200° C., producing an outlet temperature ranging between 108° C. and 109° C. The ambient conditions during the spray drying process may be 28.4° C. and 48% RH. Approximately 1400 mL of solution may be sprayed, and 208 g of spray-dried, hydrous magnesium carbonate (Mg(CO3)·xH2O) template precursor powder with a hollow-spherical morphology may be collected via a cyclonic separator.
Next, the template precursor powder may be converted into a template via thermal treatment using a muffle furnace (Vulcan 3-550 Model, 1440 W max). Approximately 10 g of the template precursor powder may be placed in ceramic boats and heated to 580° C., then held at this temperature for 13.5 hours, followed by heating to 1050° C. and holding for another 1 hour to yield approximately 3.9 g of MgO powder. The heating ramp rates for both steps may be 5° C./min and the cool-down was allowed to happen naturally overnight over 8 hours. Approximately 0.47 g of the MgO powder may be pelletized in a 15.7 mm ID hydraulic press by applying 7.8 ksi of uniaxial compression for 1 minute. The resulting disc-shaped template may have a diameter of 15.7 mm and thickness of 2.5 mm.
Next, a Thermcraft tube furnace with a 60 mm OD quartz tube may be employed in a template-directed CVD procedure. The furnace may be kept level (i.e. not tilted), with the 0.47 g pelletized template sample being placed in a ceramic boat in the heating zone prior to initiating the procedure. Ceramic blocks may be inserted outside each side of the furnace's heating zone, and glass wool may be used to fix the position of the ceramic blocks. The tube may be fitted with two stainless steel flanges. Gas may be flowed in through a gas inlet on one flange and out through a gas outlet in the other flange. After initiating a 815 sccm flow of CO2 gas, the furnace may be heated from room temperature to a temperature setting of 540° C. at a ramp-rate of 20° C./min and held at this condition for 5 minutes. Next, a 144 sccm flow of C2H2 gas may be initiated while maintaining CO2 flow, and this condition may be held for 90 minutes. The C2H2 flow may then be discontinued, and the furnace allowed to cool to room temperature under sustained CO2 flow. During cooling, the clam-shell furnace lid may be opened completely, exposing the quartz tube to the outside air. A perimorphic composite pellet obtained after cooling may be characterized. Finally, the same CVD growth procedure may be repeated twice more, with the pellet being again cooled, for a total of 3 CVD growth steps with the pellet being allowed to cool between each step. The resulting perimorphic composite pellet comprises a macroscopic, perimorphic carbon that may be tested for ambient superconductivity.
A vacuum chamber like the one associated with the Cober-Muegge microwave system utilized in Study G (
The points of the 4-point probe may be placed into static contact with the flat surface of the macroform as lightly and delicately as possible to obtain a steady, continuous sheet resistance reading. This delicate placement should be done to avoid compressing the macroform surface with the probe tips, which may be necessary due to the apparent pressure-sensitivity of the spx macroforms we tested. We theorize that this pressure-sensitivity is attributable to localized mechanical compression reducing the interlayer distance and thereby inducing interlayer electronic coupling near the voltage-sensing points of contact. Additionally, a soft, non-conductive backing underneath the carbon macroform may be utilized in order to minimize local compression. To make contact, the Sourcemeter may be turned on to get an initial reading at ambient conditions, and the chamber may then be closed and evacuated. During the evacuation of the chamber, readings of the chamber pressure and the sample's sheet resistance may be noted.
SEM images of Sample A1 confirms the presence of perimorphic frameworks.
To achieve better transparency, and to study the smaller-scale structure of the perimorphic wall in Sample A1, TEM analysis was also performed.
Closer examination of the perimorphic wall is possible in a higher-magnification view, shown in
In the highest magnification view, shown in
Care must be taken during HRTEM analysis that the fringe lines corresponding to the actual positions of the graphenic layers are not confused with the fringe lines corresponding to the z-intervals between these layers. Depending on the defocus value, the fringes associated with the actual atomic positions may be either dark or bright. Whichever color they are, the lines associated with the z-intervals will be the opposite color. In the literature, we can find examples of either dark or bright fringes being associated with graphenic layers. In order to make a confident assignment of the exact atomic positions in HRTEM images, it helps to have corroborating information about the actual molecular structure.
The presence of fringe lines indicates that this section of the perimorphic wall in Sample A1 comprises a stacked arrangement of z-adjacent graphenic regions. In the main frame of
An xy-alignment between z-adjacent graphenic regions allows smaller z-intervals and higher-density arrangements, which should in turn increase interlayer coupling and vdW cohesion. We consider this a desirable feature of a layered graphenic system as opposed to the lower-density, nonlayered network architecture exhibited by schwarzite. If density reduction is desired, this can be accomplished by introducing larger-scale modes of porosity (such as the macropores in Sample A1), while preserving a high-density layered organization at smaller scales.
Another helpful example of nematic alignment is shown in
While the layers throughout Sample A1 are nematically aligned, it is visually difficult to trace dark fringe lines in
In addition to the z-intervals between the black tracing segments, there appear to be lateral discontinuities separating the black tracing segments in the magnified inset of
The alternative (and correct) explanation is that the bright fringes (corresponding to the white tracing in the magnified inset of
This observation has a precedent in the anthracite literature. HRTEM fringes of anthracite have been analyzed to generate a model of anthracite's structural dislocations.
The simulated HRTEM fringe pattern formed by a Y-dislocation is shown below the dislocation in
Geologically-formed anthracitic networks are a natural demonstration of how structural dislocations can create a three-dimensional graphenic network. Substantially all of the carbon atoms in anthracite are members of the graphenic network resulting from these crosslinking dislocations, with the exception of an occasional CH, CH2 or CH3 group (which solid state C NMR has indicated are present only in very small quantities) attached to a ring. It is this crosslinking of the graphenic network that lends anthracite its hardness and that prevents its exfoliation or solubilization. NMR spectroscopy has been used to show that dodecylation of anthracite only affects the edge atoms of this singleton, wherein “the graphenic layers appear to merge.”
Returning to the fringe pattern shown in the magnified inset of
The case for this is further reinforced by our comparative analysis of Samples A2 and A3. Namely, if the perimorphic frameworks in Sample A1 comprised vdW assemblies, the conspicuously superior robustness of Sample A1's less crystalline particles vs. Sample A2's more crystalline particles (their relative crystallinity being ascertained by HRTEM, Raman, and XRD analysis) would conflict with findings reported in the literature. Researchers have shown that vdW assemblies of small graphenic domains are more fragile—not more robust—than vdW assemblies of larger, more crystalline domains. For example, “amorphous graphene nanocages” that possess a similar morphology to the particles in Sample A1 and comprise assemblies of small, overlapping graphenic domains (often smaller than 10 nm), are easily broken and deformed. Their fragility is explained by the weakness of the vdW interactions between these assemblies' small graphenic domains, which are easily sheared apart. Researchers' side-by-side comparison of amorphous graphene nanocages with more crystalline graphene nanocages constructed from larger domains have demonstrated the superior cohesion of the latter. However, what we actually see is a dramatic improvement in mechanical robustness in every particle throughout Sample A1 compared to the more fragile, nanocrystalline particles found in Sample A2.
Based on this, we can state that the perimorphic framework in
Since all of this networked lattice area is organized in nematically aligned layers, substantially all of this lattice area is subject to interlayer vdW interactions. For the same reason that crystalline graphene nanocages constructed from large-area domains exhibit better vdW cohesion relative to amorphous graphene nanocages constructed from small-area domains, we can infer that as we construct progressively larger anthracitic networks, we can begin to derive a considerable vdW contribution to system cohesion. This is one of the reasons that we find the anthracitic networks more appealing than schwarzite-like graphenic networks (illustrated in
More information about the bonding within the frameworks in Sample A1 can be derived from the sample's Raman spectrum. A single-point Raman spectrum, taken using a 532 nm laser at 2 mW power, is shown in
Another unfitted peak that is apparent in
In order to elucidate the underlying features of the Raman profile in
Next, the OMNIC Peak Resolve software was allowed a third peak, which was manually placed at a starting position of 1500 cm−1 prior to re-running the profile-fitting routine.
Next, the OMNIC Peak Resolve software was allowed a fourth peak, which was manually placed at a starting position of 1150 cm−1 prior to re-running the fitting routine.
Analysis of the four fitted bands indicate a split in the G band (usually found at approximately 1585 cm−1 in unstrained sp2 lattices) into the f-4 peak at 1596 cm−1 and abroad f-3 peak at 1514 cm−1. The f-4 band represents a blue-shifted mode of the G band. The increased frequency of these blue-shifted phonons is caused by compressive strain in some sp2-sp2 bonds. The much broader f-3 peak at 1514 cm−1 coincides with the D″ peak found in graphene oxide and represents a red-shifted mode of the G band. The lower frequency of these red-shifted phonons is caused by the stretching and weakening of sp2-sp2 bonds in ring-disordered regions, as described by Ferrari & Robertson. In addition to inducing tensile strain, the ring disorder of these regions disallows a uniform strain field, which broadens the f-3 band. From the split of the G band into the f-3 and f-4 peaks, we can therefore discern the presence of certain regions of compressed sp2-sp2 bonds, and certain ring-disordered regions of stretched sp2-sp2 bonds.
A blue-shifted band like f-4 is not observed in graphene oxide, in which the G peak, in addition to its normal mode at 1585 cm−1, is also present in the red-shifted mode (called the D″ peak and characterized herein by the trough height). This, in conjunction with Sample A1's lack of oxygen moieties (evidenced by the near-zero rate of mass loss below 400 C in
The f-2 peak in
Interpolation of the VDOS in an alloy structure occurs when there is strong coupling between the phases. Interpolation between the D band (associated with sp2 hybridization) and lower-frequency bands indicates the strong coupling of sp3 states and sp2 states in their immediate proximity. These regions of strong coupling activate the radial breathing mode (“RBM”) phonons found throughout the graphenic system's entire sp2 ring structure. Hence, even a trace-level presence of sp3 carbon states can be discerned in the Raman spectrum due to their activation of RBM phonons that are found throughout the much larger sp2 component. In other words, RBM phonons in grafted singletons are activated by backscattering from the sp3 states in spx rings, where the sp2 and sp3 phases are strongly coupled, and therefore the D band associated with RBM phonons is interpolated. Conversely, the preponderance of sp2 states comprising the sp2 layers between diamondlike seams are neither immediately proximal to the sp3 states, nor strongly coupled to them, and accordingly the G band, associated with sp2-sp2 vibrations, is not interpolated. Based on this analysis, the red-shifted position of the f-2 (i.e. the Df peak) in
What dictates the degree of D band interpolation is not the fraction of sp3 states within the graphenic systems, but instead the fraction of RBM phonons activated by sp3 states vs. the fraction of RBM phonons activated by sp2 edge states. Even a trace level of sp3 states may activate a majority of the RBM phonons if there are even fewer sp2 edge states. This may cause the D band to interpolate, and the degree of interpolation may be expected to increase with an increasing prevalence of sp3 states and decreasing prevalence of sp2 edge states. Of course, the respective prevalence of these two states is negatively correlated, since the spx rings are formed by the conversion of sp2 edges states into sp2 interior states or sp3 states.
Therefore, interpolation of the D band in Sample A1 can be viewed as evidence of the conversion of sp2 edge states into sp3 states associated with diamondlike seams. The conversion of the sp2 edge states into sp3 states associated with diamondlike seams also hints at a tectonic mechanism behind the formation of the seams, and this causal mechanism is explored further in connection with Sample A3 and the samples pertaining to Study B.
Outside of the f-2 peak position, another possible indication of the presence of sp3 states in the Raman spectrum is the shoulder feature associated with the Du peak. This shoulder, which appears between 1100 cm−1 and 1200 cm−1 in
A review of the anthracite literature shows red-shifted D bands in the optical Raman spectra in some grades of natural anthracite-unfitted D peaks can be occasionally found with positions below 1340 cm−l—while in other less mature or more mature grades the D band appears un-interpolated. In the less mature grades, it may be reasoned that this is because diamondlike seams have not yet been geologically formed. In more mature grades (e.g. meta-anthracites), it may be reasoned that diamondlike seams have been formed and subsequently destabilized, eliminating sp3 states and evolving screw dislocations.
To our knowledge, the basis for the D peak's occasional red-shift has neither been investigated, nor assigned to the diamondlike seams. In optical Raman, the ID
Further characterization of the anthracitic networks in Sample A1 was obtained via XRD analysis. XRD analysis was done for a sample synthesized using a procedure similar to Procedure A1, but from a magnesium carbonate feedstock powder. This feedstock powder was calcined to obtain an MgO powder with template particles indistinguishable from Sample A1's. As such, the XRD results from this carbon were analyzed to understand the crystal structure of anthracitic networks like Sample A1.
Three peaks were fitted in the range of interlayer periodicities. The three fitted peaks are referred to as Peaks I, II, and III, and are labeled in
Peak I has a maximum height at 2θ=20.995°, equivalent to a d-spacing of 4.23 Å. Like Peak II, Peak I is also broad, with a FWHM value of 4.865°. The area under Peak I is 32% of the area under Peak II, making it a significant phase of interlayer spacing. A d-spacing of 4.23 Å is too large to be associated with the interlayer phase in graphitic carbon. This peak may reflect the presence of z-adjacent, curved graphenic regions where the curvature is not in phase. Out-of-phase z-deflections disrupt the uniformity of the interlayer spacing and create expanded spaces between the curved regions. This curvature is consistent with anthracitic networks.
Peak III indicates the presence of a phase of smaller interlayer spacing, as well. With a maximum height at 2θ=30.401°, equivalent to a d-spacing of 2.93 Å, the interlayer spacing represented by Peak III is smaller than any interlayer phase in a graphitic carbon. Like Peaks I and II, Peak III is broad, with a FWHM value of 8.304°. The area under Peak III is 80% of the area under Peak II, making it a nearly equivalent phase of interlayer spacing. D-spacing values in the range of 2.93 Å are not found in graphitic carbons, which typically have a <002> d-spacing value of 3.36 Å and no other d-spacings larger than graphite's <100> d-spacing value of 2.13 Å. Heated compression of glassy carbons causes buckling of sp2 regions, sp2-to-sp3 rehybridization, and the formation of sp2/sp3 alloys with interlayer spacings between 2.8 Å and 3 Å. Sample A1's Peak III, with a d-spacing of 2.93 Å, is consistent with this, further corroborating the presence of sp3 states in Sample A1.
Consistent with Sample A1's blue-shifted mode of the G peak, its XRD profile reflects <100> compression. In the intralayer peak range, a <100> fitted peak is fitted with a maximum height at 2θ=30.401°, equivalent to a d-spacing of 2.09 Å. The peak is broad, indicating a broad range of <100> d-spacing values. A <100> d-spacing of 2.09 Å represents a compressive strain of ˜2% in the xy-plane compared to the 2.13 Å d-spacing of graphite.
The thermal oxidation profile of Sample A1 is shown in
TEM analysis of Sample A2 corroborates the deformed, fragmented appearance of the frameworks in the SEM imagery.
In
More information about the bonding structure of Sample A2 can be derived from its Raman spectra. A single-point Raman spectrum, taken using a 532 nm laser at 2 mW power, is shown in
Compared to Sample A1, Sample A2 has a much lower intensity Tru feature, with an ITr
Compared to Sample A1's average Du peak, Sample A2's average Du peak exhibits a higher intensity, with an average ID
The Gu peak is slightly asymmetrical due to the presence of a shoulder at approximately 1620 cm−1. This originates from an underlying D′ peak at 1620 cm−1, which becomes conspicuous due to Sample A2's high density of sp2 edge states. The prevalence of sp2 edge states is also indicated by the narrow Du peak centered at 1349 cm−1. This D band does not appear to be significantly interpolated with any lower-frequency sp3 bands, indicating that most RBM phonons are being activated by sp2 edge states, not by sp3 states associated with diamondlike seams. The D* peak observed in Sample A1 is also absent or negligible.
Three peaks were fitted in the range of interlayer periodicities. The three fitted peaks are referred to as Peaks I, II, and III, where the ascending numbers correspond to the ascending 2θ values at which the peaks obtain their maximum intensity values. The largest fitted peak, as measured by the area under the peak, is Peak II, which obtains a maximum height at 2θ=25.8319° and a corresponding d-spacing of 3.45 Å. The area under Peak II is set at a value of 100%. The d-spacing value of Peak II is consistent with the <002> d-spacing of turbostratic graphitic carbon, and the peak is considerably sharper than Sample A1's Peak II.
Peak I has a maximum height at 2θ=22.9703°, equivalent to a d-spacing of 3.87 Å—a contraction from the corresponding d-spacing of 4.23 Å in Peak I of Sample A1. The area under Peak I is only 13% of the area under Peak II, making it a significant, but smaller phase, whereas the Peak I phase in Sample A1 was 32% of the area of Peak II. The presence of Peak I may reflect larger z-intervals at edge dislocations, or a reduced but not eliminated presence of non-hexagonal rings. The diminishing presence of large, irregular <002> d-spacings is again consistent with the appearance of Sample A2's more aligned, planar fringe lines, as shown in
Peak III indicates a minor presence of a contracted phase of interlayer spacing. With a maximum height at 2θ=31.2063°, equivalent to a d-spacing of 2.86 Å, the interlayer spacing represented by Peak III is significantly smaller than any interlayer spacing in a graphitic carbon. Peak III is also exceptionally broad, with a FWHM value of 10.33°. The area under Peak III is only 5.1% of the area under Peak II, making it a fairly insignificant phase. This is consistent with the scarcity of Y-dislocations observed in Sample A2.
Lastly, the intralayer periodicity at 2θ=42.6906° corresponds to a <100> d-spacing of 2.12 Å, which is close to the graphitic d-spacing of 2.13 Å. This corroborates the lack of compressive strain reflected in the Gu peak's natural position at 1587 cm−1. This may indicate that compressive strain is tied somehow to the formation of crosslinking dislocations and the xy-intervals over which they occur.
The thermal oxidation profile of Sample A2 is shown in
A further practical demonstration of the degraded mechanical properties in Sample A2 vs. Sample A1 was obtain via a uniaxial compression test. In this test, the Sample A1 and Sample A2 powders were each uniaxially compressed to the same pressure. After compression, Sample A1 retained its powder form, suggesting a lack of compaction, while the Sample A2 powder was compacted into a firm, monolithic pellet.
SEM was performed to obtain a better understanding of the powders under compression.
By contrast,
Compared to the perimorphic walls in Samples A1 and A2, which exhibited a consistent appearance, the walls in Sample A3 have regions that are transparent and regions that are opaque. The transparent regions are found within the flat facets of the frameworks and at first glance appear to be holes in the perimorphic wall.
As shown by the arrows in
If no such transparent surface were in fact present to guide the framing, we would expect to see it bent, frayed, or curled irregularly by the mechanical stresses of template removal and drying. These irregularities would not be expected, however, if the framing were supported by a transparent region of the wall stretching across the facet, like a connective tissue. Instead, it would indicate the geometry of the transparent surface, which might be expected to be slightly concave due to the inward pull of the receding water during evaporative drying of the framework, creating a slight concavity. Indeed, this was the appearance of all of the framing. The conclusion from SEM analysis is that the windows observed in Sample A3 are not holes, but a more electron-transparent phase of the wall.
A phase change in the carbon from the edges of a flat facet to the central area of the facet has been observed by previous researchers. When performing CVD growth of perimorphic frameworks on NaCl cubes, a distinct phase of the wall was identified at the edges and corners of the NaCl facets (where nucleation occurred due to localized melting of the NaCl in these areas). Based on Raman analysis, these regions comprised a multilayer vdW assembly of small graphenic domains. A second phase of larger, more crystalline domains within the perimorphic wall was found in the central area of each facet—i.e. the area where there was less melting and nucleation. These perimorphic walls were broken during dissolution of the template and drying, creating platelet-like fragments. The degeneration of these frameworks stands in contrast to the intactness of the perimorphic frameworks in Sample A3, where no observable platelet-like fragments were observed in the dried carbon powder. The observation that the windows in Sample A3 do not break away and become independent platelet-like particles is a compelling indication that the walls in Sample A3 comprise an anthracitic network rather than a vdW assembly.
The perimorphic walls in Sample A3 are somewhat thinner than the walls in Samples A1 and A2. Consistent with this, Sample A3 has a higher BET specific surface area of 328 m2 g−1. This BET measurement suggests an average wall thickness of approximately 8 layers (2630 m2 g−1/328 m2 g−1=8.0). Cross-sections of the perimorphic walls reveal that they are fairly uniform in thickness and do not exhibit any discontinuities, even in the central regions of flat facets. This is shown in
Like Sample A1, Sample A3 exhibits numerous Y-dislocations. A typical fringe pattern drawn from Sample A3 and associated with a Y-dislocation is shown in the magnified inset of
These observations are corroborated by Sample A3's Raman spectra. A single-point Raman spectrum, taken using a 532 nm laser at 2 mW power, is shown in
Also similar to Sample A1, Sample A3 exhibits a relatively sharp, blue-shifted Gu peak (the usual G peak position at 1585 cm−1 is marked with a dashed line in
The ID
From our characterizations of Samples A1, A2, and A3, we can deduce the tectonic pathway by which diamondlike seams are formed during growth. We begin this discussion with the observation that the window regions of the perimorphic wall are electron-transparent, whereas the surrounding framing, and curved regions of the perimorphic wall, are not. We then connect this to an analysis of nucleation and growth of primordial domains over a templating surface. Finally, we model tectonic encounters between these primordial domains, and show how, under the right circumstances, diamondlike seams are evolved from these encounters.
The non-uniformity of electron transparency in Sample A3, as shown in
Next, we recall that, based on the interpolation of Sample A3's Du peak, a significant fraction of Sample A3's RBM phonons are activated by sp3 states, which we have associated with diamondlike seams throughout the anthracitic network. In regions of the wall with a greater density of diamondlike seams, and therefore a greater density of sp3 states, we would expect charging to increase due to discontinuities in the π cloud, through which conduction occurs. In regions of the wall with a lesser density of diamondlike seams, and therefore a lesser density of sp3 states, we would expect less charging should occur. Tying these observations together, it appears that regions of the perimorphic wall associated with higher nucleation density appear to charge more, and we attribute this to a greater density of sp3 states associated with diamondlike seams. We further attribute the greater density of sp3 states and diamondlike seam in these regions to their origin in the grafting that occurs at the tectonic interfaces of primordial domains growing over a common substrate surface. Dense, localized nucleation causes the primordial domains to proliferate, leading to increased tectonic interactions, more grafting, and therefore more sp3 states and diamondlike seams.
Next, we analyze the tectonic encounters between these primordial domains. Ring-disordered lattices possess nonzero Gaussian curvature, and their edges have an undulating geometry determined by the local lattice curvature. The ring disorder of primordial domains grown via pyrolysis at temperatures below 900° C. has been evidenced by several examples in the prior art, including the growth of ring-disordered domains on single-crystal MgO <100> wafers and single-crystal germanium <100> wafers. When two such primordial domains are grown over a common substrate surface, a tectonic encounter may occur between their edges. Since the domains' local lattice curvatures and undulating edges are not in phase, this tectonic encounter creates a stochastic, incoherent tectonic interface between the nearby edge segments. Adding to this complexity, the edges of the primordial domains can be conceptualized as a constantly self-rearranging fluid of free radicals. The incoherence of the interface, where the edge atoms of one primordial domain are not consistently above, below, or level with the edge atoms of the other domain, prevents resolution via simple subduction or sp2 grafting.
In
In the illustration of
The E1-E2 tectonic interface in
The vertical offset within an offset zone is such that opposing edge atoms cannot form sp2-sp2 bonds to their counterparts without severe lattice distortion subduction. Subduction of one edge by the other is also unfavorable. In an offset zone, under the right pyrolytic conditions, edge atoms may undergo sp2-to-sp3 rehybridization and form a sp3-sp3 bond line, grafting the primordial domains together is edge-to-edge. The formation of sp3 states to form bonds in offset zones is herein described as “sp3 grafting.”
In a level zone, the vertical offset between the two edges is small enough and the 2pz orbitals of opposing sp2 edge atoms are sufficiently aligned to allow π bonds to be formed between the edge atoms. In these zones, under the right pyrolytic conditions, the edge atoms may form a line of sp2-sp2 bonds to one another. This is similar to the sp2 grafting that has been observed between ring-ordered domains in the prior art, except that sp2 grafting at incoherent interfaces is localized at level zones.
In the illustration of
In the illustration of
The 6 rings formed via sp2 grafting and sp3 grafting are labeled in
Due to the chiral geometry imposed by their chiral chains, the spx rings R2-C and R4-C represent chiral rings. Both of these chiral rings in
In
The inversion of the edge elevations between the two offset zones also imposes the same chirality on the chiral rings R2-C and R4-C formed at the zone transitions to either side of the level zone. If the edge elevations had not been inverted between Offset Zone I and Offset Zone II, R2-C and R4-C would have had opposite chirality. This alternative scenario is illustrated in Frame II of
Following sp3 grafting within the offset zones, the sp3 atoms in
The graphenic structure G3 shown in
In the illustration of
In the illustration in
The addition reactions also result in the formation of 3 additional 6-member spx rings (designated as R7, R8, and R9 and labeled in
Like the spx rings R1, R5, and R6 located below them, the spx rings R7, R8, and R9 are in the chair conformation, and each has an orientation representing a point-reflection of the spx ring below it. Together, the z-adjacent spx rings R1 and R7 comprise a first diamondlike seam, and the other 4 spx rings (R5, R6, R8, and R9) comprise a second, distinct diamondlike seam, with the 2 diamondlike seams (isolated in the magnified inset of the H1 perspective of
In the illustration of
Located directly above and ring-adjacent to the 3 spx rings R7, R8, and R9 in
These 2 z-adjacent chiral rings are connected via a z-directional chain of sp3-sp3 bonds (comprising the sp3 member atoms labeled 1, 6, 7, and 12). Together, the chiral rings and the z-directional chain of sp3-sp3 bonds comprise a chiral column. Chiral columns, like chiral rings, are found at the inner termini of diamondlike seams in anthracitic networks. The basic architecture of a chiral column may be elucidated by comparing the magnified inset of the H2 perspective in
In the illustration of
The continued growth reflected in
The multilayer graphenic system illustrated in
In the illustration of
We can see in
The spx network illustrated in
Since these atoms are not members of rings, they cannot be members of a graphenic structure or a graphenic system.
The pyrolytic growth sequence modeled in
In Study A, we observe that the Raman D band's interpolation increases as the temperature at which pyrolysis occurs is reduced. This is consistent with the slower release of hydrogen at lower temperatures, which gives the dynamic, self-rearranging condensate at tectonic interfaces more time to relax into an energy-minimizing configuration. Sp2 or sp3 grafting, which eliminates high-energy sp2 edge states at the tectonic interfaces, is therefore promoted by lower temperatures.
In Procedure A1, the 750° C. CVD temperature allows gradual dehydrogenation and carbonization of the condensates. This facilitates some sp2 and sp3 grafting at tectonic interfaces, and as sp2 edge states are eliminated via grafting, the D band begins to show underlying, interpolated modes, as evidenced by difference between its average Du peak, which is positioned above 1345 cm−1, and its average Df peak, which positioned at 1343 cm−1. On this basis, we classify the perimorphic frameworks in Sample A1 as minimally grafted z-spx networks.
In Procedure A2, the 1050° C. CVD temperature accelerates dehydrogenation and carbonization of the condensates. High-energy edge dislocations get locked in, creating a vdW assembly. RBM phonons are activated by these sp2 edge states, and the D band of Sample A2 is therefore not interpolated. On this basis, we classify the perimorphic frameworks in Sample A1 as vdW assemblies.
In Procedure A3, a further reduction in temperature to 650° C. allows the growing condensates more time to rearrange and relax into energy-minimizing, grafted configurations that eliminate sp2 edge states. Consequently, Sample A3's Du peak, positioned at 1340 cm−1 reflects the most D band interpolation of any of the samples in Study A, and is located between the sp2 edge-activated D band at ˜1350 cm−1 and the cubic diamond peak at 1332 cm−1. On this basis, we classify the perimorphic frameworks in Sample A3 as partially grafted z-spx networks.
The samples produced and evaluated in Study B comprise perimorphic frameworks synthesized via surface replication on mesoporous or macroporous MgO templates. These samples, like Samples A1 and A3, exhibit superior mechanical properties and comprise anthracitic networks.
In Study B, lower pyrolysis temperatures were explored to demonstrate the effects of slower dehydrogenation of the free radical condensates, which it was theorized might facilitate the condensates' ability to relax into energy-minimizing grafting configurations at tectonic interfaces. Based on Study A, it was expected that this would lead to fewer sp2 edge states, which could be discerned spectroscopically via progressive interpolation of the D band. The temperature setting of the CVD furnace was varied between 640° C. and 540° C.
Evaluation of the Raman spectra of Samples B1-B3 indicates a downward tendency of the Du peak intensity (as well as the peak area) as the pyrolysis temperature is decreased. The peak FWHM does not appear drastically changed. This trend of reducing peak intensity and area signifies an overall reduction in the RBM phonons associated with sp2 rings. This is known to occur as sp3 content increases in disordered carbons—in diamondlike carbons with no sp2 rings, the D feature disappears entirely. The decreasing Du peak intensities observed in Study B can therefore be assigned to a progressive decrease in the presence of sp2 rings, which are transformed into spx rings by the sp2-to-sp3 rehybridization associated with sp3 grafting. As the pyrolysis temperature is reduced, not only do condensates have more time to relax into lower-energy sp3-grafted configurations at tectonic interfaces, but the primordial domains' ring disorder is increased, which should promote offset zones at the expense of level zones. Both of these should increase sp3 grafting and spx rings.
Evaluation of Samples B1-B3 also shows that as the CVD temperature is reduced in Study B, the Du peak also becomes progressively more interpolated with lower-frequency sp3 bands. This indicates a decreasing prevalence of sp2 edge states. As discussed in Study A, this establishes that sp2 edges are increasingly being eliminated at tectonic interfaces, consistent with the adoption of lower-energy, grafted configurations. Interestingly, the interpolation trend observed in Samples B1-B3 does not stop at the cubic diamond peak position of 1332 cm−1 but progresses to even lower frequencies.
Surprisingly, as temperature drops and grafting is promoted, it also appears that the overall level of lattice distortion in sp2 clusters is reduced. This is evidenced by the trend in the trough height for Samples B1-B3—a trend that was not observed in Study A, where it was found that Samples A1 and A3, while being synthesized at lower temperatures than Sample A2, exhibited higher troughs. This trend in Study B can potentially be explained by compression arising from the increasing prevalence of sp3 grafting and, in particular, from the increasing prevalence of more strained spx ring conformations, such as boat conformations.
Another trend observed in of Samples B1-B3 is that with decreasing pyrolysis temperatures, the Gu peak position gradually blue-shifts from its usual position at 1585 cm−1 up to 1596.6 cm−1. This indicates an overall increase in the compressive strain of sp2-sp2 bonds, and this compression is also attributed to increasing grafting. Additionally, the G band becomes narrower, indicating less variance in the strain states. Hence, Study B corroborates the correlation observed in Study A of grafting and compression. This compression also helps to explain the declining height of the trough. We can see in
Another spectral observation in Study B is that the progressive interpolation of the Du peak position to below 1328.6 cm−1 (in Sample B3) under 532 nm excitation. Because of the proximity of Sample B3's Du peak position of 1328.6 cm−1 to the cubic diamond peak position at 1332 cm−1, and because anthracitic networks are known to be prone to beam-induced heating, which could affect the Du peak position, Sample B4 was evaluated at a lower laser power setting of 0.5 mW. The Raman spectrum gathered for Sample B4 at the 0.5 mW laser power setting demonstrates conclusively that the D band is red-shifted below the 1332 cm−1 cubic diamond peak position. This interpolation below 1332 nm−1 indicates the presence of spx rings in hexagonal diamond arrangements. Hexagonal diamond has been shown to have an intense Raman peak at 1324.4 cm−1 by some workers, whereas in other instances it has been shown to have peaks between 1318 cm−1 and 1325 cm−1. Hence, Sample B4's average Du peak position of 1324.5 cm−1, and multiple point spectra with Du peak positions between 1318 cm−1 and 1320 cm−1, is strong evidence of spx rings in non-chair conformations.
In addition to its greater degree of interpolation, the Du band in Sample B4 is also conspicuously narrower than the Du bands in Samples B1-B3. This indicates that a higher fraction of its RBM phonons is being activated by backscattering at spx interfaces, and that RBM phonons activated by backscattering at sp2 edge states are being eliminated. The more these sp2 edge atoms are eliminated, and the more highly grafted the spx network becomes, the narrower this peak should become. This improvement in grafting in Sample B4 may be attributed to three factors: (i) the increased stability at lower pyrolysis temperatures of strained spx conformations required for grafting across certain tectonic interfaces; (ii) slower dehydrogenation at lower pyrolysis temperatures, allowing condensates more time to finding grafting configurations; and (iii) the use of smaller, less sterically hindered C2H2 gas molecules.
We start with the first factor, which is premised upon the idea that certain tectonic interfaces may not allow chair conformations, i.e. cubic diamond. This premise would be consistent with previously published graphene-to-diamond bonding research. In this work, it was found that for a graphene domain's edge to bond to a diamond surface, it was necessary for the atomic positions of the graphene's dangling bonds to be matched as closely as possible to the atomic positions of some line of sp3 atoms present on the diamond surface. For certain graphenic edge configurations, lonsdaleite (i.e. hexagonal diamond) surfaces offered a better-matching line of sp3 atoms than cubic diamond surfaces.
In our discussion of
In a hypothetical zigzag-zigzag interface in which the edges are sufficiently close to bond directly, such as the E1-E2 interface presented in
Since the spacing between participating edge atoms in a tectonic interface is stochastic in nature, though, we must consider that in some interfaces, opposing edge atoms may be too far apart to bond directly to each other. To illustrate this, in Frame I of
In Frame II of
In Frame III, the opposing line of sp2 edge atoms in E* undergoes sp2-to-sp3 rehybridization, forming a line of sp3 atoms, and these are bonded to the line of interstitial atoms via sp3-sp3 bonds. This line of sp3-sp3 bonds ring-connects the graphenic structures. The elevated sp3 radicals on the E** side allow continued radical addition, resulting in the formation of spx rings in the boat conformation (since chair conformations are geometrically disallowed). With continued growth, a seam may be evolved, as shown in Frame IV of
Hence, the lateral spacing at tectonic interfaces play an important role in determining the conformations of the spx rings evolved by sp3 grafting. If the spacing between zigzag edges is close enough, opposing sp2 edge atoms may be able to rehybridize and sp3-graft directly to each other, resulting in spx rings in chair conformations. If the spacing between zigzag edges is too far, an interstitial line of atoms may be inserted, and sp2 edge atoms may be rehybridized, forming two lines of sp3 atoms that can then form a sp3-sp3 bond line. This will result in less thermodynamically stable conformations that may not be stable at higher temperatures, meaning that complete grafting of tectonic interfaces may not be possible at higher temperatures. We may confidently conclude that, based on the inevitability of these interfacial configurations and their necessitation of spx rings in boat conformations, if an spx network does not exhibit D peak interpolation with sp3 modes below 1332 cm−1, it is incompletely grafted.
The insertion of interstitial atoms, as modeled in
The logic of tight atomic “packing” at tectonic interfaces applies not only to offset zones, where sp3 grafting occurs, but also to level zones, where sp2 grafting occurs. The insertion of interstitial atoms at tectonic interfaces explains the progressively higher G peak positions observed in Study B, with Sample B4 reaching an average position of 1603.3 cm−1 and point positions of 1604.2 cm−1. In procedures utilizing C2H2 feed gas at pyrolysis temperatures below 580° C., we have observed average Gu peak positions of greater than 1606 cm−1, with point positions of up to 1610 cm−1.
Other stochastically-formed tectonic interfaces may easily be envisioned, and sp3 grafting at these interfaces may evolve other spx ring morphologies. These may include 5-member rings, 7-member rings, 9-member rings, and potentially others, all of which ring-connect the participating graphenic structure. Any sp3 grafting event that evolves these spx rings may, upon further addition, form a diamondlike seam.
As an example of this, in Frame I of
Sp3 grafting therefore proceeds via sp2-to-sp3 rehybridization of these opposing sp2 edge atoms, forming two lines of sp3 atoms with atomic positions that allow the formation of a sp3-sp3 bond line between the two graphenic structures. This is illustrated in Frame II of
As shown in Frame III of
If the spacing of a zigzag-armchair interface is too large for bond formation between opposing edge atoms, interstitial atoms may need to be inserted. In such cases, sp3 grafting may lead to the formation of boat and half-chair conformations-just as it does in zigzag-zigzag interfaces with interstitial atoms. In Frame I of
In Frame II of
As shown in Frame III of
The stochastic nature of the processes makes it inevitable that there will be a variety of tectonic interfacial configurations, spx rings, and diamondlike seams, but the exemplary models detailed herein suffice to illustrate the governing principles underlying these varied, specific scenarios. They also explain the observation of Raman spectral features that are consistent with cubic and hexagonal diamond motifs.
Next, we consider more broadly the tectonic interactions and pyrolytic growth of a larger population of primordial domains, which gives rise to higher-layer tectonic activity that we have not yet considered. To illustrate this, we diagram the formation of an spx network in
In Stage I of
In Stage II of
In Stage III of
Our staged depictions of vertical and lateral growth in
The Gu peak position (as a relative indicator of compressive strain), the Du peak position (as a relative indicator of the elimination of sp2 edge states), and therefore the spectral interval between them (as an indicator of both compressive strain and the elimination of sp2 edge states) may provide a useful metric for characterizing the extent to which different spx networks have been able to form grafting bonds across the various stochastically-formed tectonic interfaces created during growth. This interpeak interval—defined herein as the distance in wavenumbers between the Gu and Du peak positions—is commonly used in the anthracite literature to determine the vitrinite reflectance via the Raman spectrum. The vitrinite reflectance, in turn, is a measure of the maturity of a coal. As coal matures, its interpeak interval expands, corresponding to increasing vitrinite reflectance. For an immature to mature coal, using 532 nm excitation, previous workers have calculated the vitrinite reflectance as: νR0%=0.0537(Gu−Du) −11.21, where νR0% is the vitrinite reflectance (as calculated by Raman parameters).
In Sample B4, the interpeak interval is 278.8 cm−1, corresponding to a vitrinite reflectance of 3.76. This vitrinite reflectance is typical of anthracite. Beyond this value, the interpeak interval saturates at approximately 280 cm−1 (varying a bit with excitation due to dispersion of the D peak), whereupon the interval begins to shrink again as anthracite matures into meta-anthracite and finally graphite. As this maturation happens, the ID
Next, we characterized Sample B4 via XRD analysis.
The most prominent feature of the XRD profile of Sample B4 is its main peak at 2θ=24.489°, which reflects a <002> d-spacing of 3.63 Å. This is significantly larger than the 3.35 Å<002> d-spacing associated with AB-stacked graphite or the 3.45 Å <002> d-spacing associated with turbostratic graphite. We attribute this expansion to forced AA-stacking at a large number of the cubic diamondlike seams distributed throughout the spx network. In AA-stacked regions, Pauli repulsion produced by alignment of the π electron orbitals can be expected to increase the minimum interlayer spacing. Indeed, the interlayer spacing of AA-stacked layers has been predicted to have 3.6-3.7 Å, which is in good agreement with the main interlayer peak at 2θ=18.454°. Additionally, we observe a related, minor <004> peak at 2θ=50.192°, reflecting a d-spacing of 1.82 Å—one-half of the <002> d-spacing of 3.63 Å.
A second interlayer peak is fitted at 2θ=18.454°, reflecting an interlayer d-spacing of 4.80 Å. These values, and the breadth of the peaks, indicate a broad range of large interlayer spacings-larger than we observed in Study A. This is explained as follows. Increased atomic packing as a result of grafting in a highly grafted x-spx network causes in-plane compressive strain that exceeds the critical buckling strain. Regions that are compressed beyond this critical buckling strain are forced to buckle in the positive z-direction, this direction representing their only degree of freedom. For this to occur requires them to overcome their vdW attraction to the underlying layer. If they are sufficiently strained, this occurs, and they bow out from the z-adjacent layer below, reaching a maximum z-deflection amplitude somewhere near the geometric center between the lateral seams anchoring their periphery. This z-deflection relieves these regions' in-plane compressive strain but also increases their interlayer d-spacing. We would expect bowing to create a broad continuum of interlayer d-spacings, and this is exactly what we observe in
With this association established, we can see signs of bowing even in the interlayer d-spacings of Sample A1 (a minimally grafted z-spx network) and Sample A2 (a vdW assembly), and we can see that these samples also exhibit states of in-plane compression based on their <100> peaks, which indicate d-spacings below 2.13 Å. From this, we can that similar phenomena are occurring in these less-grafted systems. In Sample A2, specifically, it is likely that localized spx networks are being constructed, but these do not extend throughout the whole perimorphic wall. In other words, the spx networks formed within the perimorphic walls in Sample A2 are too poorly grafted to extend the ring-connected network throughout the whole perimorphic wall.
Based on our findings presented in Experiments A and B, it is possible to speculate ex post facto about instances within the prior art where sp2 and sp3 grafting may have occurred in graphitic networks.
In one such instance, Cui employed a template-directed CVD procedure using methane (CH4) and MgO template particles at 950° C., which produced a monolayer graphenic structure that, as synthesized on the template, possessed a Du peak position of 1322 cm−1 (under 633 nm excitation). Barring any interpolation of the D band, under 633 nm excitation we would have expected the Du peak of this graphenic monolayer to be found around 1332 cm−1. As we have discussed, this would be consistent with sp3 grafting and the formation of spx rings in the chair conformation. Therefore, the reported D peak position of 1322 cm−1 reported might represent a red-shift caused by interpolation.
However, we note a few points. First, in order to satisfy ourselves on whether or not Cui's procedure produced an sp3-grafted system, we attempted to replicate the reported results. We were pleased by the close agreement in the BET and TGA characterizations of the replicated sample we were able to synthesize with these characterizations of the sample reported by Cui. Furthermore, our Raman spectral analysis (performed under 532 nm excitation) revealed a very similar Raman spectrum in terms of the ID
Second, irrespective of the of the D band interpolation in the sample reported by Cui, the sample could not be described as an anthracitic network or an spx network insomuch as the graphenic particles generated were natively monolayer, as synthesized on the template, and as such any crosslinking was lateral. The case for this was made convincingly in the prior art based on extensive BET, TGA, and XRD characterization. Hence, the vertical crosslinking between layers afforded by an anthracitic network was not realized, as these dislocations require a native, multilayer structure. It is true that the monolayer network, upon removal of the template, were reported to collapse into a bilayer structure. However, these bilayers would not have been crosslinked by dislocations, sacrificing this important third dimension of molecular-scale crosslinking present in anthracitic networks. The lack of dislocations was apparent in HRTEM imagery of the bilayers, where the fringe lines were uninterrupted, visually distinct and traceable over distances of 10 nm or more.
In another work within the prior art, Chung flame-synthesized carbon nano-onions at measured temperatures of 700° C. or less (the measured temperatures varied based on where measurements were taken). This process involved rapid chemical vapor deposition over metallic catalyst nanoparticles, creating graphitic carbon nano-onions via precipitation. Based on our ex post facto analysis, it appears that these graphitic carbon nano-onions comprised diamondlike seams. However, the mechanisms and patterns of crosslinking would have been different, given the graphitic alignment of the layers comprising the layered network (this graphitic alignment was evident in HRTEM analysis and also established by the reported <002> interlayer d-spacing of 3.45 Å). In particular, there would have been far fewer chiral rings and columns in these graphitic networks, due to the scarcity of zone transitions at tectonic interfaces between their highly ring-ordered domains. These transitions are directly related to the undulating edge geometry associated with ring-disordered domains grown via a free radical condensate growth mechanism. Additionally, these carbon nano-onions offer less versatility and diminished control over important morphological attributes compared to the growth procedures demonstrated herein. Nevertheless, it is foreseeable that certain aspects of this flame-synthesis process, such as partial oxidation, could be employed in tandem with the use of non-metallic catalysts and free radical condensate-based growth.
In exploring other pyrolytic procedures capable of synthesizing spx networks, we found that employing template-directed CVD temperatures similar to those employed in Study B, but at lower temperatures (between 325° C. and 500° C.), produced carbons with increasingly brown coloration. At 400° C. and below, incomplete dehydrogenation of the condensate during growth resulted in carbons possessing a bright brown coloration. At a temperature of 460° C., the carbons produced appeared gray with a faint brown hue.
A comparison of two samples (Samples C1 and C2) synthesized at these temperatures is shown in the drawing of
Raman characterization of Samples C1 and C2 was performed using a 532 nm laser at 0.5 mW power under an Ar blanket. This lower laser power was deemed appropriate due to the thermal instability of the samples at higher power.
Samples C1 and C2 both exhibit a decreased interpeak interval compared to the samples in Study B, which is consistent with more hydrogenation and less grafting. In Sample C1, the Du peak was interpolated, as shown in
As shown in the averaged spectrum of
The coexistence of hydrogenated and dehydrogenated phases may correspond to phases grown inside and outside of the porous template, respectively. Namely, in addition to the increased stability of C—H bonds at lower CVD temperatures, inside the porous template, where gas-exchange is diffusion-limited, we would expect an increased proportion of H2. Unable to carbonize due to the inability to release molecular hydrogen, the free radical condensate in such regions would ultimately relax back into neutral, smaller molecular weight hydrocarbon species. Workers in the field of free radical condensates have shown this phenomenon via time-of-flight mass spectroscopy. To corroborate this, Sample C2 was immersed in ethanol under gentle stirring conditions. This created a stable, amber-colored dispersion that passes through filters, indicating the dissolution of an oily phase of hydrocarbons.
Study D was performed to confirm the role of H2 gas in throttling the release of molecular hydrogen during free radical condensate growth. Procedures D1 and D2 were substantially the same, with the exception that in Procedure D1, only C3H6 and Ar were flowed into the reactor, whereas in Procedure D2, a low flow of H2 was incorporated in addition to the C3H and Ar. It was hypothesized that the presence of H2 should slow down the carbonization process and facilitate the condensate's relaxation into energy-minimizing, grafted configurations at tectonic interfaces. Raman analysis was performed using a 532 nm laser at 5 mW power.
The increased interpolation of the Du peak position in Sample D2 confirms that increasing the presence of H2 promoted the elimination of sp2 edge states in Procedure D. Based on Sample D1's Du peak position of 1341.9 cm−1, the perimorphic frameworks in Sample D1 comprise partially grafted z-spx networks. Based on Sample D2's Du peak position of 1329.5 cm−1, the perimorphic frameworks in Sample D2 comprise highly grafted x-spx networks.
From the approximately 50% reduction in carbon growth, we can also see that by slowing the condensate's carbonization, the rate of carbon growth was slowed. Hence, we find that H2 partial pressure may be used to throttle carbonization and to improve grafting-particularly at higher temperatures where carbonization is hastened. Based on this, we can infer that, in addition to the pyrolysis temperature, the C:H ratio of the carbon source gas, the rate of H2 release and diffusion from growth, the presence of an H2 feedgas, the morphology and pore structure of the substrate, the size of template particles, the activity of the substrate surface, the presence of H2 scavenging species, and numerous other factors are significant insomuch as they will all affect the dynamic equilibrium of the free radical condensate's hydrogenation and dehydrogenation.
Understanding this may allow faster kinetics to be obtained by rationally balancing these many factors. As a simple example, we have observed that we could simultaneously achieve a lower Du peak position (consistent with better elimination of sp2 edge states) and faster carbon growth kinetics when using a 700° C. CVD temperature and a 30 sccm of H2 feedgas compared to when we used a 580° C. CVD temperature without H2 as a feedgas.
Study E was performed to demonstrate the formation of helicoidal x-networks and z-networks from spx networks (in this context referred to as “spx precursors”). Samples E1 and E2 were generated using the same template material and comprised the spx precursors. Samples E1A and E2A were generated by maturing the Sample E1 and E2 spx precursors, respectively. This maturation, or sp3-to-sp2 rehybridization-induced transformation, was obtained by annealing the spx precursors prior to the removal of the MgO endomorphs—i.e. by annealing the C@MgO perimorphic composite.
Equivalent masses of the Sample E1 and E1A are shown side-by-side in
The flexibility of the perimorphic walls in Sample E1 and the surface tension of the water during drying cause the endocellular pores to collapse, so that only the sheet-like superstructure, shown clearly in
We can see in the magnified inset of
A similar comparison was made between Sample E2 and E2A. Like Sample E1, Sample E2 densified into hard, macroscopic granules, like the one shown in
Sample E2A occupied a conspicuously larger volume and was finer in consistency than the Sample E2 powder. Compared to the larger, harder granules in Sample E2, the Sample E2A powder consisted of smaller, softer agglomerates, as shown in
To understand the changes in the bonding structure created by annealing, Raman analysis was performed using a 532 nm laser at 5 mW power.
The interpolated Du peak positions in Samples E1 and E2 indicate the presence of sp3 states associated with diamondlike seams. Based on Sample E1's Du peak position of 1335 cm−1, a perimorphic framework from Sample E1 comprises a partially grafted z-spx network. Based on Sample E2's Du peak position of 1328 cm−1, a perimorphic framework from Sample E2 comprises a highly grafted x-spx network. Their interpeak intervals are typical for anthracite.
By comparison, the Du peak positions of the matured Samples E1A and E2A are 1352 cm−1 and 1347 cm−1, respectively. These fall into the sp2 D band's normal range under 532 nm Raman excitation; as such, maturation has eliminated the strong coupling of sp2 and sp3 phases in the perimorphic frameworks of Samples E1A and E2A. This indicates that the sp3 states associated with diamondlike seams have been substantially reduced or eliminated in Samples E1A and E2A. Their increased ID
Given the elimination of diamondlike seams, which provide a crosslinking mechanism to the spx networks in Samples E1 and E2, it is surprising that the particles and the perimorphic walls in the mature samples are rigidified. If these mature particles were not ring-connected, such thin-walled carbons should not have survived extraction of the templates, much less have been conspicuously rigidified compared to their spx precursors. We can therefore conclude that the mature particles are crosslinked via crosslinking structures that are more rigid than the precursors' atomically thin diamondlike seams.
Aside from the reversion of their Du peaks back to the normal D band range, Samples E1A and E2A also exhibit increased Du and Tru peak intensities (relative to their Gu peak), as shown in
To demonstrate the maturation of the spx precursor into a helicoidal network, we start by modeling the effect of sp3-to-sp2 rehybridization on diamondlike seams. Frame I of
During annealing, as shown in Frame II of
In this way, the diamondlike seams via lateral unzipping, and the associated ring-connections between z-adjacent layers are also eliminated. The singleton from Frame I of
Next, we consider the effects of maturing an spx precursor with chiral rings and chiral columns. Since we already modeled the formation of such a system (
During maturation, sp3-to-sp3 rehybridization of the sp3 sites results in bond scission. The sp3-sp3 bonds are the least stable and are destabilized first. The sp3-sp3 bonds between the two terminal atomic members of each chiral chain are broken. Each such bond represents the terminus of a lateral sp3-sp3 bond line, and its scission destabilizes the rest of the sp3-sp3 bond line. Accordingly, the linear unzipping of sp3-sp3 bond lines (previously illustrated in Frame II of
In the H1 perspective of Frame II of
This retention of lateral and vertical crosslinking is shown in Frame III of
We can see from Frame III of
The edge segment comprising the sp2 helix represents an interesting structure. While it comprises a zigzag edge configuration, it is unique in that every atomic member of the segment is bonded to three nearest-neighbor carbon atoms, whereas in a normal zigzag edge configuration only half of the edge atoms are bonded to three carbon atoms. This unique attribute of a helical zigzag results from the fact that it represents the chain of atoms created by a broken-open polygon, in which the internal angles of the broken-open polygon are all less than 1800, and thus 3 carbon neighbors are allowed at every edge site (as opposed to a normal zigzag edge, which comprises reflex angles that prevent every edge site from being bonded to three carbon atoms). This novel edge configuration may yield novel electromagnetic and thermal properties, which are known to be dependent on edge configuration in graphenic nanoribbons.
To further clarify the process by which an sp2 helix is evolved from an spx helix, we illustrate the transformation diagrammatically in
During maturation, the sp3-sp3 bond within each of the chiral rings is broken, as we previously discussed in connection with Frame II of
Next, we consider the transformation of the two-dimensional graphenic structure surrounding these one-dimensional helices. As we have established, the formation of an sp2 helix is necessarily accompanied by the formation of a graphenic helicoid, within which the sp2 helix represents an edge segment. The diagram in
During maturation, the central spx helix in Frame I of
Upon relaxation, a single, helicoidal graphenic structure is produced, as shown in Frame III of
These diagrams illustrate how maturation of an spx network with diamondlike seams and chiral rings can generate a laterally and vertically ring-connected mature network. To illustrate the principles of this transformation, we utilized a simple spx precursor comprising a single diamondlike seam and a single spx helix. However, reasonably large spx networks might comprise countless seams and chiral rings formed via tectonic interactions and grafting. In many cases, as we showed in
For this reason, it is desirable to model the transformation of a simple, exemplary spx precursor that comprises multiple seams and chiral rings. Since we already modeled the formation of such a system (
In Frame I of
Relaxation of the system illustrated in Frame II of
To better observe the ring-connections between the two helicoids in Frame III of
From these simple models, the spectral data from Study E, and the changes in mechanical behavior observed in Study E, we can conclude that the changes in bonding structure between Samples E1 and E1A, and between Samples E2 and E2A, are driven by sp3-to-sp2 rehybridization, which transforms spx networks into helicoidal networks.
This is further corroborated by XRD analysis. For this analysis, we annealed Sample B4, a powder of x-spx networks, at a temperature of 1,050° C. for 30 minutes under flowing Ar, creating Sample B4A. This matured the x-spx networks into helicoidal x-networks.
Sample B4A's XRD profile contains significant changes. First, the broad peak fitted at 2θ=18.4540 in Sample B4, which accounted for 30.4%, is not fitted in this range in Sample B4A's profile. We attributed this peak in Sample B4 to a phase of expanded interlayer spacing caused by z-directional bowing of graphenic regions due to intralayer compression beyond their critical buckling strain. At the same time, in Sample B4A, we see the emergence of an even broader fitted peak at 2θ=29.489°, corresponding to a d-spacing of 3.03 Å, with a peak area of 33.2%. These spectral changes suggest an overall shift toward smaller interlayer d-spacings, and the peak center at 2θ=29.489° indicates potential interlayer compression.
Additionally, comparing Sample B4 to Sample B4A, we note a shift in the <100> peak from 2θ=43.138° to 2θ=43.396°, respectively, corresponding to a reduction in <100> d-spacing from 2.10 Å to 2.08 Å. We also see an increase in the main <002> peak at 2θ=23.535°, corresponding to an increase in the average interlayer d-spacing from 3.63 Å to 3.78 Å.
These changes are explained by the transformed crosslinking structure. The cross-section of a diamondlike seam in the <100> plane is a line (i.e. one-dimensional), whereas the cross-section of a screw dislocation in the <100> plane is a point (i.e. zero-dimensional). Therefore, the elimination of one-dimensional pins during maturation leaves only zero-dimensional pins coinciding with the endpoints of the eliminated one-dimensional pins. With the diamondlike seams unzipped, the bowed layers are only pinned at points, instead of along entire lines, and they have more freedom to relax.
The lateral relaxation of these bowed regions has the effect of reducing the amplitude of their z-deflections (thereby eliminating Sample B4's broad peak at 2θ=18.454°, which was attributed to bowing), but obtains this by distributing intralayer compressive strain and lattice distortion more globally. This increases the average interlayer d-spacing (the d-spacing associated with the main <002> peak increases from 3.78 Å to 3.63 Å). It also is reflected in the shift of the broad interlayer peak from 2θ=18.454° to 2θ=29.489°. We see increased compressive strain in the <100> peak, the d-spacing of which is reduced by maturation from 2.10 Å to 2.08 Å.
Unlike the other fitted peaks, which are broad and represent low correlations, the peaks at 2θ=21.6600 and 2θ=35.9440 are sharp, suggesting features with high periodicity. The most likely cause for these are interlayer periodicities that are consistently formed at the screw dislocation cores of the helicoids.
Having now explored the formation of spx networks and their maturation into helicoidal networks and having understood the basic features of these anthracitic networks, we now turn to understanding tectonic zone transitions and their effect on mature, helicoidal networks, and we demonstrate how tectonic zone transitions can lead to the formation of structural variants, including spx double helices, sp2 double helices, and double helicoids.
First, we return to the helicoidal network illustrated in
In an alternative scenario, where the edge elevations between Offset Zone I and Offset Zone II are not inverted, the chiral chains in the two base-layer chiral rings possess opposite chirality. In
In Frame II of
If an spx network were subsequently grown over this base, the spx helices would have opposite chirality, and associated with this, less Eshelby twist between z-adjacent layers. If this singleton were then transformed into a helicoidal network via sp3-to-sp2 rehybridization, the screw dislocation loop formed by the two sp2 helices of opposite chirality would be less strained. From initial formation of the base-layer chiral rings to the intermediate formation of an spx network with mixed dislocations, to the ultimate formation of the helicoidal network, chirality is preserved. Anthracite researchers have observed that screw dislocation loops often involve two xy-adjacent screw dislocations with opposite chirality. We find that loops may also involve two nearby screw dislocations with common chirality.
Another potential interfacial configuration is created when the opposing edge segments crisscross without forming a level zone between the two offset zones to either side. This configuration may occur when, in spite of having similar elevations where the crisscrossing occurs, the 2p2 orbitals of opposing sp2 edge atoms are too misaligned for π bonds to form. The point at which the edges crisscross in this way is referred to as a “crossover point.” Edge atoms at a crossover point may form sp3-sp3 bonds in order to eliminate high-energy sp2 edge states, but they cannot form a sp2-sp2 bond line. We find that at these crossover points, sp3 grafting leads to the formation of chiral columns comprising spx double-helices, which upon maturation form sp2 double helices associated with double helicoids.
The pyrolytic synthesis of an spx network over a tectonic interface with a crossover point is illustrated in
In Stage II of
The two sp3-sp3 bond lines form 6 laterally adjacent spx rings, each comprising 6 atomic members. Five of the spx rings (R1, R2, R4, R5, and R6) are in the chair conformation, with the orientation of R1 and R2 comprising a point reflection of the orientation of R4, R5, and R6. As established in the analysis of
In Stages III and IV of
In
The maturation of the spx precursor GIV causes disintegration because its base is not sp2 ring-connected. The GIV base is sp2 ring-disconnected because of the absence of a level zone and sp2 grafting across the E1-E2c interface from which the base was derived. Instead, only sp3 grafting occurred across the E1-E2c interface, so the primordial domains G1 and G2 were only ring-connected by virtue of the spx ring-connections (R1, R2, R3-C R4, R5 and R6) formed from these sp3-sp3 bonds. After its formation, the base layer remains sp2 ring-disconnected while GIV is constructed over it. As a result, during maturation, the base layer of GIV is completely unzipped along this sp3-grafted interface, such that the primordial regions associated with G1 and G2 become once again disconnected at the base. For the system to remain ring-connected, these two primordial regions of the base must be ring-connected via some path of adjacent rings across the higher layers. However, each higher layer, like the base, is completely unzipped, eliminating any such path. The result is that the spx precursor is completely disintegrated into two graphenic structures, Gi and Gii, where the primordial region G1 is within Gi and the primordial region G2 is within Gii.
The unzipping of the sp2 ring-disconnected base in
Like the other spx rings formed via sp3 grafting, the chiral ring R3-C comprises 4 sp3 members and 2 sp2 members. In R3-C, however, the 2 sp3-sp3 bonds are not parallel-instead, they are point-reflected with respect to each other. This point reflection is due to the inversion of edge elevations that happens at the crossover point where R3-C is located. The 6 atomic members of R3-C are labeled 1 through 6 in Frame II of
As with other chiral rings we have modeled, the termini of the chiral chains in the chiral ring R3-C are connected via sp3-sp3 bonds. In the magnified diagram in Frame II of
Next, we consider the effects of unzipping throughout the spx precursor GIV built over this sp2 ring-disconnected base. In Frame I of
In Frame II of
In Frame III of
We illustrate the fundamental link between the interfacial zone transitions and the ultimate connectedness of a matured system in
In
In
Lattice distortion in a helicoidal network is dependent upon distance from an sp2 helix. This is illustrated by comparing the structures in
Having established the phenomena associated with maturation using simple, small-scale conceptual models, we next extrapolate what happens during maturation of an arbitrarily large spx precursor, which may be formed from numerous tectonic interfaces and grafting of numerous primordial domains. Grafting across these stochastic interfaces and subsequent higher-layer growth leads to the formation of complex, arbitrarily large spx networks. Maturation of these spx precursors forms helicoidal networks of comparable size, comprising numerous screw dislocations. The geometry of these mature networks can be intuited as networks of seamlessly conjoined helicoids-similar to a class of parametric surfaces that have been described as “rheotomic surfaces” in the field of architectural design.
A natural question to ask is whether or not a mature, screw-dislocation network comprises a singleton or an assembly—i.e. whether its membership of graphenic structures is singular or plural. This determination may be straightforward if the mature system is derived in silico from a small-scale, hypothetical precursor with a precisely defined molecular structure. However, to make this determination for a larger-scale, macromolecular precursor system would require mapping its exact molecular structure, which we cannot practically accomplish. What we can establish generally—i.e. for any real spx precursor, without having mapped its exact molecular structure—is that its maturation will result in the formation of a helicoidal network comprising either a helicoidal singleton or a helicoidal assembly. We can also establish that each outcome is consistent with our empirical observations in Study E (i.e. observations of generalized, system-level rigidification and strengthening after maturation).
The first possibility is an outcome herein described as a “singleton-to-singleton” maturation. In this type of maturation, a spx network, which comprises a singleton, is matured into a helicoidal singleton. This type of maturation would be consistent with the empirical observations in Study E (i.e. observations of increased system-level rigidity and strength after rehybridization). A singleton-to-singleton transformation is produced from spx precursors constructed upon an sp2 ring-connected base. To illustrate how a singleton-to-singleton maturation might occur in a reasonably large, complex system, we describe a first scenario in which this outcome is favored. We shall refer to this scenario as “Scenario A.”
In Scenario A, we firstly postulate that, during pyrolytic nucleation and growth of an spx precursor, a multitude of tectonic encounters occur between ring-disordered primordial domains, resulting in a multitude of tectonic interfaces. Due to the out-of-phase edge deflections of the ring-disordered primordial domains, the interfaces are incoherent and stochastic in nature. Wherever level zones occur between two primordial domains, sp2 grafting creates sp2 ring-connections between the participating domains, and wherever offset zones or crossover points occur between two primordial domains, sp3 grafting creates spx ring-connections between the participating domains.
In Scenario A, we secondly postulate that all tectonic interfaces include at least one level zone. From this it follows that, after grafting, all of the primordial domains under consideration will be sp2 ring-connected to one another, such that there will exist a path of adjacent sp2 rings connecting every primordial domain to every other primordial domain. Hence, the base itself will be sp2 ring-connected. It also follows that any tectonic interfaces that include an offset zone in addition to the level zone(s) will comprise at least one interfacial zone transition where a chiral ring will be formed. Lastly, it follows that any higher layers grown over the base will also themselves be sp2 ring-connected (by virtue of sp2 grafting across higher-layer interfaces).
In Scenario A, we thirdly postulate that continued vertical and lateral growth over the base layer forms an spx network comprising the base layer and some number of higher layers that are ring-connected to the base via diamondlike seams (formed over sp3-grafted offset zones) and via chiral columns (formed over tectonic zone transitions between sp3-grafted offset zones and sp2-grafted level zones). As we have already established, these chiral columns formed over level-to-offset zone transitions will comprise a single spx helix and will each be positioned at the terminus of a seam.
In instances consistent with Scenario A, we have already observed (
The other possible type of maturation for a spx precursor is a “singleton-to-assembly” maturation. In this type of maturation, the spx precursor, which comprises a singleton, is matured into an assembly of multiple graphenic structures. A singleton-to-assembly maturation is associated with a ring-connected, sp2 ring-disconnected base. To illustrate how a singleton-to-assembly maturation might occur in a reasonably large system, we describe a second scenario in which this outcome could theoretically occur. We shall refer to this scenario as “Scenario B.”
In Scenario B, we firstly postulate that, during pyrolytic nucleation and growth of an spx precursor, a multitude of tectonic encounters occur between ring-disordered primordial domains, resulting in a multitude of tectonic interfaces. Due to the out-of-phase edge deflections of the ring-disordered primordial domains, the interfaces are incoherent and stochastic in nature. Wherever level zones occur between two primordial domains, sp2 grafting creates sp2 ring-connections between the participating domains, and wherever offset zones or crossover points occur between two primordial domains, sp3 grafting creates spx ring-connections between the participating domains.
In Scenario B, we secondly postulate that none of the tectonic interfaces pertaining to some subset of primordial domains include a level zone. Instead, their tectonic interfaces include only offset zones and crossover points formed via the stochastic crisscrossing of the participating edges. During grafting, these primordial domains are only able to undergo sp3 grafting due to the total absence of level zones in their tectonic interfaces. It follows that only spx rings are formed at their interfaces and that this subset of domains is therefore sp2 ring-disconnected with respect to the surrounding base, of which they are part. It also follows that the base itself is sp2 ring-disconnected.
In Scenario B, we thirdly postulate that continued vertical and lateral growth over the base layer forms an spx network comprising the base layer and some number of higher layers that are ring-connected to the base via diamondlike seams (formed over sp3-grafted offset zones) and via chiral columns (formed over crossover points). As we have already established, these chiral columns formed over crossover points will each contain an spx double helix and will each be positioned at the terminus of a seam.
In a scenario like Scenario B, we have already observed (
Therefore, in Scenario B, where an spx network is constructed over an sp2 ring-disconnected base, it is theoretically possible for a singleton-to-assembly maturation to occur. However, for this outcome to be consistent with the empirical observations in Study E (i.e. observations of increased system-level rigidity and strength after rehybridization), the resulting assembly must be able to resist the shear failure observed in a typical vdW assembly. The creation of an assembly of disconnected members seems inconsistent with these observations. However, we can in fact conclude that even in the instance of a singleton-to-assembly maturation, resulting in disintegration, the resulting assembly will be interlocked so that it cannot shear apart.
This conclusion follows from our third postulate in Scenario B—i.e. that the spx network comprises at least one higher layer. So long as an spx network comprises at least one higher layer, even if a singleton-to-assembly maturation occurs, such that disintegration results in double helicoids of distinct graphenic members, the double helicoids will result in a where double helicoids are formed, even if disintegration occurs, the braid-like geometry of the double helicoids will create an open, interlocking chain preventing the individual, disconnected helicoids from being separated.
The dependency on this interlocking mechanism on the presence of higher layers is demonstrated in
In Frame II-F of
For interlocking to occur, at least one higher layer is needed in the spx precursor, such that the double-helicoid formed during maturation is not so truncated. This is illustrated in
While the graphenic structures in an individual double helicoid could theoretically shear apart via differential rotation around their common axis, this rotational mobility is impossible in a network of multiple double-helicoids. Returning to Scenario B, it follows from our postulates that the helicoidal assembly formed via a singleton-to-assembly maturation would comprise a network of many double-helicoids. Even those primordial domains postulated in Scenario B to be sp2 ring-disconnected with respect to the surrounding base would have crossover points distributed along their incoherent tectonic interfaces—a feature that we have established would create double helicoids. These arrays of double helicoids lack the rotational mobility to be sheared apart, making it necessary to break a graphenic structure in order to break the assembly.
Scenarios A and B are not intended to be limiting, but rather to demonstrate the only two theoretically possible outcomes of sp3-to-sp2 rehybridization of an spx precursor—i.e. a singleton-to-singleton maturation or a singleton-to-assembly maturation—and furthermore to demonstrate how, regardless of which outcome might pertain to a given precursor, the mature system evolved might be expected to exhibit increased rigidity and strength. Either outcome is accompanied by the formation of a helicoidal network that cannot fail via shear, but only via breakage of some graphenic region. This is consistent with our observations of the superior mechanical properties of the mature perimorphic frameworks in Samples E1A and E2A compared to the frameworks in Samples E1 and E2.
To conclude our discussion of singleton-to-singleton and singleton-to-assembly maturations, in
In
In the right-hand multigraph of
In
In the right-hand multigraph of
However, while the primordial domain associated with Node 3 is represented as disconnected in the right-hand multigraph of
This concept is illustrated in
Irrespective of whether the helicoidal network formed by maturation comprises a helicoidal singleton or a helicoidal assembly, the network geometry is analytically similar. Helicoidal networks produce very characteristic fringe patterns in HRTEM.
A preferred variant of a helicoidal network is one that averages between 2 and 5 layers.
The various anthracitic networks described in the present disclosure share certain generic attributes as a function of their layered architecture and nematic alignment. First, they provide more interlayer coupling than non-layered architectures, and we expect system cohesion to benefit substantially from t-t interactions. Compared to schwarzite or other non-layered geometries, we intuit that a denser, layered architecture at the nanometer-scale is preferred due to its combination of covalent and non-covalent modes of cohesion. Density reduction may be obtained by coupling this denser, layered architecture with mesoscale, density-reducing pore phases, following hierarchical design principles. Mesoporous and macroporous perimorphic morphologies constructed from helicoidal networks represent a way to obtain controllable density without sacrificing subnanometer-scale interlayer spacing.
Analogous to the hierarchical approach to density reduction, a hierarchical approach to crosslinking density is also appealing. With respect to the perimorphic frameworks represented in
Other benefits may be derived specifically from the helicoidal network geometry. The superelasticity and spring-like nature of graphenic helicoids has been established, with in silico studies showing a single helicoid sustaining tensile deformation of 1500% without fracture. Failure of a helicoidal network would likely initially occur via covalent breakage of network locations, following by a plastic yielding and unravelling. The mesh-like architecture should offer good toughness properties.
Helicoidal networks (and also spx networks) contain numerous edges on the surface that may be easily chemically functionalized—a fundamental requirement in many applications. Both helicoidal networks and spx networks are easily oxidized with mild oxidants (e.g. sodium hypochlorite, hydrogen peroxide) in line with the procedures described in the '580 Application. These surface edges represent the tops of the conjoined and interlocking helicoids. This is illustrated in
Another appealing surface feature of helicoidal networks is the ubiquitous presence of mouths representing entrances into the network's interlayer labyrinth. One such mouth is shown in
These mouths offer ubiquitous access points for infiltration or exfiltration of fluids, as indicated in
In systems where the primordial level zones are longer (perhaps due to less lattice curvature), longer rows of xy-adjacent sp2-sp2 bonds are formed, increasing the number of xy-adjacent sp2 rings between sp3-grafted offset zones. This will increase the average distance between the helicoids, creating a less densely crosslinked helicoidal network. In systems where the primordial level zones are shorter (perhaps due to more lattice curvature and more frequent crisscrossing), shorter rows of xy-adjacent sp2-sp2 bonds are formed, decreasing the number of xy-adjacent sp2 rings between sp3-grafted offset zones. This will reduce the average distance between helicoids, creating a more densely crosslinked helicoidal network.
Helicoidal networks comprise the preferred variant of synthetic anthracitic frameworks. They generally exhibit superior mechanical properties compared to spx networks. The difference is readily observed in applications. For example,
By comparison,
This demonstrates the utility of synthetic anthracitic networks in composite applications. In [Multifunctional Nanocomposites Reinforced w/ Impregnated Cellular Carbons] and [Multifunctional Nanocomposites Reinforced w/ Unimpregnated Cellular Carbons], the use of “cellular carbons” comprising perimorphic frameworks is shown to be advantageous compared to non-perimorphic morphologies. These applications are herein incorporated by reference. We observe in Study E that perimorphic frameworks comprising anthracitic networks may be especially advantageous in these nanocomposites.
It was demonstrated in Experiments A through E that it is possible, via directed pyrolysis reactions, to synthesize arbitrarily large spx and helicoidal networks. However, practical considerations might still restrict the size of the objects that could be made. To fabricate macroscopic anthracitic networks, it would be appealing to be able to fuse smaller, individual anthracitic networks. We now demonstrate how this may be done by creating a macroscopic preform comprising an assembly of distinct spx networks (i.e. an “spx preform”), then maturing the spx preform to ring-connect the distinct spx networks during maturation. In particular, we explore how static, non-native bilayers formed between the surfaces of adjacent spx networks may become ring-connected during maturation, extending and enlarging the anthracitic network.
We begin with two hypothetical spx networks comprising graphenic singletons, designated GA and GB. Each of these spx networks comprises a microscopic spx network, such as those demonstrated in Experiments A through E. We press GA and GB into contact with one another, such that some regions of their outermost surface layers are in static vdW contact.
Next, we postulate an individual non-native bilayer between two spx networks in static vdW contact, GA and GB. This is represented in Frame I of
While in static contact, the spx networks GA and GB are heated and matured, during which the two lines of tertiary sp3 atoms in GB are dehydrogenated and rehybridized, becoming sp2 radicals as the underlying diamondlike seams are unzipped. The geometry of the underlying helicoids pushes GB's sp2 radicals toward GA, as we attempt to illustrate in Frame II of
In this way, an assembly-to-singleton or an assembly-to-assembly maturation occurs, depending on whether the spx precursors disintegrate during maturation. However, in either scenario, a larger helicoidal network is formed that extends across the bilayer contacts of the spx precursors. The non-native bilayers are cinched together by the helicoidal geometry. If this larger helicoidal network comprises a helicoidal assembly, its graphenic member structures are interlocked with one another in braidlike double helicoids.
Sample F1 comprises perimorphic x-spx networks with a sheet-of-cells morphology similar to the samples in Study E. As observed in Study E, these frameworks' combination of flexibility and flatness causes them to dry into hard, macroscopic granules after extraction of the template. These granules are shown in
The BJH of Sample F1 was 0.289 cm3 g−1, and the BET specific surface area measured, also shown in
Sample F2 comprises a pellet shown in
Sample F3 comprises the Sample F2 pellet after being annealed at 1050° C. for 30 minutes. During annealing, the specific porosity and specific surface area is reduced to 0.028 cm3 g−1 and 233 m2 g−1, respectively, as shown in
This indicates that, during maturation, the lines of sp2 ring connections formed between the layers at bilayer contacts not only cinch the non-native bilayers together, but have a zipper-like effect, drawing together surrounding regions of the layers. This zipping effect occurs via the same mechanism at both inter-network and intra-network non-native bilayers. The zipped regions cause bottlenecking of a fraction of the mesopores (i.e. pores over 2 nm) behind micropores (i.e. pores under 2 nm), as shown in the pore distribution in
Sample F4 comprises the Sample F1 granules after a two-step sequence of annealing and then pressing (in that sequence). Unlike Sample F2, Sample F4 did not comprise a pellet-despite having been pressed under the same conditions as Sample F3, the annealed granules would not form a pellet. The BJH specific porosity and BET specific surface area for Sample F4 was 0.249 cm3 g−1 and 473 m2 g−1, respectively, as shown in
Sample F4 did not form a pellet because maturation caused the anthracitic networks to rigidify (as observed in Study E) prior to pressing them together. In other words, the annealed granules that were pressed in Procedure F4 had already matured into macroscopic, equiaxed helicoidal x-networks. The granules were densified and broken during pressing, so Sample F4 had a mixed granular-powdery consistency. However, the rigidified perimorphic walls could not obtain adequate vdW contact and cohesion, so the pressed system was not pelletized like Sample F2. Additionally, they were not collapsed to the same degree during pressing, as evidenced by the retention of the 3 to 4 nm mesopores of Sample F1.
Raman spectra of Samples F1, F2, F3 and F4 averaged over 16 points are shown in
Sample F5 is another example of a flat macroform, comprising a helicoidal network, being constructed from flat microforms. To fabricate Sample 5, non-compact perimorphic frameworks with hollow architectures similar to diagram III shown in
Sample F6 comprises a section of the Sample F5 macroform that was cut out and annealed at 1050° C. for 30 minutes.
A similar test was performed on Sample F6 by soaking a portion of it in isopropyl alcohol. As shown in
This was confirmed via Raman analysis was performed (at 2 mW power).
The presence of large regions with minimal tectonic activity also explains other spectral features. The high ITr
The lack of tectonic activity during the formation of Sample F5 explains why its ITr
The local absence of sp3 states also explains the spectral changes that occur during maturation of Sample F5. Thus far, we have observed that maturation leads to increased lattice distortion and increased trough height. However, in Sample F6, the trough height is considerably reduced compared to Sample F5. This is because of the local absence of screw dislocations in the resulting helicoidal network—in other words, the helicoids are so large that the dominant spectral effect of maturation is the elimination of ring disorder, which reduces lattice distortion and therefore reduces the trough. The combination of the increased ring order and the absence of screw dislocations is also reflected by the emergence of a 2Du peak in the Sample F6 spectra. The emergence of a 2D peak is indicative of longer-range, in-plane sp2 crystallinity. Based on Sample F6's I2D
So far in Study F, we have demonstrated a process for creating macroscopic anthracitic networks. This involves creating a static, macroscopic vdW assembly from distinct, smaller-scale anthracitic networks (i.e. “microforms”) and ring-connecting them to one another via an assembly-to-assembly or assembly-to-singleton maturation. We have demonstrated this process using flat microforms, which we have used to create both flat and equiaxed macroforms. This basic approach of cohering perimorphic microforms to create a macroform is described in the '308 Application, where the macroforms are described as “peritactic macroforms.” Study F therefore demonstrates that a peritactic macroform can comprise a single anthracitic network.
However, these are only exemplary variants of the inventive concept, which can encompass different densification techniques (e.g. mechanical compaction, evaporative drying, etc.) and forming techniques (printing, 3-D printing, molding, extrusion, injection, drawing, spinning, etc.), without limitation. These and other techniques may be used to create a peritactic macroform of any arbitrary size, geometry and aspect ratio, including elongated, flat, and equiaxed shapes. In particular, we foresee the fabrication of continuous helicoidal networks in the form of yarns, ropes, sheets, and coatings. The only requirements are to bring the spx microforms together into a vdW assembly of the desired geometry and to hold the assembly in a substantially static configuration during maturation. Maximum flexibility and contact between the spx microforms are preferred for obtaining maximum interconnectivity in the final macroform. For this reason, natively few-layer spx precursors are preferred.
The inventive concept also includes the use of microforms of different geometries. A large variety of potential microforms are described and envisioned in the '918 and '760 Applications, and these can be utilized to make different peritactic macroforms, as described in the '308 Application. These microforms may include perimorphic frameworks comprising elongated fibers, flat sheets, or equiaxed prisms, as well as more complex, hierarchical geometries (e.g. rosette-like structures). The rosette-like structures may be especially attractive due to their ability to flex and flatten into aligned plates during densification. This list of microform variants is not exhaustive—other variants may be readily envisioned. Microforms may also be used in combinations of different sizes and geometries.
As an example of one such variant, Sample F7, which is shown in
Other perimorphic frameworks that might be used as microforms are detailed in this disclosure and in the '918 and '760 Applications. These microforms, in addition to varying based on their overall particle geometry, may vary based on their compactness—i.e. their mesoscale crosslinking. This can be seen in a comparison of the elongated microforms shown in
Other microform variants may comprise rosette-like spx networks, like the one shown in
Other microforms comprise equiaxed perimorphic frameworks. In one variant, the microforms may comprise hollow spheres. These may be especially useful if a low-density, macroporous anthracitic network is desired. In another variant, the microforms may comprise perimorphic frameworks with a prismatic or polyhedral superstructure, like those shown in
Study G was performed to ascertain whether microwave irradiation could be utilized as a rapid technique for maturing spx precursors. It was hypothesized that a combination of high temperature, short annealing time, and rapid cooling was desired to mature the spx network fully, while preserving a high density of dislocations. A rapid microwave treatment, it was theorized, would offer this combination.
In Test I of Study G, a Cober-Muegge microwave system was utilized to perform a microwave treatment on the G1 carbon sample. The system consisted of a 2.45 GHz magnetron, 3000 W power supply, steel vacuum chamber, and vacuum pump. The vacuum chamber was outfitted with a rotating platform to facilitate uniform sample exposure and a gas inlet/outlet. The rotating platform could be switched on or off A quartz viewing window located near the top of the vacuum chamber allowed video observation of the sample during the microwave treatment. The microwave assembly is shown in
A 101.0 mg quantity of Sample G1 powder was placed in a medium quartz beaker (“A”). A 100.4 mg quantity of another carbon powder was placed in a small quartz beaker (“B”). The powder bed in each beaker was leveled to a uniform thickness. Beakers A and B were both then placed within a large quartz beaker in case the smaller beakers shattered from rapid heating during the microwave treatment. The large beaker was placed in the vacuum chamber in a centrally located position to maximize microwave exposure. The vacuum chamber was then sealed and vacuumed down to approximately 2 torr, at which point the chamber was refilled to ˜710 torr with nitrogen gas. This was repeated two more times to remove any remaining oxygen in the nitrogen atmosphere.
Microwave irradiation was commenced at a power level of 2400 W. This condition was held for 2 minutes and then the magnetron was switched off. The samples were then permitted to cool back down to room temperature prior to opening the vacuum chamber. The mass of the carbon collected from Beaker A was 95.2 mg and the mass collected from Beaker B was 98.5 mg.
During the 2-minute microwave irradiation treatment, the samples were observed via a video feed. This treatment occurred at approximately 1 atm. Within a few seconds of the commencement of the microwave treatment, Sample G1 began to glow red, and within 10 seconds from commencement, the red glow became bright white. This was likely the period over which rehybridization was occurring. From this point, the brightness continued to grow in intensity, with the video camera auto-adjusting its brightness settings several times to accommodate the growing intensity of light.
While temperature data was not gathered for this experiment, similarly intense white light was emitted in other treatments in which carbon sublimation and re-condensation above the sample as soot could be observed by video. This should only happen at temperatures significantly higher than 3,000° C. Some of the mass loss observed in Samples G1 and the other carbon powder can be attributed to vaporization of oxidized carbon sites (some oxidized sites are retained, despite the lack of an oxidation procedure, due to the nucleation of the carbon lattices on the template's oxygen anions) and adsorbed water. The increased mass loss in Sample G1 may be attributable to some sublimation occurring in this sample.
The remarkably intense Joule heating demonstrated by Sample G1 during microwave irradiation indicates the formation of high-density electrical currents in the carbon particles. Study G demonstrates that microwave heating may be utilized for annealing. It also demonstrates that helicoidal networks may be utilized for resistive heating applications.
In Test II of Study G, a new (i.e. not previously subjected to microwave irradiation) portion of the Sample G1 powder was subjected to microwave irradiation under a lower N2 pressure and power level. The microwave system utilized was the same as the one utilized in Test I. As before, the experiment was performed at room temperature. The lower power setting was selected in order to avoid the formation of a sustained plasma inside the vacuum chamber during microwave irradiation. A small mound of 0.103 mg of Sample G1 carbon powder was placed centrally in a quartz boat, which was placed centrally on the platform. The vacuum chamber was then sealed and vacuumed down to approximately 2 torr, at which point the chamber was refilled to ˜710 Torr with nitrogen gas. This was repeated two more times to remove any remaining oxygen in the nitrogen atmosphere. Finally, the chamber was vacuumed down to 32.5 Torr.
Microwave irradiation was commenced at a power level of 450 W. Surprisingly, the G1 carbon powder did not grow visibly hot, as it had in Test I, but instead remained black, exhibiting no signs of heating. Additionally, almost immediately upon commencement of irradiation, the carbon powder was observed to spread, adopting an extremely fine, smoky appearance that slowly filled the quartz boat. Throughout the irradiation, the powder never showed any signs of heating. Upon terminating the irradiation, the particles collapsed back into a pile at the bottom of the boat.
The absence of resistive heating, coupled with the spreading of the particles in a vacuum, may be explained by a strong diamagnetic response consistent with a resistanceless, superconducting state. Without resistance, Joule heating does not occur. The strong diamagnetic response in this superconducting state is a phenomenon known as the Meissner Effect. In a typical demonstration of the Meissner Effect, a permanent magnet is used to levitate a superconducting compound that has been cooled below its critical temperature (Tc). This occurs due to the formation of screening currents formed near the surface of the superconductor in the presence of an applied magnetic field.
In the case of Test II, we conclude that, under reduced pressure and at approximately 300K, Sample G1 enters a superconducting state, wherein microwave-induced supercurrents flow without resistance through the π electron cloud of electronically decoupled, graphenic monolayers. These supercurrents generate an opposing magnetic field, according to Lenz's law, causing the superconducting particles to repel one another and to spread out into a fine smoke. In effect, each particle becomes a superconducting magnet, and each particle repels the particles around it. This repulsion levitates particles and pushes them outward. Upon terminating the microwave irradiation, the particles stabilize back into a pile at the bottom of the boat.
While it is well-known that pyrolytic carbon is strongly diamagnetic, a diamagnetic response of this strength could not be observed at ambient pressure, nor does the diamagnetism of pyrolytic carbons explain the extraordinary lack of resistive heating under slightly reduced gas pressure. These combined phenomena demonstrate the formation of a resistanceless, superconducting state that is dependent upon gas pressure—in other words, dependent upon reduced gas-surface collisions. Test II occurred at approximately 300 K. Hence, Sample G1 comprises a demonstrated room-temperature superconductor, making it potentially the first among a theorized class of superconductors with Te of 300 K or higher.
Without being bound by theory, we propose the following explanation for the observed superconducting state. First, as we have already demonstrated, the diamondlike seams present in spx networks force AA-stacking (and also bowing), increasing the <002> distance and reducing the electronic coupling between z-adjacent graphenic layers. It has been shown that at the atomic two-dimensional limit, correlation effects become more pronounced, and superconductivity may be achieved with far lower carrier density than in bilayers and bulk structures. Electronically decoupling the layers via AA stacking therefore enables a superconducting state with fewer charge carriers. Second, we propose that the sp3 states within Sample G1 may act as dopants that increase carrier density. This concept of doping via sp3 defects has been explored in connection with carbon nanotubes. Third, we propose that gas-surface collisions at ambient pressures lead to out-of-plane phonon perturbations that break the electronically decoupled state of the atomic monolayer superconductor. This is indicative of a phonon-electron coupling mechanism that, while integral to conventional BCS superconductivity, has not heretofore been conclusively determined for high-Tc superconductors. At the atomic two-dimensional limit, we are able to observe the phonon-electron coupling mechanism experimentally. The superconducting state should be enhanced with further suppression of gas-surface collisions achieved at progressively lower pressures. It may also be enhanced with further doping.
In Test III of Study G, the Sample G1 carbon powder was exposed to microwave irradiation at low pressure in order to demonstrate superconductivity. The microwave system utilized was the same as the one utilized in Tests I and II. As before, the experiment was performed at approximately 300 K. A small mound of 0.1027 mg of Sample G1 powder was placed in a quartz boat. The powder was pushed into a small pile located in the center of the boat, as shown Frame 1 of
Microwave irradiation was commenced at a 300 W power setting. Immediately (within 1 second of commencement) the pile of carbon powder began migrating outward, visible in the camera as a slight change in the outline of the pile. This migration was continued for a couple of seconds, whereupon the magnetron was switched off and the pile stopped moving. The G1 carbon powder remained black, exhibiting no signs of heating. The pile after this initial irradiation is shown in Frame 2 of
At this point, irradiation was again commenced—this time at an increased power setting of 750 W. Again, within just 1-2 seconds of microwave exposure, the carbon powder was observed to levitate, this time migrating down the length of the boat as a black, particulate cloud. This migration, which occurred over a period of approximately 10 seconds, is shown in Frames 3 through 5 in
In Test IV of Study G, four commercial carbon powders were exposed to microwave irradiation at higher pressure. The multiwall carbon nanotube variant of the commercial carbon powder was Elicarb MW PR0940 (Thomas Swan) herein referred to as Sample G2. The multilayer graphene nanoplatelet variant was xGnP Grade C-750 (XG Sciences) herein referred to as Sample G3. The conductive carbon black variant was Vulcan XC72R (Cabot) herein referred to as Sample G4. The flake graphite variant was Microfyne (Asbury Carbons) herein referred to as Sample G5.
The microwave system utilized was the same as the one utilized in Tests I, II, and III. As before, the experiment was performed at room temperature. Piles of 101 mg, 101 mg, 101 mg and 130 mg of Samples G2, G3, G4 and G5, respectively, were placed in separate ceramic boats. The powder was pushed into a small pile located in the corner of their respective boats as shown and labeled in
The initial power setting was at 300 W. Upon commencing microwave irradiation at this power setting, Sample G2 grew visibly hot, turning a dull orange, as seen in
The microwave power setting was finally increased to 1500 W. At this power setting, Sample G2 was the hottest, displaying a bright orange-yellow glow as seen in
In Test V, the response of Samples G2 through G5 to microwave irradiation under reduced gas pressure were investigated. The sample arrangement was unchanged—the chamber was simply pumped down to 32 torr. In Test V, Samples G2 through G5 powders were irradiated again but at 32 torr. The traced outline in
Microwave irradiation was commenced at a 300 W power setting. Immediately (within 1 second of commencement), Sample G4 migrated clearly, visible in the camera as a change in the outline of the pile. Minor migration also occurred in Sample G5, although it was barely distinguishable. After a couple of seconds of migration, the magnetron was switched off and all migration stopped. The samples after this initial irradiation are shown in
Test V showed that a strong, pressure-dependent diamagnetic response was also observed in carbon black (Sample G4). This pyrolytic carbon also exhibits large <002> interlayer spacing, with an XRD report in the literature reporting the <002> peak position at 2θ=250, equivalent to an interlayer d-spacing value of 3.56 Å. We suspect that the same dislocation structures that force AA stacking faults in spx networks are adequately present in carbon black to force electronic decoupling, and that this electronic decoupling is again improved by reducing out-of-plane acoustic phonon perturbations.
In Test VI, the response of spx networks to a strong neodymium magnet under low pressure conditions were investigated to demonstrate flux pinning. A mound of powder of Sample G1 was placed on top of a “magnetic base” made from 9 neodymium bar magnets (N52 Grade with dimensions of each bar 60 mm×10 mm×5 mm). The 9 bars were arranged in a 3×3 formation to create the magnetic base. This magnetic base along with the sample was located centrally on the platform within the vacuum chamber of the microwave system. Microwave irradiation was not used in Test VI; the chamber was only used to achieve low pressure. The vacuum chamber was vacuumed down to 10 torr. After maintaining 10 torr with the sample on the magnetic base for 2 minutes the chamber was backfilled with air to gradually bring it up to atmospheric pressure. Once at atmospheric pressure, the chamber was opened, and the sample and magnetic base were taken out. On inclining the magnetic base to allow the sample to be collected it was observed the sample did not move. The magnetic base and powder were oriented vertically as shown in
A TEM micrograph demonstrating a typical perimorphic framework from Sample G1 is shown in
Intralayer compressive strain was also in the Sample G1's red-shifted Gu peak position of 1594 cm−1. Its average Du peak position was 1333 cm−1, with point spectra exhibiting Du peaks as low as 1327 cm−1, indicative of a highly grafted x-spx network with predominately cubic diamondlike seams, from which we can conclude AA stacking. The average Raman spectrum is shown in
Hence, in Study G, we demonstrate ambient superconducting powders comprising pyrolytic carbons with electronically decoupled layers, and we demonstrate that the superconducting state at the atomic monolayer limit is disrupted under ambient conditions by gas-surface collisions. We theorize that the out-of-plane acoustic phonons created by these collisions disrupt the electronic decoupling of the atomic monolayers in these pyrolytic carbons, whereas this decoupling is otherwise obtained by AA stacking faults forced by the diamondlike crosslinks. The same crosslinks pin the layers together and enforce these high-energy stacking faults, which persist where otherwise they might be minimized upon relaxation of the bilayers.
In Study G, an ambient superconducting powder exhibits both diamagnetic and flux-pinning responses to magnetic fields, indicating a Type II superconductivity. Testing at different pressures ranging from 720 to 10 torr indicate a continuum of strengthened superconductivity as gas-surface collisions are reduced and superconducting pathways are lengthened. The persistence of flux-pinning responses upon returning the powders to ambient pressure indicates that the process of evacuation has modified the particles. In Study H, we observe a similar phenomenon, which is temporary and appears related to the persistence of an internally evacuated state in some nearly impermeable regions of the porous particles for some minutes after evacuation. Reduced permeability in some regions inside the particles and granules is to be expected especially in those samples in which template-directed CVD was utilized, the endomorphic templates were extracted, and carbon-catalyzed CVD growth was then performed again on the porous perimorphic frameworks. We expect that this would begin to close many of the framework's internal pores.
Study H was performed to demonstrate that practical, macroscopic ambient superconductors could be made. Guiding Study H was our hypothesis that the size of superconducting grains in pyrolytic carbons was correlated with the size of their sp2 ring-connected regions. In Study G, the sp2 ring-connected graphenic regions of the microscopic particles in Sample G1 were likely on the same size scale as the particles themselves. In other words, the templating surface of a microscopic template being a closed surface, the spx network formed around that templating surface should comprise a ring-connected network with spx layers that would be similarly closed and sp2 ring-connected with respect to themselves.
In Study H, our objective was to generate a macroform approximating a single ring-connected spx network, with each completed spx layer of this network exhibiting sp2 ring-connectedness with respect to itself over macroscopic lengths. Complicating this was the possibility of fracturing the macroscopic spx network after its creation, which would introduce sp2 edge states in the spx layers. Based on concerns that this might happen during template extraction, we did not extract the endomorphic MgO, but simply created the mesoporous perimorphic composite according to Procedure H and then tested it. The endomorphic MgO pellet is shown in
In Test I of Study H, the macroform's initial sheet resistance upon stabilizing the 4-point probe measurement was 157 Ω/sq. The basic setup of the 4-point probe with a sample and a non-conducting pad beneath the sample is shown in
Following this, the door of the chamber was opened, and the 4-point probe was removed from the sample. Upon removal of contact, the multimeter showed an “Overflow” reading. The 4-point probe was then placed back into contact with the sample, and the reading was again 0.22 Ω/sq. Next, the sample was left for 20 to 30 minutes, after which the sheet resistance measured via the 4-point probe had returned to 157 Ω/sq. This indicates a temporal dependence of the sheet resistance. Raman spectral analysis of the sample revealed no changes from prior to the test. The Raman spectrum is shown in
Performing a number of tests like this on different macroforms, we found that the sheet resistance consistently decreased according to the natural logarithm of the pressure. However, we expect that the sheet resistance's dependency was actually on the pump-down time, which was unmeasured. During pump-down of the vacuum chamber, any diffusion constraints on the outgassing of the porous macroform would be expected to create a temporal dependence of the sheet resistance. This temporal dependence was verified other in experiments by pausing the pump-down and observing that sheet resistance continued to fall even with constant or increasing vessel pressure. This is strong evidence that, for a mesoporous pyrolytic carbon or anthracitic network, the room-temperature ability to form a Bose-Einstein condensate is determined by the pressure inside the particles' pores—i.e. the collision frequency of gas molecules with surfaces inside the macroform.
When growing pyrolytic carbons on an MgO template—and especially when growing on a macroscopic template, as we did in Study H—the differential contractions of the perimorphic carbon and endomorphic MgO phases during cooling can lead to mechanical stresses and either nanoscopic or microscopic fracturing of the spx network. Indeed, it is likely that fine fractures from cooling of perimorphic composites synthesized at high temperatures may be what facilitates endomorphic extraction for template-directed CVD processes in general. Performing a second deposition procedure appears to mend any fractures originating from the first cooling. Damaged sites in the spx network with sp2 edge states become the nuclei for new FRC growth and are healed via sp2 and sp3 re-grafting of these regions, or “mending.” Other possible ways to reduce the present of fractures from cooling is to grow a thicker perimorphic phase and to cool the macroscopic perimorphic composite slowly and uniformly.
Utilizing this “mending” technique, other types of pyrolytic carbon particles-most notably carbon black particles, glassy carbons derived from organic precursors, anthracite, coal, activated carbon, or some combination thereof—could similarly be grafted to one another to create spx macroforms. These disordered seeds act as nuclei for FRC growth, which leads to the ring-disordered lattice formation, tectonic encounters and associated grafting structures that have been demonstrated throughout the present disclosure. This mending technique should eliminate sp2 edge states and ring-connect the individual pyrolytic carbon particles or networks, causing them to coalesce. Mending these particles or networks at reduced pressure with no inert carrier gases may minimize any trapped gas left behind in sealed-off pores.
Having established the importance of evacuating any internal gases, and the ability of an internally evacuated sample to form a Bose-Einstein condensate at ambient temperature and pressure, a barrier phase may be applied to the outside of the evacuated macroform in order to prevent reentry of gas molecules. Utilizing an approach like this, ambient superconducting articles of arbitrary macroscopic length, such as filaments, may be fabricated.
Study H corroborated the observations in Study G, wherein particle-scale, ambient superconductivity was achieved. However, in Study H we were able to measure directly the decline in resistance with reducing pressure, directly corroborating the Meissner Effect and flux-pinning observed in the pyrolytic carbons of Study G. Moreover, Study H showed that at room temperature, it is possible for a porous, ambient superconductor to remain superconducting at ambient temperature and pressure conditions, so long as its pores are evacuated. We strongly suspect that the measured resistance of 0.004 Ω/sq and then subsequently 0.22 Ω/sq may not have actually been attributable to the sample as produced but may have instead been related to massive heating of the probe tips, thereby heating the contact region of the sample above its critical temperature. Other signs of heating caused by the probe tips were observed, including melting of the plastic housing (
Further improvements to the material should be readily achieved via techniques known to those skilled in the art. For example, doping the material to increase the charge carrier density should be readily achievable. Using an organic precursor, such as a polymeric binder, to bind the individual graphenic networks to one another, followed by pyrolyzing the binder and “mending” the networks may improve the ring-connectedness of macroforms. Importantly, the fabrication of infinite, sheet-like or filament-like ambient superconducting articles using roll-to-roll techniques should be possible via the basic approach of evacuated and then sealing the articles with a barrier phase, as we have described.
In the '760 Application we demonstrated the formation of perimorphic frameworks comprising graphenic structures such as hexagonal BN and BCxN. HR-TEM analysis of these networks reveals that they comprise anthracitic networks that are cohered via crosslinking dislocations, including Y-dislocations, screw-dislocations, and mixed dislocations. These materials, which are formed in a way analogous to the FRC growth of carbon, undergo the same mechanics of tectonic encounters and grafting, which in turn lead to the same anthracitic networks.
This disclosure explores how two features (i.e. catalytic activity and solid-state stability) that make refractory metal oxide templates desirable for CVD surface replication procedures can be fulfilled by novel templates that are more soluble. Specifically, we demonstrate templates in which a more soluble “templating bulk” phase and a catalytically active, substantially solid-state “templating surface” phase (as both terms are defined in the '918 Application) are combined.
The following detailed description is organized according to the following sections:
“High-solubility,” as defined herein, describes a material that can be either dissolved in deionized water to form a solution with a concentration of 5 g per liter, or reacted with deionized water to form a compound that can be dissolved in deionized water to form a solution with a concentration of 5 g per liter.
A “high-solubility template” is defined herein as a template structure that fulfills the following requirements: (1) the “templating surface” (as defined in the '918 Application and the '760 Application) remains substantially solid-state during vapor deposition growth of a perimorphic wall; (2) the templating surface comprises a catalytic component that enables the growth of a perimorphic wall; and (3) the template material is a high-solubility material.
A “catalytic component” is defined herein as a component of or on the templating surface capable of catalyzing the decomposition of a reactive gas and enabling the conformal growth of a perimorphic wall over the templating surface. Without a catalytic component, no perimorphic wall may be formed over the templating surface during vapor deposition. The catalytic component may be defects (e.g. step sites on metal oxide templating surfaces). A catalytic component may also comprise an adsorbate on the templating surface.
“Oxyanionic templates” are defined herein as templates that are solid-state under CVD conditions and comprise an oxyanionic compound. It is noted that an oxyanionic template may also be referred to as a “basic oxyanionic template” if it contains oxygen anions. Basic oxyanionic templates may result from partial decomposition of an oxyanionic precursor material.
CVD growth using high-solubility templates proceeds in a similar fashion to CVD growth on metal oxide templates. Namely, the perimorphic wall grows via dissociative adsorption of reactive gas molecules at catalytic sites on the templating surface. The perimorphic wall may then grow over the templating surface, substantially encapsulating the endomorphic template and resulting in surface replication, as described in the '918 Application and the '760 Application. Optionally, the perimorphic wall may have a layered architecture comprising two-dimensional lattices. The endomorphic template may then be extracted via dissolution in a liquid. Endomorphic extraction may result in a perimorphic framework with a simple or complex cellular morphology, depending on the geometry of the templating surface.
It is to be understood that, in practice, just like the template variants that have been described in the '918 Application and the '760 Application, high-solubility templates may take on a number of different shapes and sizes, these shapes and sizes often being determined by the process used to create the template material or a template precursor material form which the template is derived. The template's morphology may be retained after endomorphic extraction by the perimorphic framework formed around it, provided the framework is sufficiently rigid and strong. Such a framework is said to have retained its native morphology, as described in the '918 Application and the '760 Application. Alternatively, the framework may be deformed or fragmented after endomorphic extraction, resulting in a non-native morphology. In either case, the template's morphology may be an important consideration.
High-Solubility Template Considerations
A template initiates perimorphic growth over the templating surface via adsorption. From the initial adsorbate, the perimorphic wall may be grown via deposition. Perimorphic growth via deposition may beneficially occur via “autocatalyzed” or free radical condensate growth, wherein the perimorphic material itself reacts with gas-phase adsorbates during deposition, and therefore it is desirable for the template to remain solid-state under conditions in which this mode of perimorphic growth occurs.
Once the perimorphic wall has been formed, the endomorphic template should be able to be extracted. It is beneficial if this can be accomplished using aqueous liquid-phase processing. It can be beneficial for the template to be reacted with water or dissolved in water to form solutions of reasonably high concentrations of the solute. It is further beneficial if these solutions may then be used to form new template structures with acceptable and well-controlled morphological features, as discussed in the '918 Application and the '760 Application.
To meet these many requirements for each of the many applications that might be encountered, it is helpful to have a diverse portfolio of template materials. A new category of solid-state, high-solubility templates are especially useful members of a portfolio due to the ease with which they may be dissolved in water during endomorphic extraction and reconstituted in well-controlled morphologies via precipitation processes.
Mechanisms for Catalyzing Perimorphic Growth
In this section, we explore nucleation and growth mechanisms for perimorphic carbons from hydrocarbon gases on metal oxide templates, a specific type of surface replication for which a literature exists, and use this literature as a basis for understanding surface replication phenomena in the high-solubility templates disclosed herein. It is to be understood that the following discussion is not meant to be exhaustive. Nucleation and growth mechanisms are not comprehensively characterized even in the case of perimorphic carbons grown on metal oxide templates. Several different nucleation mechanisms may occur simultaneously, further complicating analysis. Any nucleation or growth mechanism, whether fully characterized herein or not, should be considered within the scope of the present disclosure.
It has been found that nucleation on metal oxide surfaces occurs at high-energy surface defects (e.g. step-sites). On other templates, nucleation mechanisms are different. For example, nucleation on metallic templates may require carbon to be dissolved into the metal, then precipitated. Nucleation on metal halide templates may require molten surfaces where inelastic collisions with gaseous molecules occur. Surface defects on the solid-state surfaces of NaCl templates, for instance, have been observed to be catalytically inactive. Therefore, it has been demonstrated that the surface defect-catalyzed mechanism is not universal; in fact, other than our own recent work in the '916 Application, where perimorphic growth was accomplished on oxyanionic templates, the surface defect mechanism has only been observed on metal oxides or metalloid oxides.
Thermogravimetric analysis (TGA) was used to analyze the thermal stability and composition of materials. All TGA characterization was performed on a TA Instruments Q600 TGA/DSC. A 90 μL alumina pan was used to hold the sample during TGA analysis. All analytical TGA procedures were performed at 20° C. per min unless otherwise mentioned. Either air or Ar (Ar) was used as the carrier gas during analytical TGA procedures unless otherwise mentioned.
Raman spectroscopy was performed using a ThermoFisher DXR Raman microscope equipped with a 532 nm excitation laser. For each sample analyzed, 16 point spectra were generated using measurements taken over a 4×4 point rectangular grid. The normalized point spectra were then averaged to create an average spectrum, with a rare point spectrum being excluded from the average due to a poor signal at that location. The Raman peak intensity ratios and Raman peak positions reported for each sample all derive from the sample's average spectrum. No profile fitting software was utilized, so the reported peak intensity ratios and peak positions relate to the unfitted peaks pertaining to the overall Raman profile.
The furnace scheme utilized for all experiments was as follows. The furnace used was an MTI rotary tube furnace with a maximum programmable temperature of 1200° C. The furnace had a 60 mm quartz reactor tube with a gas feed inlet. The opposite end of the tube was left open to the air. The furnace was kept level throughout deposition. Experimental materials in powder form were placed in ceramic boats, and the boats were placed in the center of the quartz tube (in the furnace's heating zone). The quartz tube was not rotated during deposition.
Five experiments are described below. For each of these experiments there are unique template precursor materials, template materials, perimorphic composite materials, and perimorphic materials. For exemplary purposes, perimorphic carbons were formed on these templates.
The template precursor materials include potassium carbonate (Experiments 1 and 2), potassium sulfate (Experiment 3), lithium sulfate (Experiment 4), and magnesium sulfate (Experiment 5). To produce these precursor materials, commercially sourced potassium carbonate, potassium sulfate, lithium sulfate, and magnesium sulfate powders were first dissolved in H2O. Either isopropanol or acetone was then added dropwise while stirring to induce precipitation of the solute. The precipitate was filtered and then dried.
Note that the template precursor materials in Experiments 1 and 2 comprise the same compound (K2SO4). These precursor samples differed only with respect to the batch size, with the batch size in Experiment 2 being roughly 5 times larger than the batch size in Experiment 1.
In the next stage of the experiments, each of the template precursor samples was placed in a furnace according to the furnace scheme described in Section II. Each precursor powder was then heated to the CVD temperature under a 1100 sccm flow of Ar, whereupon C3H6 flow was commenced. The powder at this temperature, and under this atmosphere of flowing Ar and C3H6, comprised the template material.
During CVD, perimorphic composite materials were formed by nucleating and growing perimorphic carbon on the templating surfaces. Similar CVD surface replication procedures have been described in the '918 Application. In each case, the CVD temperature was at least 279° C. below the melting point of the template precursor material, and no signs of melting were observed at any time. Therefore, we can conclude that each of the templates was substantially solid-state throughout CVD. Experimental parameters during CVD were according to
As shown in
In each experiment, the powder retrieved from the furnace had nucleated and grown a carbon perimorphic wall. This confirms that nucleation occurred. Since each of the templates was solid-state, any nucleation that occurred on the templating surface was unrelated to melting. Nucleation due to absorption or dissolution of carbon in these non-metallic templates can also be ruled out. We therefore attribute nucleation to surface defects, as has been observed for metal oxide templates.
The exact nature of the surface defect sites in these oxyanionic templates is not well characterized at this point, and the precise degree to which the templates were purely oxyanionic in chemical composition or may have been basic oxyanionic templates due to minor levels of decomposition is not fully characterized. However, given the extremely small mass losses that were recorded for the anhydrous sulfate samples in Experiments 1, 2, and 4, some of which loss can be attributed to adsorbed water, and given additionally that some of these sulfates melt prior to thermally decomposing (e.g. Li2SO4), and lastly given the almost negligible mass contribution of perimorphic carbon (the perimorphic wall comprising only a few graphenic layers), we can conclude that the extent of any decomposition was minor. For instance, K2SO4 thermally decomposes around 750° C. in the presence of a carbon reducing agent. While it possible that some minor decomposition occurred at 580° C., the templates in Experiment 1 and 2 were substantially K2SO4 in terms of chemical composition.
Endomorphic extraction of the templates from the perimorphic composite structures was performed in each experiment by dissolving the template in water, which was accomplished easily in small volumes of water, further corroborating the high-solubility composition of the oxyanionic templates. The resulting perimorphic frameworks were then rinsed to minimize residual ions upon drying. At this stage, an immiscible solvent might also be utilized to separate the perimorphic carbon from the aqueous process liquid in order to reduce or eliminate the need for rinsing, as described in the '918 Application and '760 Application.
SEM analysis was performed to provide a general understanding of the template and perimorphic carbon materials. Specifically, we analyze the template precursor material, perimorphic composite material, and perimorphic carbon material from Experiments 1 and 5. For the sake of brevity, and because the present disclosure focuses on the ability to synthesize the perimorphic material on the template, rather than on each template's specific morphological features (which will vary substantially not only based on the precipitation process, but also based on the thickness of the perimorphic walls), we do not report SEM analysis for all of the samples.
The inset of Frame I of
The second phase, which can be discerned at higher magnification in Frame IV, is residual MgSO4 on the surface of the perimorphic frameworks. The MgSO4 residue charges under the electron beam. The presence of this residue can be explained by the high solute concentration of the aqueous solution created during endomorphic extraction of the high-solubility oxyanionic template, as well as the large amount of retained water in the three-dimensional perimorphic frameworks. Even after rinsing, the framework contained a significant amount of dissolved MgSO4 that left a ubiquitous residue upon drying. This residue was not observed in Experiment 1 since the voluminous pores with in the perimorphic frameworks were collapsed and less water was retained in them.
Rinsing problems and the voluminous liquid waste streams associated with endomorphic extraction were noted previously in the '918 and '760 Applications. To alleviate this, a solvent-solvent separation was demonstrated to displace the detained aqueous solution held within the perimorphic carbons with an immiscible solvent. This separation technique may be especially beneficial when using a high-solubility template like MgSO4 that can form concentrated solutions (up to 35 g dissolved per 100 mL at room temperature) upon endomorphic extraction.
The perimorphic frameworks produced in Experiment 5, while their walls were only a few nanometers in thickness, demonstrated a superior ability to retain their native morphology than the frameworks produced in Experiment 1. This is attributable to the compactness and associated rigidity of the cellular substructures of the frameworks produced in Experiment 5. Namely, surface replication techniques that utilize nonporous templates, or templates with no nanoscopic pores, will result in a less compact architecture than surface replication techniques that utilize templates with finer pore structures. In the case of Experiment 5, the MgSO4 templates had finer internal pore structures due to the escape of crystalline water from the epsomite template precursor particles during heating. This is reflected in the substantial mass loss observed after thermal exposure, as shown in
The spectra in
Thus, these spectra are consistent with each of the perimorphic carbon samples comprising a lattice-engineered carbon, which has been shown to exhibit increased reactivity and to be more easily chemically functionalized. Additionally, this ring disorder can be expected to change the electronic band structure of the sp2 carbon, with sufficient disorder leading to insulating carbon structures.
Data was extracted from the average Raman spectrum of each of the samples for summary reporting. The average spectra were generated as described in Section II. In the case of Sample 3, the spectral data reported was derived from smoothing the average spectrum, which remained quite noisy, and both the smoothed and unsmoothed spectrum are shown for comparison in
The average spectrum for the perimorphic carbon produced in Experiment 3 is relatively noisy, likely because there was a relatively small amount of carbon within the Raman laser's focal point. This could be explained by uncollapsed perimorphic frameworks' lower-density, three-dimensional morphology, in which microscopic pores nearly as large as the focal point are ubiquitous. In addition to this noise, this sample's spectrum exhibits a substantially higher trough and a higher D peak position than the other samples. These features are associated with much fewer sp3 states, and we theorize herein that this is attributable to the large, atomically flat facets characterizing the K2CO3 surfaces. These nearly defect free surfaces may offer fewer nucleation sites. These flat facets can be seen in
Other oxyanionic templates may also be utilized. For example, in the '918 application, the P18-type perimorphic carbons with a hollow-spherical morphology were grown on spray-dried Li2CO3 templates at 580° C. with no signs of melting and minimal mass loss (<1.5%). If the perimorphic walls are grown thinner, crumpled and sheet-like structures can also be formed, and these can be filtered or dried on glass to form lamellar buckypapers.
Top-down and horizontal views of one such buckypaper are shown in Frames I and II of
Using similar CVD surface replication procedures as those described in Experiments 1 through 5, we have also synthesized perimorphic carbons on high-solubility aluminate (NaAlO2, melting point 1650° C.) and metasilicate (NaSiO3, melting point 1088° C.) templates at a temperature of 750° C. from propylene (C3H6) feedgas. Taken together with the other results in Study B this indicates that oxyanionic templates represent a rich source of potential template materials.
This application discloses several numerical ranges in the text and figures. The numerical ranges disclosed support ranges or values within the disclosed numerical ranges, even though a precise range limitation is not stated verbatim in the specification, since this disclosure can be practiced throughout the disclosed numerical ranges.
The above description is presented to enable a person skilled in the art to make and use the disclosure. Various modifications to the embodiments will be readily apparent to those skilled in the art, and the generic principles defined herein may be applied to other embodiment and applications without departing from the spirit and scope of the disclosure. Thus, this disclosure is not intended to be limited to the embodiments shown but is to be accorded the widest scope consistent with the principles and features disclosed herein. Finally, the entire disclosure of the patents and publications referred to in this application is hereby incorporated herein by reference.
The following applications are hereby incorporated by reference in their entirety for all purposes: PCT/US21/49195 (the '49195 Application); U.S. Provisional Patent Application 63/075,918 (the '918 Application); U.S. Provisional Patent No. 63/086,760 (the '760 Application); U.S. Provisional Patent Application 63/121,308 (the '308 Application); U.S. Utility application Ser. No. 16/758,580 (the '580 Application); U.S. Utility application Ser. No. 16/493,473 (the '473 Application); PCT/US17/17537 (the '17537 Application); PCT/US21/37435 (the '37435 Application); U.S. Provisional Patent Application 63/129,154 (the '154 Application) and U.S. Pat. No. 10,717,843 B2 (the '843B2 Patent).
| Number | Date | Country | |
|---|---|---|---|
| 63086760 | Oct 2020 | US |
| Number | Date | Country | |
|---|---|---|---|
| Parent | PCT/US2021/053316 | Oct 2021 | US |
| Child | 18129282 | US |
