JP2021505429A - Manufacture and design of composites with build-up layers - Google Patents

Abstract

In some embodiments, the composite system is a structure having a constructed 3D geometry, wherein the 3D geometry is a monolithic and deterministic structure and a matrix phase, wherein the matrix phase is at least partial. Is provided with a matrix that impregnates the structure. In some embodiments, the 3D geometry is a nano or micro-constructed 3D geometry.

Description

[Cross-reference of related applications]
This application is the benefit and priority of U.S. Provisional Patent Application No. 62 / 593,768 filed December 1, 2017 and U.S. Patent Application No. 16 / 151,186 filed October 3, 2018. All of them are incorporated as references to the extent consistent with this specification.

The present invention was carried out with government support under license numbers N00014-16-1-2431 and license numbers N00014-16-1-2827 granted by the Navy Research Bureau. The government has certain rights to the present invention.

The present invention relates to composite material systems including structures with constructed monolithic three-dimensional geometry and matrix phases, and methods of making such composite material systems. For example, the 3D geometry can include a continuous and interconnected network of features having any variety of shapes or configurations, including curved and surface features, the matrix phase being at least partial to the structure. Can be impregnated with.

Composites are composed of one or more materials that form a disordered or ordered phase and combine to provide properties or behaviors that may differ from the properties of the individual materials. Usually, these phases are classified as a reinforced phase with high stiffness or strength properties, or a matrix phase filled with the remaining volume and with poor stiffness or strength properties. Composites can be broadly characterized by the volume ratio between the reinforced phase and the matrix phase, the geometry of the phase, and the constitutive properties of each phase (Reference 1).

Composites with disordered phases can include short fibers, particles, or randomly arranged components that do not provide a continuous reinforcing phase. In these materials, the only continuous phase is the matrix, which impregnates the remaining volume around the reinforcing particles. The matrix phase serves to distribute the load between the harder / stronger strengthening particles. Alternatively, the ordered phase composite can also include aligned long fiber (discontinuous phase) or truss-like (continuous phase) geometry as the reinforcing phase and is surrounded by a continuous matrix phase. Both ordered and disordered composites can have a mechanical response that varies from isotropic (ie, probing direction-independent behavior) to anisotropic (ie, material orientation-dependent behavior).

Carbon fiber composites are a form of ordered phase composites that have been applied in aerospace, automobiles, ships, armor, and even sporting goods applications. They are widely used because of their stiffness-to-density ratio, fatigue properties, and resistance to extreme environments (Reference 2). The basic constituents of carbon fiber composites are the fibers that make up the non-continuous order-enhancing phase, with the discrete fibers bonded only through the matrix phase. The arrangement of fibers impregnated by the matrix phase forms the lamina, a two-dimensional subcomponent of the carbon composite component.

The carbon phase discontinuity and anisotropy in carbon fiber composites is associated with several fracture modes. Lamina of woven fibers with increased contact between individual fibers was made to reduce anisotropy, but discontinuous reinforcing phases are still formed (Reference 3). Isotropic properties can be achieved by laminating with unidirectional fiber layers, i.e. forming aggregates of lamina of various fiber orientations bonded via a matrix phase. Reorienting the fibers of each laminar determines the degree of anisotropy of the laminate and the (usually undesirable) stretch-bend bond. Depending on its structure, the carbon fiber laminate may be damaged by various mechanisms such as fiber buckling (during compression), interfiber breakage, and interlayer breakage (Reference 4). Fiber buckling is noticeable during lamina compressive loading due to the high fineness of the fibers in the matrix, which would otherwise not be supported in-plane. Interfiber and interlaminar breakage occurs due to the discontinuous nature of the reinforcing phase and relies on the matrix phase to distribute the load in both the thickness and in-plane directions.

Composite materials with continuous strengthening and matrix phases have been reported, most of which have foam (foaming material) as the strengthening phase (References 5-9). Form is a typical stochastic structure, which is in direct contrast to the deterministic structure. Attempts to correct the discontinuous properties of carbon fibers include adding a phase composed of particles such as a carbon matrix to promote adhesion between the fiber matrix and the interlayer interface (References 10 and 11). ) This approach can increase the strength of these materials, but does not form a continuous carbon phase, resulting in the destruction of the layers (the weakest interface). Other studies report an ordered truss network as a strengthening phase using a polymer waveguide process. The geometry of these truss-like networks is limited to a linear ray pattern starting at the edge of the material and a finite angular range between the connecting truss elements (12). The functional grades of composites by these waveguide processes are subject to the same constraints but reduce the structural integrity of the strengthening phase (References 13 and 14). Attenuating materials using these waveguide patterns have been reported, which require one or more viscoelastic phases as an energy dissipating coating (References 15 and 16).

In some cases, the production of 3D carbon structures, not limited to the architecture defined by the polymer waveguide, has been demonstrated using laminated molding (AM) of polymer samples. Glassy carbon nanostructures that are close to theoretical strength have been manufactured using two-photon lithography (TPL), but this process has extremely low throughput and limited practical applications (Reference 17). ). Cellular carbon microstructures have been made using stereolithography (SL) (18) and graphene airgel printing (19). However, each of these lithograph structures fails to directly manufacture composite parts with adjustable net shape, functional tilt, damping properties, collision absorption properties, or these or other properties defined by the 3D carbon network. It is exposed to such drawbacks.

The composite materials systems and methods of making the systems provided herein address these and other challenges associated with conventional composite materials, thereby being highly adjustable for a wide range of applications and parameter space requirements. System and method.

The present specification provides composite material systems and methods for making composite material systems useful in a wide range of applications, addressing the challenges and limitations of conventional systems and methods. Applications where these systems are useful include aerospace (such as landing gear collision absorption), automobiles (such as brake assembly vibration relaxation), pharmaceuticals (such as medical devices that require specific mechanical properties), and military (eg, such as medical equipment that requires specific mechanical properties). Body armor), maritime equipment, sports equipment, etc., but not limited to these. For example, the systems disclosed herein are useful for reducing collision energy and / or damping vibrations. The disclosed systems and methods are highly adjustable and deterministic, and they can be adapted and tuned according to the desired application and set of properties. For example, the composite material system disclosed herein may include functional tilts through a plurality of continuously interconnected three-dimensional geometries. For example, a composite system may include highly adjustable and deterministic resonator features that provide precise control and adjustability of the damping behavior of the composite system.

In some embodiments, the composite material system includes: That is, a structure having constructed three-dimensional geometry, said three-dimensional geometry being monolithic and deterministic; a matrix phase, where the matrix phase at least partially impregnates the structure. In some embodiments, the 3D geometry is a nano or micro-constructed 3D geometry. In some embodiments, the structure is at least 1%, at least 2%, at least 5%, at least 10%, at least 15%, at least 20%, at least 30%, at least 50%, at least by volume, depending on the matrix layer. 75, at least 90%, or preferably in some applications, is impregnated substantially 100% (% by volume). In some embodiments, the structure is at least 1%, at least 2%, at least 5%, at least 10%, at least 15%, at least 20%, at least 30%, at least 50%, at least by mass, depending on the matrix layer. 75, at least 90%, or preferably in some applications, is impregnated substantially 100% (% by weight).

The composite material systems, structures, and / or 3D geometries disclosed herein are a variety of physics, including collision absorption and / or damping behavior that are not available or difficult to obtain with conventional material systems. It can have physical and mechanical properties or embodiments.

In some embodiments of the systems and methods disclosed herein, the structure, composite material system, or both are selected from the range 2 × 10 ⁴ J / m ² to 4 × 10 ⁵ J / m ^2. It is characterized by the area-normalized collision energy relaxation metric (Ψ). In some embodiments of the systems and methods disclosed herein, the structure, composite material system, or both, is at least 2 × 10 ⁴ J / m ² , preferably at least 2.6 × 10 ⁵ J. It is characterized by an area-normalized collision energy relaxation metric (Ψ) of / m ² or preferably at least 3.2 × 10 ⁵ J / m ² in some applications. In some embodiments of the systems and methods disclosed herein, the structure, the composite system, or both, selected from the range of ^{1.9 × 10 6 J / kg~4 ×} 10 6 J / kg Characterized by the density-normalized collision energy relaxation metric (Ψ). The systems and methods disclosed herein, the structure, the composite system, or in some embodiments of both, from a range of ^{1.9 × 10 6 J / kg~3.2 ×} 10 6 J / kg , Preferably from the range of 1 × 10 ⁶ J / kg to 5 × 10 ⁶ J / kg for some applications, preferably 1 × 10 ⁶ J / kg to 1 × 10 ⁷ J / for some applications. It is characterized by a density-normalized collision energy mitigation metric (Ψ) selected from the range of kg, preferably from the range of 3 × 10 ⁶ J / kg to 1 × 10 ⁷ J / kg. In some embodiments of the systems and methods disclosed herein, the structure, composite system, or both are characterized by reduction of collision energy with energies selected from the range of 1J to at least 900J. ..

In some embodiments of the systems and methods disclosed herein, the structure, composite system, or both are characterized by a bounce coefficient selected from the range 0.8-0.7. In some embodiments of the systems and methods disclosed herein, the structure, composite system, or both are characterized by a bounce coefficient selected from the range 0.8-0.3. In some embodiments of the systems and methods disclosed herein, the bounce coefficient is determined when the particles are accelerated in the structure, and the particles are at least 10 more than the physical dimensions of the unit cells of the structure. It has twice as large a coefficient, preferably the particle size is 7 μm to 14 μm, and the particles have a speed selected from the range of 500 m / sec to 1100 m / sec. In some embodiments of the systems and methods disclosed herein, the bounce coefficient corresponds to a structure having a relative density selected from the range of 8% to 26%.

In some embodiments of the systems and methods disclosed herein, the composite material system is configured to absorb collision energy substantially through said structure. In one embodiment, the composite system absorbs at least 20% more collision energy than is absorbed by the structure alone, otherwise under the same conditions.

In some embodiments of the systems and methods disclosed herein, the structure, composite system, or both are characterized by at least one vibration frequency bandgap. In some embodiments of the systems and methods disclosed herein, at least one vibration frequency bandgap is at least one complete vibration frequency bandgap. In some embodiments of the systems and methods disclosed herein, at least one vibration frequency bandgap is at least one partial vibration frequency bandgap. In some embodiments of the systems and methods disclosed herein, the structure, composite system, or both are characterized by at least one partial vibration frequency bandgap and at least one complete vibration frequency bandgap. .. In some embodiments of the systems and methods disclosed herein, at least one vibration frequency bandgap is deterministic. In some embodiments of the systems and methods disclosed herein, the at least one vibration frequency bandgap is in the range of 0.1 MHz to 100 MHz, preferably 0.1 MHz for some applications. It is in the range of ~ 200 MHz. In some embodiments of the systems and methods disclosed herein, the at least one vibration frequency bandgap is 52 MHz to 55 MHz, 19 to 26 MHz, 14 to 21 MHz, 46 to 52 MHz, 1.4 MHz to 1.5 MHz. At least one of them, or 2.4 ± 0.1 MHz. In some embodiments of the systems and methods disclosed herein, at least one vibration frequency bandgap corresponds to a structure exposed to continuous vibration, pulsed vibration, or a combination thereof. In some embodiments of the systems and methods disclosed herein, at least one vibration frequency band gap is characterized by a width selected from the range 0.1-200 MHz and at least one vibration frequency band. The gap is 0.1 to 100 MHz for some arbitrary applications, 100 Hz to 10 MHz for some arbitrary applications, 0.1 to 20 MHz for some arbitrary applications, preferably 0 for some applications. .1-10 MHz, 0.1 to 5 MHz for some applications, or 0.1 to 1 Hz for some applications. For example, at least one vibration frequency bandgap corresponds to a frequency associated with ultrasonic frequencies in medical applications. The wide vibration frequency bandgap width (eg, 200 MHz) helps dissipate a wide range of stimuli / vibrations in a variety of applications. The narrow vibration frequency bandgap width (such as 1 MHz) helps to filter specific stimuli / vibrations in a variety of applications. In some embodiments, the composite system comprises a wide, narrow, or wide and narrow vibration frequency bandgap. Also note that the vibration frequency bandgap (eg, its width) may scale in proportion to the characteristic length of the architecture. This is because it is shown that a remarkable effect can occur at a frequency of ΔL = c / f (ΔL is the dimension of the characteristic microstructure, c is the speed of sound of the material, and f is the frequency) depending on the Bragg condition. In some embodiments of the systems and methods disclosed herein, the structure, composite system, or both, is at least 1.2%, preferably at least 3.5%, more preferably at least 7%. Even more preferably, it is characterized by an attenuation ratio of at least 9%. In some embodiments of the systems and methods disclosed herein, the structure, composite system, or both are characterized by a damping ratio of at least 1.2% to 3.5% in the longitudinal direction. In some embodiments of the systems and methods disclosed herein, the structure, composite system, or both are characterized by a damping ratio of at least 7% -9% laterally. In one embodiment, the attenuation ratio corresponds to the ratio between the energy dissipated in the cycle and the maximum energy stored in the cycle.

In some embodiments of the systems and methods disclosed herein, the 3D geometry comprises at least one surface feature. Exemplary surface features include, but are not limited to, sheets, surfaces, hollow spheres, hollow ovals, and shells. In some embodiments of the systems and methods disclosed herein, the 3D geometry comprises at least one surface feature, at least a portion of the at least one surface feature being characterized by a non-zero Gaussian curvature. .. In some embodiments of the systems and methods disclosed herein, the 3D geometry comprises at least one surface feature, at least a portion of the at least one surface feature characterized by nonzero mean curvature. Be done. In some embodiments of the systems and methods disclosed herein, the 3D geometry comprises at least one surface feature, at least a portion of the at least one surface feature being characterized by zero mean curvature. In some embodiments of the systems and methods disclosed herein, the 3D geometry comprises at least one surface feature, said at least one surface feature due to non-uniform Gaussian curvature or non-uniform mean curvature. Characterized. In some embodiments of the systems and methods disclosed herein, the 3D geometry comprises at least one surface feature, said at least one surface feature being characterized by a non-uniform Gaussian curvature. In some embodiments of the systems and methods disclosed herein, the 3D geometry comprises at least one surface feature, said at least one surface feature being characterized by a non-uniform mean curvature. In some embodiments of the systems and methods disclosed herein, the 3D geometry comprises at least one surface feature, said at least one surface feature being characterized by a uniform Gaussian curvature or a uniform mean curvature. In some embodiments of the systems and methods disclosed herein, the 3D geometry comprises at least one surface feature, said at least one surface feature being characterized by a uniform Gaussian curvature. In some embodiments of the systems and methods disclosed herein, the 3D geometry comprises at least one surface feature, said at least one surface feature characterized by a uniform mean curvature. In some embodiments of the systems and methods disclosed herein, the 3D geometry comprises at least one surface feature and the thickness dimension of the at least one surface feature is non-existent over the entire at least one surface feature. It is uniform. In some embodiments of the systems and methods disclosed herein, the 3D geometry comprises at least one surface feature and the thickness dimension of the at least one surface feature is uniform over the at least one surface feature. Is.

In some embodiments of the systems and methods disclosed herein, 3D geometry is characterized as spinodal geometry. In some embodiments of the systems and methods disclosed herein, the structure has a relative density selected from the range 1-1.3, optionally 1-1.5, or optionally 1-1.35. Characterized by the gradient of the normalized effective modulus with respect to.

In some embodiments of the systems and methods disclosed herein, the 3D geometry comprises a resonator. In some embodiments of the systems and methods disclosed herein, the three-dimensional geometry is characterized by a unit cell geometry, which unit cell geometry comprises a resonator. In some embodiments of the systems and methods disclosed herein, the resonator comprises microinertial features connected to at least one other feature of the three-dimensional geometry. In some embodiments of the systems and methods disclosed herein, the resonator comprises a micro-inertia mechanism, wherein another mechanism of the three-dimensional geometry includes the micro-inertia mechanism. In some embodiments of the systems and methods disclosed herein, the resonator comprises a cantilever beam feature and a microinertial feature connected to the end of the cantilever beam feature. For example, the micro-inertia mechanism may be embedded within another mechanism, such as a beam or a structural member such as a surface. For example, microinertia can be at the nodes of the beam structure or at the hinges of the surface-based geometry.

In some embodiments of the systems and methods disclosed herein, the structure is, for example, a deterministic anisotropic collision energy absorption with collision energy absorption, or a collision energy absorption metric such as a bounce coefficient, eg, at least. 1% greater, at least 20% greater, at least 100% greater, preferably at least 1000% greater for some applications, or even more preferably 10000% greater along the first direction than along the second direction (eg, 10000% greater) , X, Y, Z or between any direction or vector) is characterized by a bounce coefficient. In some embodiments of the systems and methods disclosed herein, the structure has deterministic anisotropic elasticity, eg, at least 1% greater, at least 20% greater, at least 100% greater, preferably some applications. Is at least 1000% greater, or even more preferably at least 10000% greater along the first direction than along the second direction (eg, X, Y, Z or any direction or vector. Characterized by having an elastic modulus (between). In some embodiments of the systems and methods disclosed herein, the structure is, for example, at least 1% larger, at least 20% larger, at least 100% larger, preferably at least 1000% larger in some applications. Or even more preferably attenuation such as an attenuation ratio that is at least 10000% greater along the first direction (eg, between X, Y, Z or any direction or vector) than the second direction in some applications. Or characterized by deterministic anisotropic damping with damping metric. Vibrations, such as vibrations within a particular frequency range, such as any frequency or range described herein, are deterministically determined paths or components of the composite system, their structure, or their three-dimensional geometry. Follow. In some embodiments of the systems and methods disclosed herein, the structure exhibits oscillating Bragg scattering and the structure does not exhibit oscillating local resonance. In some embodiments of the systems and methods disclosed herein, the structure exhibits oscillating local resonance and the structure does not exhibit oscillating Bragg scattering. In some embodiments of the systems and methods disclosed herein, the structure is characterized by deterministic isotropic collision energy absorption. In some embodiments of the systems and methods disclosed herein, the structure is characterized by deterministic isotropic elasticity. In some embodiments of the systems and methods disclosed herein, the structure is characterized by deterministic isotropic decay.

The systems and methods disclosed herein are compatible with a wide variety of materials, or combinations of materials. The structure can be formed from any one or more materials, for example, as compared to laminated molding.

In some embodiments of the systems and methods disclosed herein, the structure comprises a carbon allotrope material, a polymer, a ceramic material, a metallic material, or any combination thereof. In some embodiments of the systems and methods disclosed herein, the structure comprises at least one of a carbon allotrope material, a polymer, a ceramic material, or any combination thereof. In some embodiments of the systems and methods disclosed herein, the structure comprises one or more carbon allotrope materials. In some embodiments of the systems and methods disclosed herein, the structure comprises at least 50% by volume of one or more carbon materials. In some embodiments of the systems and methods disclosed herein, the structure comprises a plurality of features characterized by a core that is at least 50% by volume of one or more carbon materials.

In some embodiments of the systems and methods disclosed herein, the 3D geometry is a nodeless geometry. In some embodiments of the systems and methods disclosed herein, the structure comprises at least one hollow feature. In some embodiments of the systems and methods disclosed herein, the structure comprises at least one feature that is at least partially hollow, such as a hollow truss. For example, the spinodal shape may include at least one hollow portion or feature.

In some embodiments of the systems and methods disclosed herein, the structure is formed by laminated molding.

In some embodiments of the systems and methods disclosed herein, the 3D geometry comprises at least one longitudinal feature and at least a portion of the at least longitudinal feature is in the longitudinal direction of the feature. Characterized by a non-zero curvature along. In some embodiments of the systems and methods disclosed herein, the 3D geometry comprises at least one longitudinal feature, said at least one longitudinal feature being non-uniform along the longitudinal direction of the feature. Characterized by a curved curve. In some embodiments of the systems and methods disclosed herein, the 3D geometry comprises at least one longitudinal feature having at least one cross-sectional dimension that is non-uniform along the longitudinal direction of the feature. In some embodiments of the systems and methods disclosed herein, the 3D geometry comprises at least one feature having a non-uniform cross-sectional shape. In some embodiments of the systems and methods disclosed herein, the three-dimensional geometry comprises at least one vertical feature having a vertical axis oriented perpendicular to the thickness direction of the structure. For example, the thickness direction corresponds to the axis formed along which each layer is formed in layer-by-layer laminated modeling.

In some embodiments of the systems and methods disclosed herein, the structure defines a three-dimensional external boundary shape. The three-dimensional geometry then includes at least one feature that intersects the boundary shape at only one or zero intersections.

In some embodiments of the systems and methods disclosed herein, the three-dimensional external boundary shape defined by the structure corresponds to the shape of the composite material system. In some embodiments of the systems and methods disclosed herein, the three-dimensional external boundary shape defined by the structure is hollow. For example, the three-dimensional outer boundary shape defined by the structure may be in the form of a hollow tube, hollow cone, hollow ellipse, or other configuration.

In some embodiments of the systems and methods disclosed herein, a 3D geometry is an overall 3D geometry that includes at least a primary 3D geometry and a secondary 3D geometry, said primary and secondary 3D. The geometry is different. For example, a structure can include a structural functional inclination of a three-dimensional geometry so that at least two structures are directly continuous and interconnected and all are directly or indirectly continuous and interconnected. Includes geometry. For example, the structure may include a compositional functional gradient. Through structural function tilts (ie, multiple 3D geometries, that is, 3D geometries include at least primary 3D geometry and secondary 3D geometry), the structure or composite system has collision energy absorption behavior and / or It can have a functional gradient of damping behavior.

In some embodiments of the systems and methods disclosed herein, the matrix phase comprises at least a primary matrix phase and a secondary matrix phase. For example, the first portion of the structure may be impregnated with the primary matrix phase and the second portion of the structure may be impregnated with the secondary matrix phase. The primary matrix phase and the secondary matrix phase may be different. For example, the primary and secondary matrix phases can have different compositions. For example, the primary matrix phase may be a different polymer or resin than the secondary matrix phase. In some embodiments of the systems and methods disclosed herein, the structure comprises a closed region without the matrix phase. In some embodiments of the systems and methods disclosed herein, the structure is encapsulated within said matrix phase. In some embodiments of the systems and methods disclosed herein, the structure is encapsulated within the matrix phase so that no portion of the structure is present beyond the outer boundaries of the matrix phase.

In some embodiments of the systems and methods disclosed herein, at least some of the three-dimensional geometries are tetrahedral, Weaire-Phelan, honeycomb, auxetic, octet truss shapes. , Octagonal, 3D cage geometry, square geometry, cubic geometry, tetrahedron, space-filled polyhedron, periodic tiny surface geometry, triple periodic tiny surface geometry, spinodal geometry, chiral geometry, or a combination thereof. Characterized. In some embodiments of the systems and methods disclosed herein, at least a portion of the 3D geometry is a beam-based, shell-based, open-cell-based, or closed-cell-based representation of the 3D geometry. Square geometry, cubic geometry, tetrahedron when viewing 3D geometry with beams characterized as tetrahedrons, auxetic shapes, octet truss shapes, octahedrons, diamond lattices, 3D cage shapes, and square crystals when viewed. , Space-filled polyhedron, periodic tiny surface geometry, triple periodic tiny surface geometry, spinodal geometry, chiral geometry, or a combination thereof. In embodiments of the systems and methods disclosed herein, at least a portion of the three-dimensional geometry is characterized by beam or shell-based geometry. Here, the beam or shell-based geometry is not symmetric, non-periodic, or periodically tessellated. In some embodiments of the systems and methods disclosed herein, features include one or more of struts, beams, ties, trusses, sheets, surfaces, spheres, ellipsoids, and shells.

In some embodiments of the systems and methods disclosed herein, the 3D geometry is characterized by a plurality of features that independently have physical dimensions that are independently selected within a tolerance of 100 nm or less. .. In some embodiments of the systems and methods disclosed herein, 3D geometry is characterized by a plurality of features that independently have physical dimensions that are independently selected within a tolerance of 3 μm or less. .. In some embodiments of the systems and methods disclosed herein, the structure is 5% to 99.9%, optionally 5% to 60%, optionally 0.1% to 60%, optionally 0. It is characterized by a relative density selected from the range of .1% to 99.9%, or optionally 8% to 30%. In some embodiments of the systems and methods disclosed herein, the structure is characterized by a relative density selected from the range of 8% to 60%. In some embodiments of the systems and methods disclosed herein, the structure is by a plurality of features independently having at least one physical dimension selected from 50 μm or less, optionally in the range of 100-200 μm. Be characterized. In some embodiments of the systems and methods disclosed herein, the 3D geometry is characterized by a plurality of features, at least some of which are independently selected from the range of 50 nm to 200 μm. One or more average cross-sectional physics selected from the range preferably 2 nm to 200 μm for some applications, 10 nm to 200 μm for some applications, optionally 50 μm or less, optionally 200 nm or less, or optionally 100 μm to 200 μm. Has geometric dimensions (eg, thickness, width, diameter, etc.). For example, the spinodal geometry can be substantially hollow, including walls or shells with a thickness of substantially 10 nm. For example, the structure includes hollow features such as hollow beams or trusses with inner and outer diameters, and the physical dimensions of the cross section are the thickness of the beam or truss, or the difference between the inner and outer diameters, or the total thickness. Or the diameter of the feature. For example, the inner diameter may be substantially 250 nm and the outer diameter may be substantially 900 nm. In some embodiments of the systems and methods disclosed herein, at least some of the features are characterized by one or more average longitudinal physical dimensions selected over the range of 10 nm to 2000 μm. In some embodiments of the systems and methods disclosed herein, the 3D geometry is characterized by a unit cell geometry, the unit cell being 10 nm to 20 μm, optionally 100 nm to 20 μm, optionally 1 μm to 20 μm. It has at least one overall physical dimension optionally selected from the range of 5 μm to 20 μm, or optionally 1 μm to 200 μm. For example, the unit cell can have square geometry with lengths, widths, and thicknesses of 20 μm, 20 μm, and 5 μm.

In some embodiments of the systems and methods disclosed herein, the composite system comprises at least one additional phase. In some embodiments of the systems and methods disclosed herein, the additional phase is a void or region without structural and matrix phases. In some embodiments of the systems and methods disclosed herein, the structure, matrix phase, or both comprises an adhesion promoting additive. In one embodiment, the adhesion promoting additive facilitates adhesion between the structure and the matrix phase as compared to adhesion without the additive. Additives may be present during the formation of the structure, such as being present in the precursor material during the formation of the structure. Additives may be deposited on the structure after the structure has been formed at least partially. Additives can be introduced during the impregnation of the structure in the matrix phase. Additives can be added to the matrix phase or matrix phase precursor prior to impregnation.

In some embodiments of the systems and methods disclosed herein, the structure is characterized by elasticity, the elasticity of said structure being deterministic.

In some embodiments of the systems and methods disclosed herein, the structure is from SiO ₂ particles having a diameter selected from the range of 7 μm 14 μm and a velocity selected from the range 500 m / sec to 1100 m / sec. It is virtually undamaged by the collision. In some embodiments of the systems and methods disclosed herein, SiO ₂ structure, than the overall dimensions of the unit cell characterizing a three-dimensional geometry of the structure has at least one order of magnitude greater (at least 10-fold) in diameter Substantially undamaged by collisions from particles.

In some embodiments of the systems and methods disclosed herein, the structure is characterized as having a bending control mode. In some embodiments of the systems and methods disclosed herein, the structure is characterized as having a stretch control mode.

In some embodiments of the systems and methods disclosed herein, the structure is characterized by an average specific strength (strength to density ratio) selected from the range 0.14 to 1.90 GPa / g · cm ^3. Can be attached. In some embodiments of the systems and methods disclosed herein, the structure is characterized by an average density selected from the range 0.24 to 1.0 g / cm ³ . In some embodiments of the systems and methods disclosed herein, the structure is characterized by an average Young's modulus selected from the range of 0.16 to 18.6 GPa. In some embodiments of the systems and methods disclosed herein, the structure is characterized by an average Young's modulus selected from the range of 0.16 to 440 GPa. In some embodiments of the systems and methods disclosed herein, the structure is characterized by compressive strength selected from the range of 5 MPa to 20 GPa. In some embodiments of the systems and methods disclosed herein, the structure is characterized by a fracture strain value of 20% or greater. In some embodiments of the systems and methods disclosed herein, the structure is characterized by a fracture strength value of 1 GPa or greater.

In some embodiments of the systems and methods disclosed herein, the carbon homologous material is from glassy carbon, graphite carbon, amorphous carbon, pyrolytic carbon, graphite, carbon black, and any combination thereof. It is selected from the group. In some embodiments of the systems and methods disclosed herein, the carbon allotrope material comprises pyrolytic carbon.

In some embodiments of the systems and methods disclosed herein, the structure comprises a coating. In some embodiments of the systems and methods disclosed herein, the coating comprises a metal, ceramic, or a combination thereof.

In some embodiments of the systems and methods disclosed herein, the matrix phase is selected from the group consisting of polymers, epoxies, carbon allotropes, ceramics, metals, viscous fluids, or any combination thereof. Contains one or more materials.

In one aspect, it is a method of making a composite material system, wherein the method is a step of preparing a structure by laminating modeling, the structure having a constructed 3D geometry, and the 3D geometry being monolithic. And the step of impregnating the structure with the matrix phase such that the structure is at least partially impregnated with the matrix phase, thereby making the composite system. ,including. In some embodiments, the structure is at least 1%, at least 2%, at least 5%, at least 10%, at least 15%, at least 20%, at least 30%, at least 50, depending on the matrix phase. By volume, at least 75% by volume, at least 90% by volume, or preferably substantially 100% by volume in some applications. In some embodiments, the structure is at least 1%, at least 2%, at least 5%, at least 10%, at least 15%, at least 20%, at least 30%, at least 50, depending on the matrix phase. By weight%, at least 75% by weight, at least 90% by weight, or preferably in some applications, it is impregnated with substantially 100% by weight. In some embodiments of the systems and methods disclosed herein, the 3D geometry is a nano or micro-constructed 3D geometry. In some embodiments, the method further comprises the step of designing the three-dimensional geometry using computer-aided design techniques. In some embodiments, the design step comprises determining the three-dimensional geometry based on a computational spinodal decomposition. In some embodiments, the design step comprises determining at least one vibration frequency bandgap of the structure such that at least one vibration frequency bandgap of the structure is deterministic. In some embodiments, the designing step comprises determining the elasticity of the structure so that the elasticity of the structure is deterministic. In some embodiments, the designing step comprises determining the restoration factor of the structure so that the restoration factor of the structure is deterministic. In some embodiments, the designing step determines the area-normalized collision energy relaxation metric (Ψ) of the structure so that the area-normalized collision energy relaxation metric (Ψ) of the structure is deterministic. Including steps. In some embodiments, the laminate molding process is a digital light processing (DLP) technique, a continuous liquid interface manufacturing technique, a microstereolithography (μ-SLA) technique. Two-photon polymerization lithography technology. Interference lithography technology. It is selected from the group consisting of holographic lithography technology, stimulated emission impairment (STED) lithography technology, other butt photopolymerization technology, material extrusion technology, powder bed fusion technology, material injection technology, and combinations thereof. In some embodiments, the laminate molding process is a three-dimensional lithography technique. In some embodiments, the steps of preparing the structure include forming a three-dimensional framework. The preparing step further includes a step of processing the three-dimensional framework to prepare the structure. The method comprises a pyrolysis process. In some embodiments, the method further comprises applying a coating to the structure. In some embodiments, the method further comprises etching a portion of the structure such that the structure comprises at least one hollow mechanism and at least a portion of the at least one hollow mechanism comprises the coating. .. In some embodiments, the impregnation step comprises a process selected from the group consisting of sonication, vacuum exposure, and any combination thereof. In some embodiments, the method further comprises a step of post-treating the composite material system, the step of post-treatment comprising a curing process. In some embodiments, the step of preparation involves selecting a precursor material selected from the group consisting of resins, metals, ceramics, polymers, and any combination thereof. In some embodiments, the step of preparation is to select a precursor material selected from the group consisting of organic resins, hybrid organic inorganic resins, metals, metal alloys, ceramics, polymers, and any combination thereof. Including. Useful organic resins include, but are not limited to, acrylic-based, thiol-based, polyurethane-based, and epoxy-based resins. Useful hybrid organic-inorganic resins include precursors derived from siloxanes and metal alkoxides, such as those described in US Patent Application Publication No. 2018/0088462, which is incorporated herein by reference. Not limited. In some embodiments, the precursor material further comprises an additive. Useful additives include metal and / or ceramic particles (eg nanoparticles), other inorganic and / or organic additives, solutions of inorganic and / or organic materials that form a single phase with the resin, or suspensions. Includes liquids, resins, inorganic particles, inorganic fibers, organic binders, polymer powders, metal powders, ceramic powders, metal wires (eg microwires or nanowires), metal salt solutions, metal ion solutions, and any combination thereof. However, it is not limited to these. For example, precursors for extrusion techniques can include thermoplastic polymers and low melting point metals, as well as additives including inorganic particles, fibers and the like. For example, material feeds for powder bed-based technologies can include metal, ceramic, and polymer powders. For example, precursors for the use of binder jet-based techniques may include these powders in combination with organic binders. For example, precursors for direct energy deposition techniques can include metal powders and metal wires. As additional examples, other raw material / precursor materials for laminate molding technology include metal and ceramic nanoparticles ink for direct ink writing and magnetohydrodynamic printing, metal and polymer sheets for laminate base processes, electroplating bases and Includes metal salts or metal ion solutions for photoreduction, precursor gases of focused beam based methods (FEBID / FIBID), and droplets of molten metal for laser-induced forward movement and magnetohydrodynamic printing.

In some embodiments, the pyrolysis process is carried out over a temperature range selected from the range of 500 ° C. to 3000 ° C. and over a period selected from the range of 1 hour to 336 hours. In some embodiments, the pyrolysis process is carried out over a temperature range selected from the range of 500 ° C to 900 ° C, optionally 500 to 1300 ° C. The structure is optionally exposed to any atmosphere during thermal decomposition, which is substantially devoid of vacuum and / or inert gases, or oxygen and water vapor. In some embodiments, the pyrolysis process provides isotropic contraction of the three-dimensional framework to the structure selected from the range of 15% to 80%.

In some embodiments, the method further comprises the step of applying an external stimulus to the structure to alter at least one vibration frequency bandgap of the structure.

In some embodiments, the method does not include the step of etching the template.

In some embodiments of the systems and methods disclosed herein, the composite material system is preferably 1000 kg / m ³ or less, 1500 kg / m ³ or less for some applications, 1200 kg / m ³ or less, More preferably less than 1000 kg / m ³ for some uses, more preferably 900 kg / m ³ or less for some uses, more preferably 500 kg / m ³ or less for some uses, more preferably for some uses It has a density of 300 kg / m ³ or less, more preferably 150 kg / m ³ or less for some applications, or even more preferably 100 kg / m ³ or less for some applications. In some embodiments of the systems and methods disclosed herein, the composite ^{system, 100kg / m 3 ~1500kg / m} 3, 100kg / m 3 ~1000kg / m 3 of less than the range or during any, Has a density selected from a subrange of. In some embodiments of the systems and methods disclosed herein, the composite material system or structure is at least 100 MPa, optionally at least 495 MPa for some applications, preferably at least 660 MPa for some applications, more preferably. Is characterized by a Young's modulus of at least 1000 MPa for some applications, at least 1.82 GPa for some applications, or at least 2.2 GPa for some applications. In some embodiments of the systems and methods disclosed herein, the composite system or structure is characterized by Young's modulus selected from the range of 100 MPa to 5 GPa, or any partial range in between. In some embodiments of the systems and methods disclosed herein, the composite material system or structure is at least 5 MPa, preferably at least 8 MPa in some applications, preferably at least 11 MPa in some applications, more preferably. Characterized by yield strength of at least 15 MPa in some applications, at least 25 MPa in some applications, at least 60 MPa in some applications, or at least 100 MPa in some even more preferred applications. Be done. In some embodiments of the systems and methods disclosed herein, the composite system or structure is characterized by a yield strength selected from the range of 5 MPa to 100 MPa, or any partial range in between. In some embodiments of the systems and methods disclosed herein, the composite system or structure exhibits catastrophic failure with strain (ε) selected from the range 0.1 to at least 0.5. It is characterized by not having. In some embodiments of the systems and methods disclosed herein, the composite system or structure is characterized by a flow stress selected from a range of at least 70 MPa or optionally 70-500 MPa. In some embodiments of the systems and methods disclosed herein, the composite system or structure is at least 10 MPa, preferably at least 20 MPa in some applications, or preferably 10 MPa to 100 MPa in some applications. Characterized by bending strength selected from a range or any subrange between them. In some embodiments of the systems and methods disclosed herein, the composite material system or structure is at least 1 GPa, preferably at least 1.4 GPa for some applications, preferably at least 2 GPa for some applications. Preferably at least 3.3 GPa for some applications, preferably at least 3.9 GPa for some applications, preferably at least 5 GPa for some applications, or more preferably 100 MPa to 10 GPa for some applications or these. Characterized by flexural modulus selected from any subrange between.

Also disclosed herein is a composite material system having any one or any combination of embodiments of the composite material system and method disclosed herein. Also disclosed herein are methods of making material systems having any one or any combination of composite material systems and embodiments of methods disclosed herein.

Without wishing to be bound by any particular theory, the beliefs or understanding of the underlying principles of the devices and methods disclosed herein can be discussed herein. Regardless of the final accuracy of the mechanical explanation or hypothesis, it is recognized that embodiments of the present invention can nevertheless be valid and useful.

It is a flowchart of the method of manufacturing the composite material which has an arbitrary continuous phase. An embodiment of a composite having a continuous three-dimensional phase with arbitrary geometry, a variable cross section in a single truss element. An embodiment of a composite having a continuous three-dimensional phase with arbitrary geometry, a non-zero curvature and horizontal truss element. It is an embodiment of a composite material having a continuous three-dimensional phase of shell or surface shape and having one or more matrix phases. It is an embodiment of a modular three-dimensional structural element made of the composite material described in the previous embodiment. No molding process is required during the strengthening phase and the geometry is predefined. An embodiment of a structural element made of composite material as described in the previous embodiment, in which the in-plane and thickness geometry is functionally tilted. The obtained member has a continuous three-dimensional phase. An embodiment of a structural element made of composite material as described in the previous embodiment, in which the in-plane and thickness geometry is functionally tilted. The obtained member has a continuous three-dimensional phase. An embodiment of a phase microarchitecture in which the resonator dissipates vibrations and provides damping to the material. The obtained member has a continuous three-dimensional phase. An embodiment limited to the practice of samples of composites with continuous three-dimensional phases with arbitrary geometry, structural elements made from precursor resins via DLP3D printing and corresponding micrographs. It is an embodiment limited to the implementation of a sample of a composite material having a continuous three-dimensional phase with arbitrary geometry, a three-dimensional continuous strengthening phase of carbon after pyrolysis and a corresponding micrograph. It is an embodiment limited to the implementation of a sample of a composite material having a continuous three-dimensional phase having an arbitrary geometry, and is a composite structural element having a continuous three-dimensional phase and an epoxy matrix phase. It is a figure which shows the manufacturing process of the continuous strengthening phase of arbitrary shape by the laminated molding. It is a schematic diagram of a structure having a three-dimensional geometry or a unit cell thereof, that is, a tetrahedron. It is a schematic diagram of a composite material system having a structure having a three-dimensional geometry or a unit cell thereof, i.e. a structure of FIG. 9A impregnated with a matrix phase. It is a polymer octet lattice, and is a photomicrograph of a complete lattice suspended on a bed with a spring structure. It is a polymer octet lattice, and is an enlarged view showing a unit cell of a 5 μm octet. It is a pyrolytic carbon lattice and is an octet lattice with a relative density of 26%. It is a pyrolytic carbon lattice, and is an enlarged view showing a submicron unit cell. An octet grid with a relative density of 26%. It is a pyrolytic carbon lattice and is a dodecahedron lattice with a relative density of 17%. It is a pyrolytic carbon lattice and is an enlarged view of the final tetrahedral dodecahedron geometry. It is a figure which shows the influence of the SiO ₂ beads of 7 μm on the octet lattice of a relative density 27%. Both the collider and the grid in the initial frame before the collision are highlighted in red. Subsequent frames reveal the bounce of the collider from the grid, ejecting small debris. Collision speed and bounce speed were 1060 and 560 m / sec, respectively. It is a PMMA resist coating for better bonding the lattice to the substrate and shows the initial carbon lattice. It is a PMMA resist coating for better bonding the lattice to the substrate, and shows a carbon lattice with a submicron PMMA coating. A post-collision confocal microscopic image showing a collision site and slight permanent deformation, which is a PMMA resist coating for better adhesion of the grid to the substrate. FIG. 14A shows an approximation of a collision as a plane wave passing through a grid material, showing a plane wave passing through a grid of thickness x _lat and a substrate of thickness x _s . FIG. 14B shows an approximation of collision as a plane wave through the lattice material, which is an x-t diagram of the worst case elastic plane wave of a carbon octet lattice with a relative density of 60%, with a collision time of (average) 4 ns. It shows that elastic waves are not enough to reach the substrate and reflect off the projectiles. It shows a floating sample manufacturing process, where the original Si substrate was patterned and etched to create stanchions with a height of about 50 μm. It shows the floating sample manufacturing process and shows nanowires in which samples that remain pyrolyzed are tethered to a substrate. Demonstrating the floating sample manufacturing process, the released sample was taken by a nanomanipulator. Demonstrating the floating sample manufacturing process, the samples were secured to the stanchions using Pt adhesive. It shows a floating sample collision experiment, and the same repulsive behavior was observed, but there was no complete penetration in the thickness direction or catastrophic sample damage. The summary of the collision experiment is shown. It can be observed trends between the collision energy _mv ^{0 2} and bounce coefficient _v r / _{v 0.} There was no significant difference between the different architectures, but there was a decrease in repulsive force in the low relative density samples. Floating sample experiments are shown in yellow and PMMA coated experiments are shown in blue. It shows a square unit cell of interest, showing an unchanged square unit cell with an elliptical cross-section beam in the horizontal direction and a circular beam in the vertical direction. It shows a square unit cell of interest and shows a slightly buckled geometry. It shows a square unit cell of interest, showing an unchanged square unit cell with an elliptical cross-section beam in the horizontal direction and a circular beam in the vertical direction. A square unit cell of interest is shown, which is a top view of the unit cell in FIG. 18A. A square unit cell of interest is shown, which is a top view of the unit cell in FIG. 18B. A square unit cell of interest is shown, which is a top view of the unit cell in FIG. 18C. FIG. 19A shows the dispersion relation of the square unit cells, and shows the unchanged unit cells. FIG. 19B shows the dispersion relation of the square unit cells and shows the slightly buckled unit cells. FIG. 19C shows the dispersion relation of the square unit cells, and shows the completely buckled unit cells. FIG. 19D shows the irreducible Brillouin zone in reciprocal space, with partial bandgap shaded in green. It shows the auxiliary unit cell and shows the unchanged unit cell. It shows an auxiliary unit cell and shows a unit cell with an added resonator. FIG. 21A shows the dispersion relation of the auxetic units, showing an unchanged unit cell without a bandgap. FIG. 21B shows the dispersion relation of the auxetic unit, shows the unit cell provided with the resonator, and shows the appearance of the band gap. FIG. 22A shows a bespoke ultrasonic setup, showing an image of a transmission experiment in which one transducer drives a signal through a material of interest and another transducer picks up the transmitted signal. FIG. 22B shows a bespoke ultrasonic setup, showing a picture of the vacuum chamber and the arrangement of the transducers. FIG. 22C shows a bespoke ultrasonic setup, a CCD image of the setup in vacuum with a rubber pack between the transducers for verification purposes. FIG. 23A shows a frequency sweep experiment, showing an unchanged unit cell. FIG. 23B shows a frequency sweep experiment, showing a unit cell with a resonator. FIG. 23C shows a frequency sweep experiment and shows the amplitude of the frequency sweep transmission signal. A bandgap centered around 2.4MHz has been found and is highlighted in gray. FIG. 23D shows a frequency sweep experiment and is an enlarged view of a non-contact sample. FIG. 23E shows a frequency sweep experiment, which is an enlarged view of a sample slightly distorted by contact with a transducer. FIG. 24A shows a chirp transmission experiment, showing an input chirp signal including 1-3 MHz. FIG. 24B shows a chirp transmission experiment, which is an FFT of a transmission signal via a rubber pack, and the frequency component of the chirp was confirmed. FIG. 24C shows a chirp transmission experiment, which is an FFT of a transmission signal through a sample including a resonator, and shows a bandgap of about 2.4 MHz at the center, which is the same as the sweep experiment. It shows the elastic surface of the microstructure of some samples and shows the columnar structure showing the hard direction aligned with the z-axis. It is a cubic structure showing a priority direction aligned with the x, y and z axes. It is a lamellar structure showing that the magnitudes of rigidity differ by several orders of magnitude along the x-axis, y-axis, and z-axis. It has a quasi-isotropic structure. Showing stiffness scaling as a function of relative density, spinodal structures have higher absolute coefficients and more desirable, lower scaling indices. FIG. 27A is a curvature probability distribution for the strut structure. FIG. 27B is a probability distribution of curvature for an octet truss with fillets of 0.5r and r at the nodes, where r is the radius of the tube. The absolute curvature of the octet structure is significantly higher than the curvature of the spinodal structure. It is a nano / microscale shell-based spinodal structure, a polymer structure made by two-photon lithography. It is a nano / microscale shell-based spinodal structure, a structure coated after FIB milling, which exposes the polymer again. It is a nano / microscale shell-based spinodal structure, which is a shell-based spinodal structure after etching the inside of the polymer. It is a shell-based spinodal structure on a macro scale and is a polymer columnar spinodal structure. It is a shell-based spinodal structure on a macro scale, and is a top view of FIG. 29A. It is a shell-based spinodal structure on a macro scale and is the resulting carbon-enhanced phase. It is a shell-based spinodal structure on a macro scale, and is a top view of FIG. 29C. Constructed plates for blast impact testing, polymer precursor plates and resulting pyrolysis plates. It is a constructed plate for a blast impact test and is a photomicrograph of an octet carbon structure. It is a constructed plate for a blast impact test, and is a photomicrograph of a dodecahedron structure. An example cube of the strengthening phase, a polymer octet cube. An example cube of the strengthening phase, the pyrolytic carbon octet cube from FIG. 31A. An example cube of the strengthening phase, a pyrolytic carbon 3D Kagome beam. It is an enlarged view of FIG. 31A. It is an enlarged view of FIG. 31B. It is an enlarged view of FIG. 31C. It is a tubular constructed component and is a polymer tube with a dodecahedron structure. It is a top view of FIG. 32A. It is an enlarged view of FIG. 32B. It is a pyrolytic carbon tube having a decahedron structure before impregnation. It is an enlarged view of FIG. 32D. It is an enlarged view of FIG. 32D. FIG. 33A shows the manufacture of pyrolytic carbon micropillars and the characterization of the microstructure and is a schematic representation of the manufacturing process. This process involves TPLDLW of columnar pillars from an IP-Dip polymer resin, followed by thermal decomposition under vacuum at 900 ° C. FIG. 33B is an SEM image of a typical micropillar before thermal decomposition. FIG. 33C is an SEM image of a typical micropillar after thermal decomposition, showing considerable volume shrinkage. FIG. 33D is a bright-field TEM image of pyrolytic carbon. The diffraction pattern in the inset shows an amorphous structure. FIG. 33E is an HRTEM image of two regions contoured by the solid square of FIG. 33D. It indicates the presence of some sub-nanometer-sized voids (indicated by red arrows). FIG. 33F is an HRTEM image of two regions contoured by the solid square of FIG. 33D. It indicates the presence of some sub-nanometer-sized voids (indicated by red arrows). FIG. 33G is a Raman spectrum of a pyrolytic carbon micropillar. Typical G and D bands at energies 1359 cm ^-1 and 1595 cm ^-1 show sp ² hybrids. FIG. 33H shows the EELS of pyrolytic carbon, and the shaded green and purple parts correspond to the 1s-π ^* peak and the 1s-σ ^* peak of carbon, respectively. Quantitative analysis of the data shows that pyrolytic carbon contains about 96.55% sp ² bonds. FIG. 34A shows a uniaxial compression and tensile test of a pyrolytic carbon micropillar, which is compressive stress-strain data for a pyrolytic carbon pillar with a diameter of 4.6 to 12.7 μm. All of these micropillars were elastically deformed to a maximum strain of about 20 to 30% and showed a slight plastic strain (about 8 to 10%) before breaking. The dashed line shows the slope of the linear part. FIG. 34B is an SEM image of the typical pyrolytic carbon micropillars shown in FIG. 34A before and after compression, showing the occurrence of brittle fracture by multiple fragments. FIG. 34C shows a typical stress-strain data set from in-situ deformation of a 2.25 μm diameter pyrolytic carbon pillar, with significant plastic deformation up to 33.6% strain occurring. The inset shows an SEM image of the micropillars before compression. A series of snapshots taken during the in-situ deformation are displayed on the plot, with the numbered frames corresponding to the red arrows of the same number on the stress-strain curve. The SEM images to the right of the stress-strain data are front and back views of the compressed micropillars. Nucleation and propagation of split rhagades correspond to strain bursts indicated by the blue arrows on the stress-strain curve. FIG. 34D is tensile stress strain data of a bone-shaped sample of a pyrolytic carbon dog having a gauge diameter of 0.7 to 2.0 μm. FIG. 34E is an SEM image of a typical tensile test piece before and after the experiment. FIG. 34F is a statistical distribution of tensile fracture strength. FIG. 35A shows the change in strength with the diameter of the pyrolytic carbon micropillar and the super-large elastic limit, and shows the change in the compressive strength with the increase in the diameter of the micropillar. In the figure, the blue dashed line represents the average compressive strength of micropillars with a diameter of less than 2.3 μm. FIG. 35B is a 20-cycle force-displacement curve of a deformable pillar with a diameter of 1.28 μm with a maximum compressive strain of about 23%, showing near complete recovery in all cycles except the first cycle. ing. The SEM image shows the pillars before and after deformation from 20 load cycles. FIG. 36A shows an atomic simulation of uniaxial compression and tension of a pyrolytic carbon nanopillar, showing the atomic composition and cross-sectional morphology of a simulated sample with a diameter of 20 nm. FIG. 36B is a compressive stress-strain curve of a pyrolytic carbon nanopillar. FIG. 36C is a tensile stress-strain curve of a pyrolytic carbon nanopillar. FIG. 36D is a snapshot of pillars deformed with different compression strains. Von Mises Color-coded according to atomic strain. FIG. 36E is a snapshot of pillars deformed with different compression strains. Von Mises Color-coded according to atomic strain. FIG. 36F is a snapshot of pillars deformed with different compression strains. Von Mises Color-coded according to atomic strain. FIG. 36G is a snapshot of pillars deformed with different compression strains. Von Mises Color-coded according to atomic strain. FIG. 36H is a snapshot of pillars deformed by different tensile strains. Von Mises Color-coded according to atomic strain. FIG. 36I is a snapshot of pillars deformed by different tensile strains. Von Mises Color-coded according to atomic strain. FIG. 36J is a snapshot of pillars deformed by different tensile strains. Von Mises Color-coded according to atomic strain. FIG. 37A outlines the combination of ultra-high strength / specific strength and high deformability of pyrolytic carbon micropillars and is an ashby chart of the strength vs. density of various structural materials including pyrolytic carbon micropillars. FIG. 37B is a diagram showing a comparison of tensile and compression specific strengths between our pyrolytic carbon micropillars and other structural materials. FIG. 37C is a summary of the specific strength and fracture strain of pyrolytic carbon micropillars and other structural materials. The excellent combination of specific strength and deformability of our pyrolytic carbon exceeds that of almost all other materials. It is a compressive stress-strain curve of nanopillars simulated at different densities with a diameter of 10 nm. It is a plot of Young's modulus (GPa) vs. density (g / cm ³ ) corresponding to a material or structure from a related technique and a particular embodiment of the invention. It is an image corresponding to the atomic arrangement of a nanopillar formed of a carbon allotrope material. It is an image corresponding to the atomic arrangement of a nanopillar formed of a carbon allotrope material. It is an image corresponding to the atomic arrangement of a nanopillar formed of a carbon allotrope material. It is an image corresponding to the atomic arrangement of a nanopillar formed of a carbon allotrope material. It is an image corresponding to the atomic arrangement of a nanopillar formed of a carbon allotrope material. The nanopillars shown in FIGS. 40A-40E have a diameter of 10 nm and a density of 1.0-1.8 g / cm ³ . The presence of sp carbon, sp ² carbon, and sp ³ carbon has been identified. A density estimation model and a comparison of pyrolytic carbon densities reported in recent literature. a and b are diagrams of the packed structure of the curved graphene layer in the pyrolytic carbon, L represents the size of the curved graphene layer, and L _s represents the interlayer distance between adjacent layers. c is a typical open structure unit cell composed of two graphene layers. d is the density of pyrolytic carbon (ρ _PC ) as a function of the ULS ratio L / L _s . The solid line curve was obtained from the prediction based on the equation (1), and the broken line curve was obtained from Reference 26. Current extended models predict that the density of pyrolytic carbon micropillars is 1.0-1.8 g / cm ³ . An in-situ compression experiment of a pyrolysis micropillar without a residual ring is shown. a, b and c are snapshots of an in-situ compression test of pyrolytic carbon pillars without residual rings, where c is a fission crack nucleating and rapidly propagating under high compressive stress to destroy the micropillars. Destruction. d is the corresponding compressive stress-strain curve. It shows the effect of the residual carbon ring on the compression of pyrolytic carbon micropillars. a, b and c are snapshots of the in-situ compression test of the pyrolytic carbon pillars with the remaining rings. d is the corresponding compressive stress-strain curve, and the slight burst marked with "b" corresponds to the bulge of the ring edge due to high stress concentration. Large strain bursts marked with an "c" represent cleavage of the pillars and detachment of the ring. It shows the bond structure of pyrolytic carbon pillars used in atomistic simulations. The sp ² bond is much more ubiquitous than the sp and sp ³ bond. The sp bond is mainly localized at the edge of the curved graphene layer. sp ³ bonds are usually either connecting adjacent graphene layers to each other, it is formed in the high energy surface of the graphene layer. It is a figure explaining the fracture mechanism of the pyrolytic carbon nanopillar under uniaxial tension, and is the snapshot of the extended nanopillar with the strain of 56.3-60.5%. Nanoscale cavities (indicated by orange arrows) nucleated and grew during elongation and then fused with each other to form nanoscale cracks. A snapshot of nanopillars stretched with a strain of 61.0-61.8%. As the tensile strain increases, nanoscale cracks propagate along the direction perpendicular to the tensile direction, resulting in a smooth fracture surface. All atoms in FIGS. 45A and 45B are color coded by atomic von Mises strain. It is a figure which shows the influence of the initial defect on the tensile strength of a pyrolytic carbon pillar. a and b are the atomic configurations of the simulated sample containing initial cracks of length 4 nm and 8 nm, respectively, with all initial cracks indicated by white flakes. c and d are a series of snapshots of pillars with initial cracks of lengths of 4 nm and 8 nm, respectively. The destruction of both nanopillars always began with the growth and expansion of existing nanocracks. Both samples after fracture had a smooth fracture surface, indicating a brittle fracture mode. All atoms at c and d are color coded by atomic von Mises strain. A summary plot of the strength and fracture strain of pyrolytic carbon micropillars and other structural materials. FIG. 48A shows the fabrication of pyrolytic carbon nanolattices and the characterization of microstructures, and is a schematic diagram of the manufacturing process of pyrolytic carbon nanolattices. FIG. 48B is a CAD representation of the unit cell of the octet truss. FIG. 48C is an SEM image of an octet nanolattice with a strut diameter of d = 435 nm. FIG. 48D is a CAD representation of the Isotras unit cell. FIG. 48E is an SEM image of an isotras nanolattice with a vertical strut diameter of d ₁ = 460 nm and a tilted strut diameter of d ₂ = 523 nm. FIG. 48F is an HRTEM image of pyrolytic carbon extracted from the nanolattice, showing the amorphous nature of the pyrolytic carbon. FIG. 49A shows an in-situ uniaxial compression experiment of a pyrolytic carbon nanolattice, showing the mechanical response of a pyrolytic carbon octet truss nanolattice with different relative densities obtained from in-situ compression. FIG. 49B shows the mechanical response of pyrolytic carbon isotras nanolattices with different relative densities obtained from in-situ compression. FIG. 49C is an SEM image of the octet truss nanolattice before compression at d = 382 nm. In the figure, the initial detectable structural defects caused by the manufacturing process are enclosed. FIG. 49D is an SEM image of the octet truss nanolattice after compression at d = 382 nm. FIG. 49E is an SEM image of the uncompressed isotras nanolattice at d ₁ = 538 nm and d ₂ = 612 nm. In the figure, the initial detectable structural defects caused by the manufacturing process are enclosed. FIG. 49F is an SEM image of the compressed isotras nanolattice at d ₁ = 538 nm and d ₂ = 612 nm. It is a logarithmic map of the mechanical properties of the pyrolytic carbon nanolattice and shows the Young's modulus of the pyrolytic carbon nanolattice plotted against density on a logarithmic scale. It shows the compressive strength of the pyrolytic carbon nanolattice plotted against density on a logarithmic scale. In the charts of FIGS. 50A and 50B, for comparison, an alumina hollow nanolattice (Reference 11), an airgel microlattice (Reference 23), a glassy carbon nanolattice (Reference 24), a cell-shaped carbon microstructure (Reference 25), Includes several previously reported micro and nano-constructing materials such as SiOC microlattices (Reference 26). It shows a finite element simulation of uniaxial compression of a pyrolytic carbon nanolattice using different unit cells, and shows a simulated configuration of an octetlas nanolattice containing existing defects introduced by imposing initial deflection on the struts. Shown. The inset shows an enlarged view of the local structure, including the initial deflection of the stanchions. It shows a simulated configuration of an Isotras nanolattice containing existing defects introduced by imposing initial deflection on the stanchions. The inset shows an enlarged view of the local structure, including the initial deflection of the stanchions. It shows a simulated configuration of a nanolattice of a tetrahedral truss containing existing defects introduced by imposing initial deflection on the struts. The inset shows an enlarged view of the local structure, including the initial deflection of the stanchions. It is a compressive stress-strain curve of the octet truss nanolattice. It is a compressive stress-strain curve of the Isotras nanolattice. It is a compressive stress-strain curve of a tetrahedral truss nanolattice. It shows a comparison of the specific strengths of our pyrolytic carbon nanolattices with other micro and nanolattices reported so far. FIG. 53A shows an in-situ compression test of a polymer nanolattice, which is a compressive stress-strain curve of an octet truss nanolattice with d = 1.12 μm. FIG. 53B is an SEM snapshot of the deformed octet truss nanolattice under various compressive strains. FIG. 53C is an SEM snapshot of a deformed octet truss nanolattice under various compressive strains. The circled area indicates buckling of the struts during compression. FIG. 53D is an SEM snapshot of a deformed octet truss nanolattice under various compressive strains. FIG. 53E is a compressive stress-strain curve of an isotras nanolattice with d ₁ = 1.30 μm and d ₂ = 1.49 μm. FIG. 53F is an SEM snapshot of the deformed isotras nanolattice under various compressive strains. FIG. 53G is an SEM snapshot of the deformed isotras nanolattice under various compressive strains. The circled area indicates buckling of the struts during compression. FIG. 53H is an SEM snapshot of the deformed isotras nanolattice under various compressive strains. FIG. 54A shows the density of Young's modulus-to-pyrolytic carbon nanolattices, which is the Young's modulus-to-density of octets and isotras pyrolytic carbon nanolattices on a logarithmic scale. The gradient of the power law to be scaled is shown for each architecture. The error bars represent the standard deviation from the mean of some data for samples of comparable density. FIG. 54B shows the compressive strength vs. pyrolytic carbon nanolattice density, which is the compressive strength vs. density of the octet and isotras pyrolytic carbon nanolattices on a logarithmic scale. The gradient of the power law to be scaled is shown for each architecture. The error bars represent the standard deviation from the mean of some data for samples of comparable density. It is a figure which shows the relative decrease of the intensity as a function of the range of the initial deflection of the nanolattice with the initial deflection. FIG. 56A shows a comparison between finite element modeling and experimental results, showing Young's modulus vs. relative density from finite element modeling and experimentation. The dependence of the Young's modulus of the nanolattice on the relative density from finite element modeling is consistent with the dependence from experimental measurements. FIG. 56B shows a comparison between finite element modeling and experimental results, showing intensity vs. relative density from finite element modeling and experiments. The dependence of nanolattice strength on relative density from finite element modeling is consistent with the dependence from experimental measurements. FIG. 57A is a 3D constructed carbon structure, a cross-section SEM image and a cross-section energy dispersive spectroscopy (EDS) spectrum. The scale bar is 5 μm. FIG. 57B is experimental Raman spectrum data (.), Which is a fitting curve (dotted line) for each band and a linear combination (red line) of these peaks. FIG. 57C is an X-ray diffraction (XRD) pattern. FIG. 57D is a transmission electron microscope (TEM) high resolution image and diffraction pattern (entrance). The scale bar is 5 nm. FIG. 58A is a line analysis of EDS in cross section. FIG. 58B shows particles ground from a 3D constructed carbon structure used for XRD analysis. FIG. 59A is a typical stress-strain curve of compression, with Roman numerals corresponding to the individual events shown in FIG. 59B. FIG. 59B is a photographic image of the compressed 3D constructed carbon structure at the local fracture shown by the red dotted circle. I is the image at the time of initial contact, II-a, b, c are the images at the time of partial local collapse shown in the red line, III is the image before the second stress release, IV is shown by the red line. The image at the time of partial layer collapse, V is the image before the third stress release, and VI is the image at the time of half layer collapse. The board and upper load cell are grayed out. Stress-strain curves of five samples of 3D constructed carbon structure. An image showing a constructed three-dimensional structure with nodeless geometry according to a particular embodiment of the invention, an additional exemplary nodeless geometry is found in Abuida et al. ("Triple periodic minimum". Effective Conductivity and Modulus of New Foams with Surfaces, Mechanics of Materials, vol. 95, April 2016, pp. 102-115), which is incorporated herein by reference. It is an image showing a constructed three-dimensional structure having nodeless geometry according to a specific embodiment of the present invention. It shows the impregnation of the shell-based spinodal-enhanced phase. It shows the impregnation of a reinforced phase with a dodecahedron tube structure. It shows the impregnation of the reinforcing phase with an octet cubic structure. It is an epoxy impregnated composite material from the reinforced phase shown in FIG. 62A. It is an epoxy impregnated composite material from the reinforced phase shown in FIG. 62B. It is an epoxy impregnated composite material from the reinforced phase shown in FIG. 62C. FIG. 63A is an octet carbon material in compression, which is a sample before compression. FIG. 63B is a sample destroyed after a catastrophic event. FIG. 63C is the final geometry of the tested sample and shows the fracture surface. FIG. 63D shows the stress-strain response of two identical samples. FIG. 64A shows a compressed octet carbon epoxy material, which is a sample before compression. FIG. 64B shows a densified sample after s = 0.5. FIG. 64C shows the stress-strain response of the sample. FIG. 65A is a diagram illustrating four-point bending of the carbon octet material, which is a sample before the experiment. FIG. 65B is a sample after a catastrophic destruction event. FIG. 65C is a morphology of the final sample. FIG. 65D shows the resulting stress-strain response, which corresponds to the outermost edge of the material. FIG. 66A is a diagram illustrating four-point bending of an epoxy-impregnated carbon octet material, which is a sample before the experiment. FIG. 66B is a sample with the highest bending load. FIG. 66C is a morphology of the sample after unloading showing no visible damage. FIG. 66D shows the resulting stress-strain response, corresponding to the outermost edge of the material.

In general, the terms and phrases used herein have their technically recognized meanings and can be found by reference to standard texts, literature and contexts known to those of skill in the art. The following definitions are provided to clarify their particular use in the context of the present invention.

The term "monolithic" refers to a system, structure, shape, or other element that is a single interconnected continuous element. In one embodiment, the monolithic element is formed or composed of a joint or seamless material. In one embodiment, the term "interconnection" refers to a system, structure, shape, or other element, each of which is a first portion or feature of (i) a system, structure, geometry, or other. Directly connected to a second part or second feature of an element, or (ii) a system, structure, geometry, or to a second part or second feature of a system, structure, geometry, or other element. , Or indirectly connected through a third part or third feature of other elements. In one embodiment, the interconnected system, structure, geometry, or part or feature of other elements is not completely separated from the rest of the system, structure, geometry, or other element. In one embodiment, the term "continuous" refers to a system, structure, geometry, or other element, each of which has its first part or feature a second part of the system, structure, geometry, or other element. Or to a feature, directly or indirectly combined, fused, or belong to the same uninterrupted phase. In one embodiment, the two features that are simply connected by superficial contact (contact, eg, touch), but otherwise separated from each other, are not continuous. In one embodiment, two distinct features, such as fibers or particles, which are simply in contact or woven together, can be interconnected, but not continuous with each other. In one embodiment, a structure or geometry consisting of multiple features, such as fibers or particles, each of which is simply in contact with or interwoven with another feature, such as fibers or particles, may be an interconnected structure or geometry. It is not a continuous structure or geometry.

The term "deterministic" is characterized by at least one known feature and / or at least one property (eg, vibration frequency bandgap) and / or 20% of the determined or desired value, preferably 10. Refers to a system, structure, shape or other element that is controlled within%, more preferably within 5%, more preferably within 1%, or even more preferably within 0.1%. In one embodiment, the deterministic geometry is 20%, preferably within 10%, more preferably within 5%, more preferably within 1% of the determined or desired value before or during the formation of the structure. Or more preferably characterized by one or more features, each independently having at least one predetermined physical dimension within 0.1%. For example, a deterministically constructed 3D geometry of a structure contains multiple features such as trusses and has one or more physical dimensions (eg, width, thickness, diameter, length) and its values. Is within 20%, preferably within 10%, more preferably within 5%, more preferably within 1% of the value of one or more physical dimensions designed by CAD technology or the like or determined prior to formation. , Or even more preferably within 0.1%. Probabilistic geometries and structures such as random and natural forms are not deterministic.

The term "construction" refers to a system, structure, geometry, or feature that is designed and has features formed according to that design. In one embodiment, the designed structure is formed according to a deterministic or deterministic process. In one embodiment, substantially all features and their physical dimensions are designed or predetermined so that substantially all features and their physical dimensions are substantially equivalent to those of the design according to the design. Is formed in.

The term "three-dimensional geometry" refers to geometry characterized by a three-dimensional geometry composition. In one embodiment, the structure has three-dimensional geometry when a three-dimensional system of physical space is needed to fully describe the physical dimensions of the unit cells of the structure. 3D geometry can be nano-designed and / or micro-designed. In one embodiment, the structure characterized by nanodesigned three-dimensional geometry has at least one physical dimension (eg, length, width, diameter, or height) whose value ranges from about 1 nm to less than 1 μm. A structure characterized by one or more features having. The one or more "features" include, but are not limited to, beams, stanchions, ties, trusses, seats, shells, and nodes. In one embodiment, the structure characterized by the nanodesign three-dimensional geometry has at least one physical dimension (eg, length, width, or height) whose value ranges from about 1 nm to less than 1 μm. It is a structure characterized by a unit cell. In one embodiment, the structure characterized by the microdesigned three-dimensional geometry has at least one physical dimension (eg, length, width, or height) whose value ranges from about 1 μm to 1000 μm. A structure characterized by one or more features having. In one embodiment, the structure characterized by microconstructed three-dimensional geometry is a unit cell having at least one physical dimension (eg, length, width, or height) that is a value in the range of 1 μm to 1000 μm. It is a structure characterized by.

As used herein, the "features" of a system, such as a composite material system by embodiment, structure, or geometry such as three-dimensional geometry by embodiment, are beams, struts, ties, trusses, seats, shells, spheres. Refers to, but is not limited to, elements such as, ellipsoids, nodes, or combinations thereof. In one embodiment, fillets, bevels, chamfers, or similar attributes are part of the feature, but not the feature itself. For example, fillets, or rounded inner or outer corners, are part of one or more features, but not the "features" themselves as used herein. For example, the fillet between the first truss and the second truss is part of the first truss, part of the second truss, or part of each of the first and second trusses. The fillet itself is not a "feature" of the three-dimensional geometry or structure used herein. A "longitudinal feature" is at least 50 in length (or size along its longitudinal axis) than each of the other characteristic size dimensions (ie, width, height, thickness, or diameter). % Refers to a larger element. Exemplary longitudinal features may include, but are not limited to, beams, stanchions, ties, and trusses. In one embodiment, surface features are features that can be better characterized as flat and / or curved planar features than longitudinal features. In one embodiment, the surface features correspond to features that can be approximated or characterized as mathematical two-dimensional manifolds with uniform or non-uniform thickness. In one embodiment, the surface feature is an open surface that corresponds to a feature that can be approximated or characterized as a mathematical two-dimensional manifold with uniform or non-uniform thickness. Exemplary surface features include, but are not limited to, sheets and shells.

"Matrix phase" refers to a material or combination of materials that can at least partially impregnate the structure of a composite material system. The matrix phase may or may not be uniform. The matrix phase may or may not be homogeneous. At least partial impregnation of the structure refers to at least partial filling of the voids in the structure. In one embodiment, at least partial impregnation of the structure refers to at least partial filling of accessible voids in the structure. An inaccessible void region of a structure refers to a closed void region to the matrix phase (eg, a hollow truss or hollow portion of spinodal geometry) that does not undergo a first etching or another fracture process on the structure. And do not impregnate. In some embodiments of the systems and methods disclosed herein, the matrix phase is not a coating such as a coating deposited by ALD, sputtering, or electrophoretic deposition. In some embodiments of the systems and methods disclosed herein, the matrix phase is not an electrolyte, such as an electrochemical cell electrolyte, including a solid electrolyte.

"Vibration frequency bandgap" refers to the frequency or frequency range corresponding to the vibration (or vibration) of a structure, composite material system, or structure, and the magnitude or energy of vibration in the frequency or frequency range is at least 10. Double (1 digit), at least 20 times, at least 50 times, preferably at least 100 times (2 digits), preferably at least 200 times for some applications, more preferably at least 500 times for some applications. Yes, it is less than the magnitude or energy of vibration at frequencies outside the "vibration frequency band gap". In some embodiments, the vibration frequency bandgap can be characterized by a midpoint frequency and / or frequency width. In one embodiment, the partial vibration frequency bandgap exists along one or more directions (eg, X, Y, Z, or any direction or vector in between), but along all directions. Does not vibrate frequency bandgap. In one embodiment, the complete vibration frequency bandgap is a vibration frequency bandgap that exists along all directions (eg, X, Y, Z, or any direction or vector in between).

The term "cross-sectional physical dimensions" refers to the physical dimensions of a feature measured on the horizontal or cross-sectional axis. In one embodiment, the horizontal axis is perpendicular to the vertical axis of the feature. In one embodiment, the physical dimensions of the cross section correspond to the width or diameter of a feature such as a beam, strut, or tie. In one embodiment, the longitudinal physical dimensions are the dimensions of the feature along the vertical axis of the feature, the vertical axis being perpendicular to the cross-sectional axis. If necessary, the longitudinal physical dimensions are measured between the two nodes. If necessary, longitudinal physical dimensions are measured between the physical edges of the structure.

The term "unit cell" refers to the minimal arrangement, composition, or geometry of a plurality of features, so that the entire structure or its three-dimensional geometry characterized by the unit cell is formed by repeating the unit. Can be done. For example, repeating unit cells in three dimensions can form three-dimensional structure. The entire structure can be a three-dimensional structure such as a three-dimensional porous structure.

"Young's modulus" is a mechanical property of a material, device, or layer that refers to the ratio of stress to strain of a given substance. Young's modulus is given by the following equation.

Here, E is Young's modulus, L ₀ is the equilibrium length, ΔL is the change in length under load stress, F is the loaded force, and A is the area to which the force is loaded. Young's modulus can also be expressed by Lame-parameters using the following equation.

Here, λ and μ are Lame-parameters. Young's modulus can be measured according to methods conventionally known or not yet known in the art. For example, Young's modulus is at Roylance ("Stress-Strain Curves", MIT Course, August 23, 2001, https://web.mit.edu/course/3/3.11/www/modules/ss.pdf). Corresponds to the slope of the linear portion of the stress-strain curve described in (accessed at application).

The term "average", when used with respect to material or structural properties, refers to the calculated arithmetic mean of at least two, or preferably at least three, identical measurements or calculations of the above properties. For example, the average density of structures is the arithmetic mean of at least two similarly performed measurements of the density of the above structures.

The term "density" refers to volumetric mass density. Density is expressed in units of mass per volume (eg, g / cm ³ ). When referring to a material, the term density corresponds to the volume-mass density of the material. When referring to a structure, the term density corresponds to the volume-mass density of the structure, which is a function of the geometry of the structure and also the function of the material from which the structure is formed, such a structure. An increase in porosity corresponds to a decrease in the density of the structure. The density of structures, such as structures with three-dimensional geometry, according to one embodiment of the present invention can be measured according to methods conventionally known or not yet known in the art. For example, the density of the structure can be determined for a disc-shaped sample by determining the mass, height, and diameter and dividing the determined mass by the volume of the sample.

The term "relative density" refers to the volume fraction of a solid material in a composite system, structure, or feature. In one embodiment, the relative density corresponds to the ratio of the density of the structure to the density of the solid material (or combination of materials) that makes up the structure. Relative density can be expressed as a percentage (ratio of densities) or a percentage (ratio of densities x 100%). In one embodiment, the relative density of the pre-pyrolysis structure or its three-dimensional geometry is substantially the same as that after pyrolysis.

The term "specific strength" refers to the ratio of strength to density of a material, system, structure, or feature, where strength refers to the force per unit area at the point of failure of a material, element, or structure. Specific strength is sometimes referred to as strength-to-weight ratio. In one embodiment, "strength" refers to compressive strength. In one embodiment, "strength" refers to tensile strength. In one embodiment, the compressive strength is the maximum stress that the material can withstand under crushing loads. In one embodiment, the compressive strength of a material, structure, or element that is broken by crushing can be defined as an independent property within a fairly narrow range. In one embodiment, the compressive strength of a material, structure, or element that does not grind during compression is the amount of stress required to deform the material to any amount. In one embodiment, the compressive strength of a material, structure, system, feature, or element that is not crushed by compression can be calculated as stress at 0.2% strain offset from the linear portion of the stress-strain curve. In one embodiment, compressive strength is calculated by dividing the maximum load on a material, structure, or element by the original cross-sectional area of the material, structure, or element being inspected.

The term "stiffness" refers to the degree to which a material, structure, system, or feature resists deformation in response to applied force. Rigidity is the displacement generated by the force applied to a material, structure, or element and the force applied along the same degree of freedom (eg, the same axis or direction) indicated by the material, structure, or element. Corresponds to the ratio. The term "specific stiffness" refers to the ratio of stiffness to the density of a material, element, or structure. In one embodiment, the stiffness of a material, structure, or element is Young's modulus of the material, structure, or element.

According to certain embodiments, the structure has a nodeless (ie, nodeless) geometry. The nodeless geometry has extraordinary mechanical restoring force. The mechanical restoring force can be understood, for example, from the viewpoint of strain against fracture (strain vs. fracture) and strength against fracture (strength vs. fracture). In one embodiment, the breaking strength of a material, element, or structure corresponds to the compressive strength of the material, element, or structure. In one embodiment, the structures of the invention have a fracture strain of 2% to 5% and optionally 2.9% to 3.5%. The strain to fracture can be determined according to methods conventionally known or not yet known in the art. For example, strain vs. fracture can be determined from the strain value corresponding to a linear portion such as the third linear portion of the stress vs. strain data up to the sudden stress loss (fracture) of the structure.

The term "laminated form" refers to the process of forming a structure or feature by depositing or stacking materials. The terms "additive manufacturing process" and "additive manufacturing process" can be used interchangeably. The laminating process may involve layer-by-layer (layer-by-layer) deposition of material to form complex three-dimensional structures or elements. The deposited material may include, but is not limited to, inorganic materials, organic-inorganic composites, polymers, metals, or combinations thereof. Exemplary layered modeling processes include, but are not limited to, 3D printing, stereolithography (SLA), fused deposition modeling (FDM), and two-photon lithography. In some embodiments, the laminating process does not require a removal process to form a structure or element. Examples of removal machining processes include, but are not limited to, cutting, machining, electrical discharge machining, engraving, molding, grinding, drilling, and etching. In one embodiment, the layered build process includes or is supported by computer-aided design (CAD).

In one embodiment, the term "defect" refers to imperfections or unintended features due to manufacturing, such as local deformations, cracks, beam junction offsets, beam bulges, struts, and pits or voids. Refers to a characteristic.

The term "node" may refer to a junction or intersection of multiple features such as a beam or strut. The structure can have three-dimensional geometry, which is a nodeless geometry.

According to one embodiment, the term "core", when referring to a feature of a structure having three-dimensional geometry, refers to the internal volume of the feature, including the outer surface of the feature and excluding the outer surface of the feature. In one embodiment, the feature core corresponds to the internal volume of the feature except that of any coating, especially the coating introduced after the pyrolysis process present on it.

The term "prepolymer" refers to a monomer or a mixture containing one or more monomers in which the monomer has reacted to a medium molecular weight state. The prepolymer can be further polymerized to a fully cured high molecular weight state. In some embodiments, the terms prepolymer and monomer can be used interchangeably.

As used herein, the term "polymer" is often referred to as a significant number of repeat units (eg, 3 or more repeat units, and in some embodiments, 10 or more repeat units, a number). Repeated structures linked by covalent chemical bonds characterized by 30 or more repeat units in the embodiment) and high molecular weight (eg, 10,000 Da or more, in some embodiments 50,000 Da or more, equal to 100,000 Da). Refers to a molecule consisting of units. A polymer is usually a polymerization product of one or more monomer precursors. The term polymer includes homopolymers, or polymers consisting essentially of a single repeating monomer subunit. The term polymer also includes copolymers formed when two or more different types of monomers are bonded into the same polymer. Copolymers can contain two or more monomer subunits and can include random, block, brush, brush block, alternating, segmenting, grafting, tapering and other structures. Useful polymers include organic or inorganic polymers that can be amorphous, semi-amorphous, crystalline or semi-crystalline. Polymer side chains capable of cross-linking the polymer (eg, physical cross-linking) can be useful in several applications.

The term "substantially" refers to a property within 10%, within 5%, within 1%, or equivalent to a reference property. The terms "substantially equal," "substantially equivalent," or "substantially unchanged" when used in combination with a reference value that describes a characteristic or condition, with respect to a given reference value. It means that it is within 10%, arbitrarily within 5%, optionally within 1%, optionally within 0.1%, or arbitrarily equal. For example, if the value of the ratio is within 10%, optionally within 5%, optionally within 1%, or optionally equal to 1, the ratio is substantially equal to 1. The term "substantially greater", when used in combination with a reference value that describes a property or condition, is at least 2%, optionally at least 5%, or optionally at least 10% greater than the provided reference value. Point to a value. The term "substantially less" is at least 2%, optionally at least 5%, or optionally at least 10% lower than the provided reference value when used in combination with a reference value that describes a property or condition. Point to a value.

In one embodiment, the compositions or compounds of the invention, such as alloys or precursors of alloys, have been isolated or substantially purified. In one embodiment, as is understood in the art, the isolated or purified compound is at least partially isolated or substantially purified. In one embodiment, the substantially purified compositions, compounds or formulations of the invention are in some cases 99% for some uses, in some cases 99.9% for some uses, and in some cases for some uses. It has a chemical purity of 99.99% and, in some cases, 99.999 for some applications.

In one embodiment, the term "reduced energy" refers to energy that is redirected from a composite system, structure, feature, or material and does not cause failure of the composite system, structure, feature, or material (eg, pre-collision). Particle energy + particle energy after collision only if the velocity vector is different from the first). In one embodiment, the term "collision energy" refers to the energy of the collider before a collision. In one embodiment, the term "absorbed energy" refers to the difference between the collision energy and the repulsion energy of a collision energy (eg, a particle). In one embodiment, the term "specific energy absorption" refers to the ratio of strain energy density (defined as W, W = ∫σdε) to material density (ρ). In one embodiment, the term "specific energy absorption" refers to the ratio of strain energy density (defined as W, W = ∫σdε) to composite system, structure, material, or feature density (strain).

The following description provides many specific details of the devices, device components and methods of the invention to provide a complete description of the exact nature of the invention. However, it will be apparent to those skilled in the art that the present invention can be practiced without these particular details.

In one embodiment, the composite system has at least one monolithic structure (or "reinforced phase") with three-dimensional geometry and the centerline of the truss element (as opposed to the waveguide process) is of material. It does not extend from the edge to the whole and can start and end at any point in the material. Similarly, the centerline of the truss element can be placed in any direction within the material, including perpendicular to the thickness direction, as opposed to a waveguide process where this is not possible. Similarly, a particular truss element can have any cross section and can vary throughout the truss element. Finally, the centerline of a particular truss element can have a non-zero curvature. One or more matrix phases fill the volume around the strengthening phase. The matrix phase can consist of different material classes including, but not limited to, polymers, ceramics, carbons, and metals.

In one embodiment, the composite system has a continuously reinforced phase with a three-dimensional shell or surface geometry with negative, zero, or positive Gaussian curvature. The walls or membranes of the shell geometry can have varying thicknesses throughout the material. The surface geometry can follow a closed cavity separated from the outer matrix phase. One or more matrix phases fill the volume around the strengthening phase. Reinforced phases can be composed of different material classes, including but not limited to polymers, ceramics, carbons, and metals.

In one embodiment, modular three-dimensional structural elements of any shape are made of the composite material described in the previous embodiment. The reinforced phase of the composite will have the geometry of the final structural component as a continuous phase. The structural component topologies can have zero or multiple holes (ie, monolithic composite or tubular components, respectively). The resulting holes can be impregnated with different matrix phases or left unchanged.

In one embodiment, structural components of any shape are made of the composite material described in the previous embodiment and have a functionally inclined geometry of one or more phases. The continuous lattice structure or surface of the strengthening phase can vary in the thickness direction and in-plane while maintaining continuity. The cross section and thickness can also be changed without affecting the continuity of the phase.

In one embodiment, the microstructure of the continuously reinforced phase of the composite provided in the previous embodiment can function as a resonator and have the characteristic of providing damping to the material. The cavity is surrounded or isolated by the matrix phase.

The method of making a three-dimensional composite system with an arbitrary architecture includes the step of designing an arbitrary structure (which can be periodic) by a computer-aided design (CAD) tool and the step of selecting a desired precursor resin. Includes the step of exposing the resin to the desired layer-by-layer pattern features of a layered modeling technique, including but not limited to SLA and DLP. If necessary, additional resin is then removed and the sample is post-cured by UV and heat treatment, followed by a pyrolysis process in a specified temperature profile and controlled environment. The structure is then impregnated with one or more materials and a vacuum process and a sonication process are utilized to finally form a composite material having a continuous phase of any shape.

The composite system disclosed herein provides an improvement over typical carbon fiber composites in that weak interlayer interfaces are eliminated, resulting in excellent material response under bending and compression. The fully interconnected reinforcement phases can provide advantages for collision absorption applications where the in-plane properties of thin materials can determine the degree of damage. Moreover, the above method provides a clear advantage in the manufacture of structural parts by producing the strengthening phase in the final desired geometry instead of avoiding the molding process.

The following is a description of exemplary, exemplary, and embodiments of composite material systems and methods disclosed herein.

The fabrication process embodied in FIG. 1 begins with the design of one or more geometries that act as the reinforcing phase of the composite. The precursor resin can determine the constitutive properties of the resulting reinforced phase, but the architecture determines the structural response (ie, stretching or bending dominates). The structure is manufactured using laminated modeling techniques such as stereolithography (SLA) and digital optical processing (DLP) 3D printing, and any continuous shape can be made. The printed structure can then be placed in the furnace and the pyrolysis process can be carried out in a vacuum or an inert atmosphere. The resulting reinforced phase can be a carbon or ceramic material. Any coating process can be performed prior to impregnating the resulting reinforced phase with the selected matrix phase. During the impregnation step, additional processes such as vacuum degassing and sonication help ensure complete impregnation of the matrix phase. Finally, the resulting composite can be post-cured by UV or heating processes.

In the embodiment presented in FIG. 2A, a subset of the carbon-enhanced phases 10 produced by the above process is embedded in the epoxy matrix 16. In this case, the stretch-dominant architecture was chosen because of its high stiffness-to-density ratio and the achievable bending-dominated architecture depending on the polymer waveguide pattern. A given truss element of this structure also has a variable cross section, strengthening the grid nodes 12 and increasing rigidity as compared to a constant cross section structure. In this particular geometry, the truss element 14 can start and end at any point in the material, not necessarily at the edges, as required by the waveguide process. A variant of this embodiment is shown in FIG. 2B, where a given truss element is allowed to have a non-zero curvature of 18, and the truss element can be oriented perpendicular to the thickness direction 20.

The embodiment presented in FIG. 3 shows a subset of continuous 3D carbon phases made of shells or surfaces 22 impregnated with epoxy matrix phase 26. The surface of the strengthening phase can have a non-zero Gaussian curvature that cannot be achieved by the waveguide process. This geometry was achieved using the surface of an octet truss structure. In this geometry, isolation regions 28 can be created to form the gas phase in order to prevent impregnation from the matrix phase 26. Finally, the surface thickness of this structure is designed to be indefinite (non-constant) 24, which is not possible in the waveguide process.

Another embodiment shown in FIG. 4 shows a modular structural element made of a composite as described in a previous embodiment. Unlike typical fiber-based composites, the reinforced phase 30 is designed to have the same shape as structural components (eg, tubing) and does not require an additional molding process. In this case, a bend-dominated architecture was chosen to add adaptability and increase the energy absorption capacity of the part. The structural component has a primary matrix phase 32, but the internal volume of the tube can be left empty or filled with a secondary matrix phase 34.

5A-5B present an embodiment in which the functional tilt of the structural component (made from a composite material as described in the previous embodiment) is achieved. The structural part has one continuous reinforcing phase with several different geometries 36, 40, 42 and 46. A bending-dominated tetrahedral decahedron architecture is used in the upper region where the cross-sectional area of the elliptical truss element increases from 36 to 40, 42. The domain transition 38 from region 36 to region 40 is continuous and the connectivity has not changed (ie, the truss element does not abruptly terminate as in the case of some functional gradients made through the waveguide process). The lower region 46 consists of a stretch-dominated octet truss structure with circular cross-section truss elements and a continuous transition 44 to the upper region.

In one embodiment, as shown in FIG. 6, the design of the cavity within the phase microstructure creates a rigid damping composite. The unit cell 48 has an octet truss structure having a resonator 50 inside the unit cell. The resonator consists of a cantilever with microinertia added to the free end and is tuned to resonate at a predetermined frequency ω determined by the length of the beam and the amount of microinertia at the free end. The unit cell can be formed by a combination of truss elements and shells, which can separate the resonator from the primary matrix impregnating the reinforcing unit cell 48. If a shell is present, the internal volume of the octahedron in the octet 52 is left empty to allow free vibration of the resonator, otherwise the damping properties of the hard material are possible.

The embodiments implemented are shown in FIGS. 7A-7C. In this embodiment, a sheet having a three-dimensional continuous octet truss structure 54 having an overall size of 12.5 × 4 × 0.13 cm is produced by DLP3D printing using an Autodesk Ember printer and PR-48 resin. The unit cell size of the resin structure 56 is ~ 1.3 mm and has a circular cross-section truss element. The resin structure is then thermally decomposed at a peak temperature of 1000 ° C. in a furnace in an inert atmosphere. The resulting carbon structure 58 undergoes isotropic shrinkage to 40% of its original volume with only a 7% reduction in its original mass. The resulting carbon unit cell 60 retains its original octet truss geometry and a characteristic size of around 500 μm. The carbon structure is then impregnated with the epoxy matrix phase 62 to become the reinforced phase 64 of the composite structure sheet.

FIG. 8 shows one possible method of producing the reinforced phase of the composite described in the previous embodiment by DLP or SLA 3D printing. The printer head 66 is coupled to the edge of the strengthening phase and pulls the structure vertically during the printing process. The print screen 70 has a layer of uncured precursor resin that polymerizes when any exposure pattern 72 is projected from below. Note that this allows truss elements 68 with arbitrary cross sections starting and ending from the edges of the structure, in addition to the truss elements perpendicular to the construction direction.

In general, specific exemplary embodiments of composite material systems with three-dimensional, arbitrarily designed, fully interconnected phases, and methods of making them are presented. A particular embodiment will be described in which the strengthening phases are fully interconnected, the cross section changes, and the truss elements are composed of non-zero curvature (impregnated with a continuous matrix phase). These composite systems allow the production of interconnected 3D reinforced phases with shells of non-zero curvature and varying thickness. In some embodiments, modular structural parts made of the above composites with functionally graded continuous phases are presented. In addition, certain embodiments present enhanced vibration damping in composites through the design of resonators within the microstructure without the need for a dissipative viscoelastic phase that affects the stiffness of the material.

Literature corresponding to background and explanation

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Reference 7: RJ Grylls and CM Austin, “Article made of ceramic foam joined to a metallic nonfoam, and its preparation”, Patent US 6 582 812B1, 04 24, 2003. Available online: https://patents.google. com / patent / US6582812B1 /

Reference 8: G. Rettenbacher, J. Reiter, F. Feuchtenschlager, P. Schulz, and P.-F. Louvigne, “Multilayer composite armor”, Patent US 70 226 045B2, 0411, 2006. Available online: https: //patent.google.com/patent/US70226045B2/

Reference 9: J. Zeschky, F. Goetz-Neunhoeffer, J. Neubauer, SH Lo, B. Kummer, M. Scheffler, and P. Greil, “Preceramic polymer derived cellular ceramics”, Composites Science and Technology, vol. 63, no. 16, pp. 2361-2370, 2003.

Reference 10: M. Tehrani, AY Boroujeni, TB Hartman, TP Haugh, SW Case, and MS Al-Haik, “Mechanical characterization and impact damage assessment of a woven carbon fiber reinforced carbon nanotube-epoxy composite”, Composites Science and Technology, vol. 75, pp. 42-48, 2013. Available online: http://dx.doi.org/10.1016/j.compscitech.2012.12.005

Reference 11: SS Wicks, RG de Villoria, and BL Wardle, “Interlaminar and intralaminar reinforcement of composite laminates with aligned carbon nanotubes,” Composites Science and Technology, vol. 70, no. 1, pp. 20-28, 2010. Online Available at: http://dx.doi.org/10.1016/j.compscitech.2009.09.001

Reference 12: AJ Jacobsen, WB Barvosa-Carter, AF Gross, R. Cumberland, KW Kirby, and D. Kisailus, “Composite structures with ordered three-dimensional (3d) continuous interpenetrating phases,” Patent US 8 320 727B1, 1127, 2012. Available online: https://patents.googie.com/patent/US8320727B1/

Reference 13: SS Yang and AJ Jacobsen, “Micro-truss materials having in-plane material property variations,” Patent US 9 405 067B2, 08 02, 2016. Available online: https://patent.google.com/patent / US9405067B2 /

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Reference 15: TA Schaedler, AJ Jacobsen, W. Carter, and G. McKnight, “Constrained microlayer cellular material with high stiffness and damping”, Patent US 9 217 084B2, 12 22, 2015. Available online: https: // patent.google.com/patent/US9217084B2/

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Reference 17: J. Bauer, A. Schroer, R. Schwaiger, and O. Kraft, “Approaching theoretical strength in glassy carbon nanolattices”, Nature Materials, vol. 8, no. February, 2016.

Reference 18: X. Chen, G. Zhao, Y. Wu, Y. Huang, Y. Liu, J. He, L. Wang, Q. Lian, and D. Li, “Cellular carbon microstructures developed by using stereolithography,” Carbon, vol. 123, pp. 34-44, 2017. Available online: http://dx.doi.org/10.1016/j.compscitech.2017.07.043

Reference 19: Q. Zhang, F. Zhang, S. P. Medarametla, H. Li, C. Zhou, and D. Lin, “Three-Dimensional Printing of Graphene Aerogel,” Small, vol. 12, no. January, 2016.

The present invention can be further understood by the following non-limiting examples.

Example 1: Collision response of 3D carbon architecture

The manufacture and design of materials with high stiffness-to-density or strength-to-density ratios have been studied through the use of bubble materials. In particular, the beam-based lattice architecture has made it possible to fabricate lightweight but highly rigid materials (References 1 and 2) and strong materials approaching theoretical limits (Reference 3). Many of these studies have focused on the static response of these materials, and few have studied the dynamic response of lattice structures. In particular, some studies have studied the dynamic compression of the lattice structure on the μm scale (Reference 4), and another study has studied the effects of macroscale structures such as lattice core sandwich plates (Reference 5). ). The length scales and tessellations associated with these studies do not provide proper separation of scales where the length scales of the boundary conditions are much larger than those of the inherent microstructure. Such separation allows the true material properties to be investigated as opposed to the properties of the individual structures.

In this example, the supersonic collision response of a carbon lattice structure (ie, the form of a reinforced phase in the composite of the present invention) will be described while maintaining proper separation of scales. Two-photon lithography manufacturing processes are used to create three-dimensional lattice structures with nanometer-scale capabilities of various unit cell geometries and relative densities (8-26%) at supersonic speeds (500-1100 m / s). Observe the marginal damage after the collision at. These results show how the lightweight design phase (before impregnation) can provide great resilience to collisions.

1.1 Sample design and production

Since the architecture leads to different mechanical properties in the static state (References 6 and 7), the purpose was to investigate the effect of the architecture on the supersonic collision of the lattice architecture. Polymer grids of rigid octets and non-rigid tetradecahedron unit cells were manufactured using a two-photon lithography process (Nanoscribe) with unit cell sizes in the range of 5-10 μm (see Figures 10A-10B). .. In addition, the beam radii were modified to achieve three different relative density samples of 8%, 17% and 26%.

A sufficiently large tessellation (approximately 60 x 60 x 15 unit cells) was selected so that the effective sample size was much larger than the size of the unit cell and the grid could be approximated as the effective material. The polymer sample was then pyrolyzed to 900 ° C. in vacuum to give a monolithic carbon lattice with 80% isotropic shrinkage while retaining the original geometry (FIGS. 11A-11D). reference).

For the smallest initial unit cell (ie, 5 μm), using a beam diameter reduced to about 200 nm, the resulting carbon unit cell had submicron dimensions. The geometry of the final unit cell corresponds to the shape of the original polymer, although there is a small warp during the pyrolysis.

1.2 Crash test

The resulting carbon structure (ie, the strengthening phase) was exposed to supersonic collisions by accelerating SiO ₂ particles with diameters in the range of 7-14 μm. In all cases, the particle size was at least an order of magnitude larger than the characteristic unit cell size. The method adopted is defined as Laser-Induced Particle Collision Testing (LIPIT) (References 8 and 9), allowing controllable collision velocities of up to 1 km / sec while acquiring the collision process with a high speed camera. FIG. 12 shows a characteristic experiment of an octet grid with a relative density of 26% and a collision velocity of 1060 m / sec.

For all grids tested, projectile bounces, collision velocities and bounce velocities were measured. Due to the low adhesion between the grid and the substrate, the sample delaminated within a few milliseconds after collision, requiring modification of the sample to allow post-fracture characterization. To effectively connect the sample to the substrate, a thin layer of PMMA resist was applied onto the substrate, resulting in a coating of hundreds of nanometers that bonded the grid to the substrate (see FIGS. 13A-13C).

1.3 Theoretical elastic wave velocity

Simplification that planar elastic waves are emitted from the collision site through the thickness of the lattice, as shown in FIG. 14A, to ensure that the observed behavior is substrate independent (ie, independent of substrate stiffness or thickness). Suggest the problem that was raised. It is assumed that the effective Young's modulus of a non-elongated octet lattice is given by the following equation (1) (Reference 7).

Here, Es is the Young's modulus of the constituent material, and r / l is the ratio of the radius to the length of the strut. In this case, the elastic wave velocity can be approximated by the following equation (2).

Here, ρ is the effective lattice density. Using worst-case values such as r / l = 0.2, ρ = 1252 kg / m ³ (corresponding to carbon octets with a relative density of 60%) and a sample thickness of 14 μm, elastic waves pass through the sample twice. It takes about 12 nm to (that is, reciprocate). Since the high-speed camera frame can measure the approximate collision time, the average collision time of 4ns means that the information about the substrate is not transmitted to the particles before bouncing. In other words, the bounce behavior is simply a function of the grid material, not the substrate. This is summarized in the xx diagram shown in FIG. 14B.

To experimentally test this claim, we designed a floating sample experiment in which the same grid was mounted on a Si strut a few microns away from the substrate. This experiment mimics a macroscale drop tower collision experiment in which a plate with fixed boundary conditions is affected by accelerated mass (see Figures 15A-15D). In this way, the grid was completely separated from the previously placed Si substrate.

When the same experiment was performed on a floating sample at a collision speed of 588 m / sec, the same repulsion behavior and a bounce speed of 320 m / sec were obtained (see FIG. 16). This result confirmed that the observed response was substrate independent.

1.4 Energy absorption scaling behavior

Crash tests on samples of both rigid and non-rigid architectures at different relative densities gave the results shown in FIG. Experiments at low collision energy (i.e., mv 0 ^_2, where v ₀ is the impact velocity), the highest coefficient of restitution is obtained (i.e., coefficient of restitution is the ratio of the rebound velocity relative collision velocity v _{r /} v ₀ ).

Rigid and non-rigid architectures may have up to an order of magnitude difference in elastic modulus in the static state (Reference 7), but no definitive difference was observed under supersonic collision conditions. Octet and tetrahedral dodecahedron samples had similar bounce coefficients across the investigated kinetic energy states.

The trends observed in FIG. 17 appear to show the correlation between the collision energy and the bounce coefficient, independent of the architecture. In experiments with high collision energies and different relative densities, it seemed that the repulsive force decreased as the relative density decreased. This is probably due to the lower effective material strength and more energy consumed by local damage. Despite this slight local damage at high collision energies, these results show the advantages of using an architecture for collision mitigation. This is because energy can be more easily dispersed from the collision site and local damage can be reduced. If the collision process does not allow the time between the arrival of elastic waves at the projectile and the occurrence of bounces, the process is likely to be inertial rather than architectural, with careful design and material distribution. Control is still needed.

Additional Note: Since the matrix phase corresponds to additional inertia (ie, mass), the matrix phase can increase energy dissipation or mitigation. With respect to damping, viscoelastic matrices such as polymers can further dampen vibration and collision energies. The matrix also helps prevent cracks from opening / propagating, thus increasing the strength of the material. The specific values of these increases can be highly dependent on the architecture and material choices.

One example is a "coated sample" (covered with a thin layer of epoxy), which corresponds to a composite system having a structure partially impregnated with a matrix phase (such as epoxy). Referring to FIGS. 13B and 17, the addition of a thin polymer layer reduced the bounce coefficient to about 0.65 (35% reduction) of the bounce coefficient of the uncoated sample. This translates to 58% less kinetic energy of the particles during repulsion. The energy relaxation criteria do not change because the particles bounce and maintain structural integrity, but are converted to dissipation / absorption (ie, thermal or permanent deformation) compared to a sample without a matrix (absorbed energy). ) Energy increased by 20% (Kyushu energy was 66% of collision energy by uncoated sample, but 86% was absorbed by coated sample). Since this was a thin coating, it could be the lower limit of increased absorption when adding a matrix in one embodiment. Note that the absorbed energy corresponds to the difference between the collision energy and the kinetic energy of the collider after rebound.

In general, inclusion of the matrix reduces the bounce coefficient and increases energy absorption. This means that the viscoelastic / plastic properties of the matrix absorb some of the energy so that less energy is returned to the collider when moving in the opposite direction. Having a matrix improves all damping characteristics compared to structures without a matrix phase. For example, the matrix can increase the vibration frequency bandgap width or reduce the transmission intensity of vibrations at some frequencies. From a static point of view, in the presence of a matrix phase, the strength of the material is significantly increased compared to a structure without a matrix phase, and fracture can be catastrophic / brittle to ductile.

Example 2: Material attenuation due to architecture

Careful design of the constituent materials can result in interesting dissipation behavior, which can lead to energy dissipation. Studies have shown large 3D-printed effective materials that can attenuate vibrations using dispersion mechanisms such as Bragg scattering and local resonance (Reference 10) and microscale materials that dissipate ultrasonic waves in water (Reference 10). 11) is shown. Proper design of the composition, taking into account the density and rigidity of the material, leads to effective material attenuation by the rigid, non-dissipative constituent material.

In this embodiment, a careful design of the configuration is utilized in order to dissipate the vibrational energy in a rigid non-attenuating constituent material by utilizing the Bragg scattering and local resonance mechanism.

2.1 Bragg scattering: square lattice

Starting with a square unit cell such as that shown in FIG. 18A, the volume expansion of that member changes the unit cell to that shown in FIGS. 18B-18C, depending on the degree of volume expansion, buckling instability. Will bring.

The introduction of curvature into the beam not only changes the shape of the unit cell, but also its effective mechanical properties, especially in the in-plane direction. The effect of this buckling geometry on the dynamic properties of the material composed of these square unit cells will be investigated numerically.

For this numerical study, it was assumed that each unit cell had a polymer core with a Si coating. The horizontal beam polymer had a minor radius of 0.25 μm, the main radius was 0.9 μm, the Si coating was 0.4 μm, and the vertical beam polymer radius was 0.9 μm, with the same coating. The dimensions of the original square unit cell were 20 x 20 x 5 μm.

The natural frequency analysis of the 3D unit cell at each stage of buckling was performed using the commercially available finite element package COMSOL Multiphysics. Each unit cell was divided into a horizontal beam region and a vertical beam region, each containing an elastic material model of the corresponding homogenized beam properties. The homogenized properties were obtained using the weighted volume average from the known volume of each material (ie, polymer and silicon), the corresponding Young's modulus, density, and Poisson's ratio of each material.

Bloch boundary conditions were applied to the corresponding faces of the unit cells using the square unit cells shown in FIGS. 18A-18F. Using the irreducible Brillouin zone shown in FIG. 19D, the fractional vector passed through the boundary and the intrinsic frequency was calculated for each state. Tetrahedron elements were used and mesh convergence was confirmed in all cases.

Figures 19A-19D show the appearance of partial bandgap in the x and y directions (due to symmetry), corresponding to frequencies that cannot propagate throughout the material consisting of the unit cells. This is due to the Bragg scattering made possible by the buckled geometry, assuming that the bandgap does not appear in the case of non-buckling in FIGS. 19A. The bandgap width of the vibration frequency can change in proportion to the feature length of the architecture. Specifically, the Bragg condition is remarkable at the frequency of ΔL = c / f (ΔL is the characteristic dimension of the microstructure, c is the speed of sound in the material, f is the frequency, and shows a linear relationship with the frequency). Indicates that the impact may occur.

These results show the adjustability of the designed structure, which allows for a dispersion mechanism that can lead to damping. The absolute value and width of the bandgap frequency obtained can be adjusted by changing the size of the unit cell and the constituent materials. The results shown here are obtained with a fully elastic material model, which means that the same behavior can be achieved with hard materials that do not decay in nature, such as metals, ceramics, and carbon.

2.2 Local resonance: Authentic design with resonator

In addition to Bragg scattering, local resonance can generally be used to allow bandgap at lower frequencies. In this study, the auxetic unit cell presented by Krodel et al. Is used (Reference 12). As shown in FIGS. 20A-20B, a resonator (ie, a concentrated mass attached to the cantilever) is added to the unit cell.

We performed numerical studies as was done in Section 2.1, assuming a fully polymer unit cell with dimensions of 60 x 60 x 210 μm. An elastic material model was used (ie, no damping of the constituent materials was assumed) to calculate the dispersion relation in the Γ-X direction (see Figure 19D).

The dispersion relation of the unmodified auxetic unit cell (see FIG. 21A) indicates that there is no bandgap in the direction of interest, while adding a resonator (see FIG. 21B) creates a bandgap at 1.5 MHz. be introduced. Similar to Section 2.1, the width and position of this bandgap is fully adjustable based on unit cell dimensions, materials, and resonator parameters.

To experimentally verify the numerical results from FIGS. 21A-2B. Samples of the same dimensions, materials, and parameters were made and tested in a custom vibration transfer experiment in vacuum (see Figures 22A-22C).

Using the configuration presented above, a continuous sine wave with frequencies varying between 1 and 3 MHz was transmitted through the sample (see FIGS. 23A-23E).

When a frequency sweep was performed on the resonator unit cell, as shown in FIG. 23C, a band gap centered on a frequency of about 2.4 MHz was shown. The bandgap occurred at frequencies slightly higher than those predicted by the simulations in FIGS. 21A-21B. This is because the grid was distorted by establishing proper contact with the transducer, which can change the dispersion behavior (see Distortion Difference Between 23D-23E).

To further examine this bandgap, additional transmission experiments were performed in which chirps were transmitted through the grid (instead of continuous waves). In this case, the chirp contained frequencies from 1 to 3 MHz, and a Fast Fourier Transform (FFT) was applied to the transmitted signal to analyze its frequency content (see Figures 24A-24C).

These experiments and simulations on polymer auxetic lattices show the possibility of adding local resonance as a mechanism to introduce damping into the material. Since the material properties of the simulation are perfectly linear elastic, this behavior can be extended to various materials such as metals, ceramics and carbon.

Example 3: Fully adjustable elasticity with spinodal decomposition derived architecture

Design materials with beam-based architectures have been shown to be effective in achieving high stiffness-to-density ratios (References 1 and 6), but they are still far from theoretical limits. Furthermore, their mechanical properties deviate from theoretical predictions due to the presence of nodes that also function as stress concentrations that can cause failures.

This example illustrates the use of spinodal decomposition to create shell-based microstructures that lack nodes and achieve better mechanical properties and higher mechanical adjustability than beam-based architectures.

3.1 Elastic surface adjustable

Using the spectral method and anisotropic energy functionals (13), numerical spinodal decomposition can result in microstructures with fully adjustable elasticity (for exemplary methods to illustrate computational spinodal decomposition, see A. Vidyasagar, S. Krodel, and DM Kochmann, "Microstructure Patterns with Adjustable Mechanical Anisotropy Obtained by Simulating Anisotropy Spinodal Decomposition," Science, vol. 474, no. 2218, p.20180535, 2018). Using only the boundaries of the resulting microstructure (ie, the shell), a 3D elastic surface was calculated using the commercially available finite element code Abaqus, demonstrating the complete adjustability of the effective Young's modulus of the microstructure. (See FIGS. 25A-25D).

The elastic surfaces shown in FIGS. 25A-25D show unparalleled elastic adjustment capabilities not achievable with commonly studied beam-based architectures, as shown by numerical simulations of these structures. In addition to this highly adjustable behavior, spinodal structures can approach theoretical limits and exhibit a significantly greater Young's modulus than other structures such as trusses and 3-period minimal surfaces (References 14-16). See FIG. 26).

By applying periodic boundary conditions to columnar structures, P-cell minimum surfaces, and hollow octet trusses with equal relative densities, spinodal structures come closer and have better modulus when explored in the z direction. It shows and is approaching the focal boundary. further,

Format (E ^* is the effective Young's modulus, the Young's modulus of E _s of the constituent material, [rho is the relative density ([rho is [rho bar), a scaling exponent) Applying fitting, the lower (more desirable) spinodal structure Produces a scaling index.

As mentioned above, one obvious advantage of spinodal structures is the lack of those nodes, which reduces the stress concentration at which cracks can begin. This results in a surface with a quasi-constant low curvature, as opposed to a truss with an infinite curvature at the nodes (see FIGS. 27A-27B).

The curvature distribution of the octet truss above shows a bimodal distribution even when fillets are applied to the nodes. This limits the maximum curvature of the structure, but the absolute value is much larger than the maximum curvature of the spinodal structure.

It should be mentioned that the effective modulus of the structure, geometry, and / or system disclosed herein is closer to the theoretical limit (foic boundary) than the effective modulus of a typical truss or transverse fiber (foil boundary). That is, it approaches the black foil line in FIG. 26). With respect to elastic adjustability, the composite systems disclosed herein, as having a spinodal-shaped structure, have deterministic anisotropy (or) of elasticity, damping, collision energy absorption, and / or other properties or features. It should also be mentioned that it is characterized by isotropic) if necessary.

3.2 Manufacture of shell-based spinodal materials

We used a two-photon lithography process to create a polymeric spinodal structure on a microscale. Metals or ceramics were deposited from 5 nm to 5 μm using deposition techniques such as atomic layer deposition (ALD) and magnetron sputtering. In this case, the resulting material is a composite material with a shell-based reinforcement phase. Alternatively, the polymer core can be removed and left empty or replaced with another matrix. Focused ion beam (FIB) milling was used to remove small sections of the coating to expose the polymer under the newly applied coating. Finally, introducing the structure into the etching chamber such as O ₂ plasma, to remove the inside of the polymer core was produced shell-based spinodal structure.

These structures were manufactured on a micrometer to centimeter scale and using the DLP3D printing method shown in FIGS. 29A-29D. The resulting shell-based polymer structure was then pyrolyzed to 1300 ° C. in vacuum to give a shell-based carbon-reinforced phase that could be impregnated into the matrix.

Example 4: Micrograph of a pyrolysis plate for blast collision test

30A-30C are plates designed for blast collision testing, FIG. 30A is a polymer precursor plate and the resulting pyrolysis plate, FIG. 30B is a photomicrograph of the octet carbon structure, and FIG. 30C is a dodecahedron structure. It is a micrograph of.

Example 5: Enhanced Phase Blocks of Various Architectures

31A-31F are cubes of an example of the strengthening phase, FIG. 31A is a polymer octet cube, FIG. 31B is a pyrolytic carbon octet cube from FIG. 31A, FIG. 31C is a pyrolytic carbon 3D cage beam, and FIGS. 31D-31F are It is an enlarged view of FIGS. 31A to 31C.

Example 6: Direct Manufacture of Tubular Constructed Components

32A to 32F are tubular constructed components, FIG. 32A is a polymer tube having a dodecahedron structure, FIG. 32B is a top view of FIG. 30A, FIG. 32C is an enlarged view of FIG. 31, and FIG. 32D is before impregnation. 32E to 32F of the pyrolytic carbon tube having the dodecahedron structure of FIG. 32A are enlarged views of FIG. 32A.

References corresponding to Examples 1-6

Reference 1: X. Zheng, H. Lee, TH Weisgraber, M. Shusteff, J. DeOtte, EB Duoss, JD Kuntz, MM Biener, Q. Ge, JA Jackson, SO Kucheyev, NX Fang, and CM Spadaccini, “Ultralight , ultrastiff mechanical metamaterials, ”Science, vol. 344, no. 6190, pp. 1373-1377, jun 2014. Available online: http://www.ncbi.nlm.nih.gov/pubmed/24948733 http: // www.sciencemag.org/cgi/doi/10.1 126 / science.1252291

Reference 2: LR Meza, S. Das, and JR Greer, “Strong, lightweight, and recoverable three-dimensional ceramic nanolattices,” Science, vol. 5, no. 6202, pp. 1322-1326, 2014. Available online : www.sciencemag.org/content/345/6202/1322/suppl/DC1

Reference 3: J. Bauer, A. Schroer, R. Schwaiger, and O. Kraft, “Approaching theoretical strength in glassy carbon nanolattices,” Nature Materials, vol. 8, no. February, 2016.

Reference 4: JA Hawreliak, J. Lind, B. Maddox, MI Barham, MC Messner, N. Barton, BJ Jensen, and M. Kumar, “Dynamic Behavior of Engineered Lattice Materials,” Scientific Reports, vol. 6, p. 28094, 2016. Available online: http://www.nature.com/articles/srep28094

Reference 5: CJ Yungwirth, HNG Wadley, JH O'Connor, AJ Zakraysek, and VS Deshpande, “Impact response of sandwich plates with a pyramidal lattice core,” International Journal of Impact Engineering, vol. 35, no. 8, pp. 920-936, 2008.

Reference 6: LR Meza, GP Phlipot, CM Portela, A. Maggi, LC Montemayor, A. Cornelia, DM Kochmann, and JR Greer, “Reexamining the mechanical property space of three-dimensional lattice architectures,” Acta Materialia, vol. 140 , pp. 424-432, 2017. [Online] Available: http://dx.doi.Org/10.1016/j.actamat.2017.08.052

Reference 7: CM Portela, JR Greer, and DM Kochmann, “Impact of node geometry on the effective stiffness of non-slender three-dimensional truss lattice architectures,” Extreme Mechanics Letters, vol. 22, pp. 1 10-138, 2018. Available online: https://doi.Org/10.1016/j.eml.2018.06.004

Reference 8: J.-h. Lee, D. Veysset, JP Singer, M. Retsch, G. Saini, T. Pezeril, KA Nelson, and EL Thomas, “High strain rate deformation of layered nanocomposites,” Nature Communications, vol. . 3, no. May, pp. 1 164-1 169, 2012. Available online: http://dx.doi.org/10.1038/ncomms2166

Reference 9: D. Veysset, AJ Hsieh, S. Kooi, AA Maznev, KA Masser, and KA Nelson, “Dynamics of supersonic microparticle impact on elastomers revealed by real time multi frame imaging,” Nature Publishing Group, pp. 1-6 , 2016. Available online: http://dx.doi.Org/10.1038/srep25577

Reference 10: KH Matlack, A. Bauhofer, S. Kr odel, A. Palermo, and C. Daraio, “Composite 3D-printed meta-structures for low frequency and broadband vibration absorption,” PNAS, pp. 1-5, 2015 Available online: http://arxiv.Org/abs/151 1 09465 {%} 0A http://dx.doi.org/10.1073/pnas.1600171 1 13

Reference 11: S. Krodel and C. Daraio, “Microlattice Metamaterials for Tailoring Ultrasonic Transmission with Elastoacoustic Hybridization,” Physical Review Applied, vol. 6, no. 6, p. 064005, 2016. Available online:

Reference 12: S. Krodel, T. Delpero, A. Bergamini, P. Ermanni, and DM Kochmann, “3D auxetic micro-lattices with independently controllable acoustic band gaps and quasi-static elastic moduli,” Advanced Engineering Materials, vol. 16 , no. 4, pp. 357-363, 2014.

Reference 13: A. Vidyasagar, S. Krodel, and DM Kochmann, “Microstructural patterns with tunable mechanical anisotropy obtained by simulating anisotropic spinodal decomposition,” Proceedings of the Royal Society A: Mathematical, Physical and Engineering Science, vol. 474, no. 2218, p. 20180535, 2018. Available online: http://rspa.royalsocietypublishing.Org/lookup/doi/10. 1098 / rspa.2018.0535

Reference 14: L. Zhang, S. Feih, S. Daynes, S. Chang, MY Wang, J. Wei, and WF Lu, “Energy absorption characteristics of metallic triply periodic minimal surface sheet structures under compressive loading,” Additive Manufacturing, no. August, 2018. Available online: https://linkinghub.elsevier.com/retrieve/pii/S2214860418304688

Reference 15: BD Nguyen, SC Han, YC Jung, and K. Kang, “Design of the P-surfaced shellular, an ultra-low density material with micro-architecture,” Computational Materials Science, vol. 139, pp. 162- 178, 2017. Available online: http://dx.doi.Org/10.1016/j.commatsci.2017.07.025

Reference 16: O. Al-Ketan, R. Rezgui, R. Rowshan, H. Du, NX Fang, and RK Abu Al-Rub, “Microarchitected Stretching-Dominated Mechanical Metamaterials with Minimal Surface Topologies,” Advanced Engineering Materials, vol. 1800029, p. 1800029, 2018.

Example 7: Collision energy calculation and comparison with Kayler

Compare the collision behavior of octet carbon structures manufactured through the two-photon lithography process described in the technical report (ρ = 26% (ρ has a relative density of ρ bar)) with Kevlar 170 g / m ² satin woven fabric. (Reference 1).

The area normalized energy relaxation metric is defined by the following equation.

Where W is the absolute energy (absorbed and / or redirected) and A is the area associated with the effect. Using the value of one sheet of this type of Kevlar sheet (area density ρ _{a, Kev} = 0.17 kg / m ² ), it was obtained as compared to Ψ _Lat = 2.61 × 10 ⁴ J / m ² . The value was Ψ _Kev = 3.26 × 10 ⁴ J / m ² . This difference in metric is mainly due to the difference in scale related to the experiment. Also, the Kevlar sheet had a projectile punctured and lost physical integrity, but the grid had a small permanent deformation that could not be punctured by the collider.

Performing the final one normalization based on the areal density of each material (lattice ρ _{a, Kev} = 0.008 kg / m ² ) is a measure of energy relaxation per kilogram (or "density normalization collision"). The energy relaxation metric ") is provided to be 1.9 × 10 ⁶ J / kg and 3.2 × 10 ⁶ J / kg, respectively, for Kevlar and lattice.

Documents related to Example 7

Reference 1: F. Figucia, US Army R & D Command (1980).

Example 8: Carbon by design through atomic level architecture

Summary: Designing and manufacturing materials with a combination of high strength, high deformability / ductility, large elastic limits, and low densities has been a long-standing challenge, as these properties can be mutually exclusive. there were. Here, by controlling the material of the precursor and the conditions of pyrolysis, pyrolytic carbon micropillars having a specific kind of atomic level architecture were prepared. According to nanomechanical experiments, pyrolytic carbon micropillars have a tensile strength of about 2.5 GPa, a compressive strength approaching the theoretical limit of about 11.0 GPa, a substantial elastic limit of 20 to 30%, and a low density of 1.0. It indicates ~1.89g / cm ^3, which corresponds to a specific strength of 8.07GPa / g · cm ³ greater than the characteristics of all the existing structural materials. Pyrolytic carbon micropillars with a diameter of less than 2.3 μm behaved like rubber and maintained a large compressive strain of about 50% without catastrophic failure, while the large ones suffered brittle fracture with a strain of about 20%. showed that. Large-scale atomic simulations show that these excellent mechanical properties, at least in part, are local deformations of 1 nm curl graphene fragments within pyrolytic carbon microstructures, interactions between adjacent fragments, and between carbon atoms. It was revealed that the existence of strong covalent bonds makes it possible.

In modern advanced material design, the fabrication of high performance materials that combine high strength, substantial deformability, large elastic limits, and low densities has long been a goal and challenge. Almost all structural materials have two sets of obvious contradictions: high strength and high deformability / ductility, high strength and low density. For example, metals and alloys are ductile and can withstand fracture strains in excess of 10% by adjusting dislocation plasticity during deformation (Reference 1). However, their yield strength is usually limited to about 100 MPa, and their elastic limit is only about 2%. Ceramics have higher strength (up to a few GPa), but fracture strain is usually less than 5% due to the absence of movable lattice dislocations during deformation (Reference 1). Metal and ceramic materials generally have densities greater than 2.7 g / cm ³ . Polymers (Reference 2) and porous materials (foam (Reference 3), nanolattice (Reference 4), nanosponge (Reference 5), etc.) are lightweight and have much lower densities than most metals and ceramics. There is. These materials are significantly deformable and can withstand elastic strains of more than 50% (References 2 to 5), but their strength is about 10 MPa.

Numerous studies (References 6-13) have shown that the mechanical properties of materials (strength and ductility, etc.) are highly dependent on their microstructure and the dimensions of intrinsic and extrinsic. Therefore, adjusting the microstructure or the intrinsic and extrinsic dimensions is an effective way to change the mechanical properties of the material. For some polycrystalline metals, the intensity was increased from about 100 MPa to about 1 GPa by reducing the particle size and incorporating nanotwinned structures (References 6 and 7) at the atomic level. High-entropy alloys (HEAs) containing five or more major elements with approximately the same atomic concentration show high yield strength of 1-3 GPa and fracture strain of 10-30% (Reference 8) due to solid solution. It is controlled by mixing a plurality of principal components on a lattice scale (Reference 8). Single crystal metals with an extrinsic dimension (ie, sample size) of less than 10 μm exhibit a so-called “smaller and stronger” size effect (References 9-11). Examples include Au nanowires / nanopillars with a diameter of tens of nanometers showing ultra-high tensile strength of 5.6 GPa, which is close to the theoretical limit (Reference 10). This ultra-high strength is associated with pure, nearly defect-free crystal microstructure and / or depletion of dislocation sources at the nanoscale (Reference 9). For ceramics, a recent study (Reference 12) found that micro-sized shape memory zirconia pillars with few grain along the gauge section withstand about 7% quasi-elastic strain by undergoing martensitic phase transformation. The maximum compressive strength of these ceramic pillars was 1.5 to 2.5 GPa. For polymers, when a strong, hard phase (in the form of nanofibers or nanoparticles) is introduced into the polymer matrix, the resulting polymer-based composite usually has a strength of up to about 0.5 GPa (Reference 13, Reference 13). 14).

Carbon-based materials contain a large number of allotropes due to the unique electronic structure of carbon atoms, which allows the formation of sp, sp ² and sp ³ hybrid bonds. The mechanical and physical properties of carbon materials can vary significantly as a result of various bonding structures. As two typical carbon allotropes, graphene and carbon nanotubes having a 100% sp ² bond have been reported to have ultra-high tensile strengths up to 100 GPa (Reference 16). The mechanical properties of these two allotropes are very sensitive to defects such as vacancies, pentagon-heptagon pairs, and grain boundaries, and the stress concentration around the defects can significantly reduce their strength (References 16-). 20). Although they are not practical for structural applications on larger scales, their three-dimensional (3D) assemblies exhibit hyperelastic behavior by buckling and bending of basic building blocks and scale to macroscopic levels. It can be uploaded (References 21 to 24). The porous microstructure of the 3D graphene assembly allows these construct materials to be extremely lightweight, having a low density of 0.001 to 1.0 g / cm ³ and an excellent elastic limit of up to 50%. However, the strength is about 10 MPa (References 21 to 23). Recently, various pyrolytic carbon materials (References 25-28) have been synthesized by pyrolysis using polymer precursors. The bulk pyrolytic carbon sample prepared at 1000 ° C. (Reference 26) had an optimum hardness of 4 GPa and a density of 1.1 to 1.4 g / cm ³ . Micro-sized glassy carbon synthesized at a high temperature of 400 to 1000 ° C. and a high pressure of 10 to 25 GPa (Reference 27) showed a compressive strength of 9 GPa and a density of 2.0 to 2.5 g / cm ³ . The pyrolytic carbon material usually has a cleavage surface with a fracture strain of less than 3% (Reference 27). A glassy carbon nanolattice with a characteristic strut size of about 200 nm and a density of 0.3-0.7 g / cm ³ (References 28, 29) is a photoresist-based microarchitecture made by photolithography. It was produced by thermal decomposition using, and achieved a compressive strength of about 300 MPa with a fracture strain of less than 10%. The ultrastructure of these pyrolytic carbon materials is usually composed of curved carbon layers or fullerene-like fragments with dimensions of a few nanometers, their mechanical properties and performance being the first precursor, pyrolysis. It is strongly dependent on the later atomic level microstructure and the processing temperature and pressure (References 25 and 26). These studies suggest that multiple properties of a material, including density, strength, and deformability, can be improved simultaneously by designing and controlling atomic-level architectures and reducing characteristic dimensions. It also emphasizes both the promises and challenges associated with the design and manufacture of high performance materials that retain a combination of high strength, significant ductility, large elastic limits, and low densities.

In this exemplary embodiment, we disclose a pyrolytic carbon micropillar with a diameter of 0.7-12.7 μm by two-photon lithography and pyrolysis. By characterization based on transmission electron microscopy (TEM), Raman spectroscopy, and electron energy loss spectroscopy (EELS), these micropillars control conditions for 1 nm sized curl graphene fragments, precursor materials, and thermal decomposition. It became clear that it was composed of an atomic level architecture achieved by doing so. In-situ nanomechanical tests show ultra-high elastic limits of 20-30% pyrolytic carbon, high tensile and high compressive strengths of 2.5 and 11.0 GPa, low densities of 1.0-1.8 g / cm ³ . and shown to have an ultra-high ratio maximum intensity of up to 8.07GPa / g · cm ^3, the sample diameter less than 2.3 .mu.m, it is given a strain greater than 40%, the plastic without substantially destroying It was shown to be deformed and behave like rubber. The experimentally obtained microstructure was incorporated into a large-scale atomic simulation to investigate the deformation mechanism underlying the excellent mechanical properties of pyrolytic carbon pillars under uniaxial compression and tension.

FIG. 33A is a columnar micropillar 6-50 μm in diameter and 12-100 μm high printed using two-photon lithography direct laser lighting (TPLDRW) from IP-Dip, a commercially available acrylic-based photoresist. The schematic diagram of the manufacturing process of is shown. During manufacturing with TPLDLW, the geometry and dimensions of the sample can be precisely controlled. Subsequently, when thermally decomposed in vacuum at 900 ° C. for 5 hours, the polymer sample is completely carbonized and the volume shrinks by 98% (Reference 29). The resulting pyrolytic carbon pillars have diameters in the range of 1.28 to 12.7 μm (20 to 25% of the pre-pyrolytic dimensions) (FIGS. 33B to 33C). The remaining carbon ring visible on the silicon substrate represents the footprints of the original pillars and the constraints imposed by the substrate during pyrolysis. Some samples were manufactured with caps to fit grips for uniaxial tensile testing. Details of the synthesis will be described later in this example.

FIG. 33D contains a representative high resolution TEM (HRTEM) image of a pyrolytic carbon pillar, including an electron diffraction (SAED) pattern of the selected region in the inset, revealing its amorphous microstructure. The magnified TEM image of FIGS. 33E-33F is a large number of 1.0-1.5 nm sizes randomly distributed throughout the pillar volume, producing sub-nanometer-sized voids (indicated by red arrows in FIGS. 33E-33F). Indicates the presence of curl atom fragments. The dimensions of the carbon layer fragments of the pyrolytic carbon sample and the spacing between adjacent layers are much smaller than those previously manufactured (about 4-6 μm and 1.67 to 1.99 nm, respectively) (References 26 and 27). .. These microstructural features provide a useful basis for estimating the density of pyrolytic carbon micropillars by extending the reported geometric model developed for non-graphitized glassy carbon. (Reference 26). In this geometric model, the density depends on the average size and curvature of the carbon layers and the spacing between adjacent layers. Using this model, the density of pyrolytic carbon micropillars in this study was measured to be 1.0-1.8 g / cm ³ . This is close to the density of low density type I glassy carbon (References 27 and 30). Details of the relevant method will be described later in this embodiment. Figure 33G includes a typical pyrolysis shows the Raman spectrum of the carbon micro-pillars, two prominent peaks at Raman shifts 1359Cm ^-1 and 1595cm ^-1 corresponding respectively to D and G peaks of graphite .. The ratio of the integrated area under the integrated area and G band under D band (I D _{/ I G)} represented by the following formula (1), substantially characteristic of a fragment of the carbon layer curled observed in HRTEM images Crystal size L could be calculated (FIGS. 33E to 33F).

Here, alpha is a constant (2.4 × 10 ¹⁰ ), and λ is the wavelength (nanometer unit) of the laser used in the Raman experiment. Using this formula, the characteristic crystal size of the carbon layer fragment was calculated to be 2.4 nm. This is essentially consistent with the size of 1.0-1.5 nm determined from HRTEM observations. It should be noted that in the evaluation of the crystal size of the carbon layer, the HRTEM observation has higher accuracy than the approximate prediction from the equation (1) based on the Raman spectrum. Subsequent calculations calculated the characteristic crystal size of the curled carbon layer to be 1.0-1.5 nm, as derived from HRTEM observations. As shown in FIG. 33H, EELS is revealed the presence of ^{1s-s *} peak at 1s-sigma ^* peaks and 285eV in 292 eV, which are the characteristic s and π bonded to sp ² hybridized carbon It is consistent. The proportion of sp ² bonds was estimated by using the two-window method (Reference 32) and using all sp ² glassy carbon raw materials as reference materials (Reference 27). The proportion of sp ² bonds is as high as 96.5%, indicating the superiority of sp ² hybrids in pyrolytic carbon micropillars. This result shows that the pyrolytic carbon material treated at high temperature mainly contains disordered sp ² bonds because the sp ³ mixed amorphous carbon becomes unstable above about 700 ° C. (Reference 30). This is consistent with the observation (Reference 27). The results also indicate that these bonds correspond to a layer of graphene. Details of estimation and analysis based on Raman spectra and EELS data will be described later in the description of the method of this example. The microstructural properties described above reveal that our pyrolytic carbon is an assembly of nanometer-sized curled graphene fragments interspersed with sub-nanometer-sized voids. Overall, this particular delicate microstructure was designed and created by selecting precursor materials and controlling the dimensions / geometry and pyrolysis conditions of the printed sample.

To characterize the mechanical properties of pyrolytic carbon micropillars, we conducted a series of nanomechanical experiments. Non-in-situ uniaxial compression experiments were performed with nanoindenters with flat punch indenter tips with a diameter of 120 μm. FIG. 34A shows all compressive stress-strain data sets for micropillars with diameters from 4.6 μm to 12.7 μm. It seems that all the micropillars were deformed smoothly until they were broken, first elastically deformed to a strain of about 20 to 30%, and then plastically deformed to a further strain of about 8 to 10% before fracture. A slight misalignment on the top surface of the micropillar caused non-linear behavior under the initial strain of about 1-3%. Based on the fitting of the linear elastic portion of the stress-strain curve in FIG. 34A, Young's modulus was estimated to be 16-26 GPa. The breaking strength of these micropillars increased from 3.8 GPa to 5.6 GPa as the diameter decreased. FIG. 34B shows an SEM image of a typical micropillar with a diameter of 7.17 μm before and after deformation, showing that the μ-pillar was broken into small pieces by brittle fracture.

We also performed similar in-situ compression experiments on micropillars with diameters of about 2 μm or less. The in-situ compression experiment was performed on a custom-made on-site nanomechanical device (SEMenter) capable of accurately controlling deformation by simultaneous video capture (Reference 33). FIG. 34C shows the compressive stress-strain response of a micropillar with a diameter of 2.25 μm, which is a vast plateau-like strain of up to 25%, followed by a linear elastic mode of strain of up to 10%. The plastic region, and when the strain increases by more than about 18%, is characterized by a final stage in which the stress rapidly increases from 5.48 GPa to 12.63 GPa. This stress-strain curve resembles a rubber curve. After unloading from the maximum stress of 12.63 GPa, releasing about 10% elastic strain relieved the micropillars partially. FIG. 34C shows a series of snapshots of this sample during the experiment, where the numbered frames correspond to the red arrows of the same number in the data. It was found that the micropillars gradually shortened and thickened (pyrolytic carbon micros with a diameter of 2.25 μm) until the maximum load strain reached 43.6%, without localization or catastrophic failure. In-situ compression of the pillars takes place). During compression, the micropillars gradually shorten and thicken as the compression strain increases. A slight tilt may occur during compression. After unloading, the micropillars have an elastic recovery of about 10% strain. SEM images obtained from the front and back views of the pillars reveal vertically aligned split microcracks, which are prone to nucleation under large compressive stresses and stress-strain data (FIG. 34C). It led to a slight distortion burst, as indicated by the blue arrow in. In this Example 1, the compressive strength of such a sample corresponds to the stress at the first burst. FIG. 42 shows the detailed in-situ deformation process of another micropillar with a diameter of 2.26 μm under compression, obtaining nucleation and propagation of split microcracks. The d of the corresponding stress-strain data in FIG. 42 shows similar characteristics to the plot in FIG. 34C. The obvious difference between these two datasets is the large strain burst seen in Figure 42d, which can be caused by the rapid propagation of microcracks. Similar signs of deformation and breakage are observed during compression of almost all 2 μm diameter micropillars. To eliminate the potential effect of the remaining carbon ring (FIG. 34B), the ion beam (FIB) milling is focused to remove the ring from the sample (see precompressed FIGS. 34C and 42d). FIG. 43 shows a compressive deformation of a 1.86 μm diameter micropillar holding the remaining carbon ring, which expands from the substrate during compression and separates from the substrate, 36% as shown in FIG. 43d. The degree of distortion resulted in a substantial distortion burst. The maximum stress achieved in FIG. 43d is comparable to that of FIGS. 34C and 42d, suggesting a marginal contribution to the strength of the remaining carbon ring.

Uniaxial tensile experiments on dog bone-shaped specimens made using the same procedure were performed in-situ within the SEMenter, which enables tensile tests that cannot be achieved with ordinary nanoindenters (Reference 33). FIG. 34D summarizes the tensile stress-strain data for samples 0.7-2.0 μm in diameter. After linear elastic loading, it was confirmed that all samples were fractured to 10-25% elongation through brittle fracture (in-situ tensile tests of pyrolytic carbon micropillars 1.5 μm in diameter are performed. The pillars stretch and break, resulting in a smooth fracture surface; tensile fracture strain is up to about 26%). A typical smooth fracture surface is shown in FIG. 34E. The statistical distribution of tensile strength of all pyrolytic carbon samples tested is shown at 34F, which fits the Weibull distribution of two parameters.

Here, σ ₀ and m are material parameters. This distribution gives a characteristic intensity of 1.78 GPa and a low Weibull coefficient of 3.42. This indicates that the fracture strength varies widely. This high variability in the fracture strength of pyrolytic carbon samples suggests that their fracture is likely due to internal defects.

FIG. 35A shows all experimental data obtained from compression experiments, where strength is defined as compressive fracture stress. This plot reveals that for samples larger than 2.3 μm, the compressive strength σ _y increases as the diameter D decreases, according to the power law σ _{y to} D ^−0.37 (FIG. 35A). This scaling law is in good agreement with the theoretical predictions of σ _{y to} D ^−0.40 , and is an asymptotic analysis of a fracture mechanics-based model (Reference 34) that describes the compressive fracture of a characteristic semi-brittle column of diameter D. It is derived from. In this model, the pillars are found to be damaged through the propagation of split cracks of initial length h, similar to experimental observations (eg, FIGS. 34C and 42). This model also provides an equation for the theoretical limit σ _th (Reference 34).

Here, E is Young's modulus and Γ is destructive energy. Using the Young's modulus E = 19.5 GPa (average Young's modulus obtained from compression experiments of all samples) and fracture energy Γ = 29.9 to 61.9 J / m ² of glassy carbon reported in Document 35. , Eq. (2) was used to calculate the theoretical limit range of σ _th = 4.0 to 13.5 GPa of the split crack with an initial length h = 100 nm to 1 μm. This prediction range is similar to the experimentally obtained compression strength of 3.8 to 11.3 GPa (Fig. 35A), and the strength of pyrolytic carbon pillars with a diameter of less than 2.3 μm approaches the theoretical limit. It means that you are. In the highlighted region above the average intensity (shown by the dashed blue line in Figure 35A), the micropillars can withstand ultra-high compressive stresses of 7.2-11.3 GPa and high compressive strains of over 40%. it can. The large variation in the compressive strength of the micropillar with D <2.3 μm mainly results from the variation in the length of the initial split microcrack h. The compressive strength of micropillars with D <2.3 μm is 3.5 times higher than the corresponding tensile strength of 0.8-2.5 GPa on average. This tension-compression asymmetry is consistent with the theoretical predictions of asymmetry coefficients of 2.5-4.4 that occur in high-strength covalent isotropic materials, as determined by recent elliptic fracture criteria ( Document 36). Compression and tensile experiments have revealed that pyrolytic carbon micropillars with D <2.3 μm exhibit high deformability, that is, compressive strain of more than 40% before fracture and tensile strain of about 20%. Repeated compression experiments on these samples showed near complete recovery after each cycle beyond the first cycle. FIG. 35B shows a 20-cycle force-displacement data set of 1.28 μm diameter micropillars with a maximum compressive strain of 23%. These data, combined with the pre-deformation / post-deformation SEM images shown in the inset of FIG. 35B, restored the micropillars by 95% of their original height after 20 cycles of compression to 23% distortion. It is shown that.

To elucidate the underlying mechanism that allows for the observed large deformability and ultra-high strength of small-scale pyrolytic carbon by LAMMPS (Reference 37), it has a diameter of 10-20 nm and is constant. Large-scale molecular dynamics (MD) simulations of uniaxial compression and tension of pyrolytic carbon pillars with an aspect ratio (= 2) were performed. During simulation, nanopillars is compressed or decompressed at a constant strain rate and 300K of constant temperature along the axial direction 5 × 10 8 ^/ sec. Throughout the simulation, adaptive intermolecular reactions empirical bond-order force fields (38) were used to describe interatomic interactions. This force field can capture the formation and breakage of carbon bonds (Reference 38). A complete description of the atomic simulation will be described in the following manner. The simulated sample is composed of many fragments of a curl graphene layer about 1 nm in size and has a density of 1.4 g / cm ³ , which is the TEM of the experimental sample, as shown in FIG. 36A. Consistent with observation. These fragments were linked by covalent bonds or van der Waals interactions. The magnified image in FIG. 36A has an interval of about 0.4 nm between adjacent graphene fragments, and there are several sub-nanometer-sized voids adjacent to them, as in the HRTEM images in FIGS. 33D-33F. Indicates to do. Hybridization of carbon atoms in graphene usually involves sp bonds predominantly concentrated in the edges of the graphene layer, with sp ³ bonds generally interconnecting adjacent graphene layers or forming on a high energy surface (FIG. 36A). .. In certain simulated samples, the proportion of sp ² binding was at least an order of magnitude higher than the proportion of sp and sp ³ (see Figure 44) binding, demonstrating the superiority of sp ² binding, with the above analysis by EELS. Match. Figures 36B-36C show the compressive and tensile stress-strain responses determined from MD simulations, revealing trends and stresses similar to experimental data. This result partially confirms the microstructure and density similarity between the simulated and experimental samples. Figures 36D-36G show some snapshots of the cross sections of the simulated deformation samples at various compression strains. In the early elastic stages, the curl graphene layers approached each other and some were significantly bent (Fig. 36D). As the loaded compressive strain increased, some graphene layers slipped compared to adjacent graphene layers and the graphene layers were abruptly broken under shear (FIGS. 36D-36E). As shown in FIG. 36B, such individual fracture events caused stress fluctuations in the mechanical response with a strain of 21.5%. At 50% compressive strain, sub-nanometer-sized voids collapsed, causing nanopillar densification. Due to the interlayer slip and shear of the adjacent graphene layers, the nanopillars were slightly tilted (Fig. 36F). During unloading, the nanopillars showed a recovery associated with the release of the accumulated elastic strain energy. The distance between graphene layers increased and sub-nanometer-sized voids partially resumed (Fig. 36G). The recovered strain is 19%, which is comparable to the experimental results (Fig. 36B) (in addition, an atomistic simulation of uniaxial compression for a pyrolytic carbon nanopillar with a diameter of 20 nm is performed. In the initial compression stage, curl The graphene layers approach each other and some graphene layers bend significantly. As the compressive strain increases, some graphene layers slide against adjacent layers and the graphene layers are suddenly destroyed under shear. At 50% strain, sub-nanometer-sized voids tend to close and the pillars tend to densify. Sliding and shearing between adjacent graphene layers creates a slight tilt in the nanopillars during unloading. , Nanopillars show constant elastic recovery. The distance between graphene layers has increased and sub-nanometer-sized voids have partially reopened.)

Another similarity of the experiment is that all the simulated nanopillars exposed to tension were destroyed after undergoing a nearly linear elastic deformation (Fig. 36C). 36H-36J show a series of snapshots of cross sections of samples stretched with different strains. It was confirmed that a large number of nanoscale cavities nucleate, expand under tension, and then coalesce to result in the formation of nanoscale cracks (FIGS. 36l and 45A). Ultimately, these nanoscale cracks propagated in the direction perpendicular to the tensile load, resulting in a smooth fracture surface (FIGS. 36J and 45B). This cleavage fracture is similar to the experimental observation shown in FIG. 34E (in addition, an atomic simulation of uniaxial tension on a pyrolytic carbon nanopillar with a diameter of 20 nm is performed. During the tension, the curved graphene layer is stretched. Numerous nanoscale cavities become nuclei, grow and coalesce to form nanoscale cracks. Finally, these nanoscale cracks propagate in the direction perpendicular to the tensile direction and are smooth. It becomes a fracture surface). FIG. 36C shows that the tensile strength of the nanopillar without initial cracks exceeds 20 GPa, which results from the considerable force required to break the strong covalent bond. The strength drops to about 12 GPa after introducing the cracks into the nanopillars, indicating that the presence of initial defects / imperfections significantly reduces the tensile strength of the pyrolytic carbon pillars. FIG. 46 shows the deformation process of nanopillars with nanocracks with initial lengths of 4 nm and 8 nm. It can be seen that their fractures are always due to the growth and expansion of existing nanocracks, leading to smaller fracture strains and smoother fracture surfaces than nanopillars without nanocracks. The tensile strength of the simulated sample is much higher than the tensile strength of the experimental sample, which is about 10-11 orders of magnitude difference in the strain rate loaded in the experiment and simulation, about 1-2 of the sample size. This is a common phenomenon caused by digit differences and unequal defect concentrations. The MD simulation also revealed details of the mechanism for compression and tension of pyrolytic carbon pillars. During compression, sub-nanometer-sized void closure, structural densification, and slip / shear with graphene layer fragments accommodate large deformations. Under stretching, samples with initial defects are destroyed by the coalescence and expansion of existing defects. For samples without initial defects, tensile deformation is governed by nucleation, growth, coalescence, and propagation of the resulting nanoscale cracks (FIGS. 36H-36J and 45A-45B). These basic deformation mechanisms reasonably explain the observed high deformability, high elastic limits, and high strength of pyrolytic carbon micropillars.

In order to examine the properties of the pyrolytic carbon materials examined in this study in their content, FIG. 37A plots the material parameters of the strength vs. density of various structural materials, which are described in the conventional structural materials (Reference 4). , 25, 26, 39) and recently reported high-strength nanomaterials (References 6, 10, 40-43) are also included. In this plot, the strength of pyrolytic carbon in this study is bulk pyrolytic carbon (PyCs) (References 25, 26, 39), graphite, carbon fiber reinforced carbon (C / C) (Reference 40), graphene oxide paper. (GOP) (Reference 41), Copper Nanopillars (Cu-NPs) (Reference 42), Gold Nanopillars (Au-NPs) (Reference 43), Bulk NanoTwin Crystal Copper (NT-Cu) (Reference 6), and Document 28. It is approaching the upper limit of uniaxial strength of the proposed structural material. The strength of the pyrolytic carbon micropillar is comparable to that of carbon microfiber (Reference 44) and gold nanowire (Au-NW) (Reference 10), but its density is about 79 of carbon fiber and Au-NW, respectively. % And 7.3%. FIG. 47 shows shape memory zirconia (Reference 12), SU-8 composite material (Reference 14), carbon microfiber (Reference 44), GOP (Reference 41), Cu-NPs (Reference 42), NT-Cu (Reference 6). ), And similar property plots of strength and fracture strain of various materials, including Zr-based metal glass (MG) (Reference 45). The pyrolytic carbon micropillars in this study show an excellent combination of high strength and high deformability, which is a classic between strength and deformability that has plagued all materials so far. It means overcoming the trade-off. The pyrolytic carbon micropillar has a high tensile strength of 2.5 GPa and a high compressive strength of 11.0 GPa, and has a low density of 1.0 to 1.8 g / cm ³ , so that it has high strength and low strength. The competition with density has been partially overcome, leading to an ultra-high specific strength of 8.07 GPa / g · cm ³ . FIG. 37B shows the inherent specific tensile strength and specific compression specific strength of various materials, where the specific strength of pyrolytic carbon micropillars is an order of magnitude higher than GOP, NT-Cu, and Au-NW, and carbon. It shows that it is comparable to microfiber. Their ratio compressive strength, typical hard ceramic (Document ₄₆₎ (B 4 C, SiC, and the like _Al 2 _{O 3),} metal nanopillars (Cu-NPs (Reference 42) and Au-NPs (Document 43)) , And carbon materials (PyCs (References 25, 26, 39), Graphite, and C / C (Reference 40)), which exceed that of diamond with the highest specific compression strength ever (Reference 28). .. FIG. 37C shows an Ashby plot of specific strength vs. fracture strain of various other materials, including pyrolytic carbon and titanium alloys, magnesium alloys, carbon fiber reinforced polymers (CFRP) and diamonds. In particular, our pyrolytic carbon occupies an unexplored space in the Ashby diagram that no other material can reach. Our experiments and simulations have revealed that pyrolytic carbon micropillars exhibit high deformability, ultra-high elasticity limits, and a unique combination of ultra-high strength and specific strength. These excellent mechanical properties of pyrolytic carbon micropillars result from their microstructure and constituent materials. As a basic building block, a 1 nm sized curl graphene layer has high in-plane stiffness and out-of-plane flexibility as well as high strength. The dense assembly of these graphene layers forms pyrolytic carbon micropillars by covalent bonds or van der Waals interactions. As a result, pyrolytic carbon micropillars maintain large elastic deformation and can withstand large compression and decompression. These results provide a new design route for assembling nanometer-sized curl graphene fragments into high-performance carbon materials.

The pyrolytic carbon micropillars of the present invention will exhibit 1.5 to 8.2 times higher compressive strength and at least an order of magnitude greater fracture strain than existing bulk and microsized pyrolytic carbons (References 26, 27). Please note. These differences in mechanical properties can be attributed to differences in microstructure and sample size between these materials. First, the crystal size of the carbon layer fragments and the spacing between adjacent layers of pyrolytic carbon are those of existing bulk and micro-sized pyrolytic carbons (approximately 4-6 nm and 1.67 to 1.99 nm). Much smaller than (References 26 and 27). These various microstructures are caused by various pyrolysis precursor materials and conditions (such as temperature and duration). Second, the diameter of high-strength, highly deformable pyrolytic carbon is several microns, which is two to four orders of magnitude smaller than the diameter of bulk and micro-sized pyrolytic carbon (over hundreds of microns). 26, 27). Therefore, the design / control of the atomic level microstructure and sample dimensions has significantly improved the mechanical properties of pyrolytic carbon.

In summary, we synthesized new pyrolytic carbon micropillars derived from polymer photoresists via DLW and pyrolysis. These micropillars are composed of curled graphene fragments with an average size of about 1.0-1.5 nm. In both compression and tensile tests, these micropillars were found to exhibit high deformability, ultra-high elasticity limits, and an exceptional combination of ultra-high strength and specific strength. Large-scale MD simulations provided details on the mechanism of deformation of pyrolytic carbon pillars. That is, compressive deformation was dominated by structural densification and slip / shear of the graphene layer, whereas tensile deformation was dominated by expansion of initial defects or nucleation, growth and coalescence of nanoscale cavities. These deformation mechanisms carry a unique combination of desirable properties such as high deformability, high elastic limits and high strength. Our results reveal an important relationship between the microstructure, deformation mechanism, and mechanical properties of pyrolytic carbon materials, thereby providing a potential route for designing and synthesizing new high performance carbon materials. To do.

Method

Sample Preparation: The pyrolytic carbon micropillar fabrication process involves two steps: two-photon lithography and high temperature pyrolysis. First, pillars were synthesized using a 3DTPL DLW (Nanoscribe Photonic Professional) using a dip-in laser lithography configuration, a 63x objective lens, and a commercially available IP-Dip photoresist. For thermal decomposition, the printed polymer sample is heated to 900 ° C. in a vacuum tube furnace at a heating rate of 7.5 ° C./min, then maintained at the target temperature for 5 hours, and finally at room temperature at natural rate. Cooled down to. After pyrolysis, the pillar dimensions are reduced to about 20% to 25% of the original value, which corresponds to 98% volume shrinkage. The diameter D of the pyrolytic carbon pillars for the compression experiment was varied from 1.28 to 12.7 μm. Dog bone-shaped samples with a gauge section of 0.7-2.0 μm were also synthesized using the same procedure as in tensile experiments. The aspect ratio (that is, height to diameter) of the pyrolytic carbon sample was 1.4 to 1.8 for compression and 1.5 to 4.3 for tension.

Ultrastructure characterization: The ultrastructure of the pyrolytic carbon micropillars was characterized by the FEI Technai TF-30 TEM at an acceleration voltage of 300 kV. EELS was performed on FEI Technai TF-20 at an acceleration voltage of 200 kV and the relative ratio of sp ² and sp ³ bonds was estimated. Samples for TEM analysis are prepared using a specific region liftout procedure, the separated lamellas are mounted on a TEM grid, and a voltage of 15 kV and a current of 10 μA in an ion beam (FIB, FEIVersa) is used for 60. The final thickness was reduced to a final thickness of .73 nm. Raman spectra were collected at room temperature using a Raman spectrometer (Renishaw M1000 Micro) equipped with a 514.5 nm laser.

Nanomechanical experiments: Uniaxial compression and all uniaxial tensile experiments on samples with a diameter of 1.18 to 2.28 μm are performed with a custom-made in-situ nanomechanical instrument (SEMENTOR) (Reference 33) using a flat punch indenter tip with a diameter of 10 μm. It was performed at a constant nominal strain rate of 10 ^-3 / sec. Larger samples with diameters of 4.6 to 12.7 μm are Nanoidenter G200XP, Agilent / Keysight Technologies, using a 120 μm diameter flat punch, 0.02 due to load limits in the in-situ device. Compressed at a constant load factor of ~ 0.2 mN / sec. An additional compression experiment was performed on a sample with a diameter of 2.21 to 12.7 μm in G200, and the results of the in-situ experiment were individually verified.

Estimating the Density of Pyrolytic Carbon Micropillars from TEM Analysis: HRTEM images reveal that pyrolytic carbon micropillars consist of randomly distributed curved graphene layers of nanometer size. FIG. 41 provides a comprehensive set of images related to estimating the density of these materials. FIG. 41a shows the distribution of curved graphene segments. FIG. 41b shows individual representative graphene segments, where the average end-to-end length is L and the spacing between adjacent layers is L. In order to estimate the density of pyrolytic carbon (ρ _PC ), it was constructed based on an existing geometric model (Reference 26). The density ρ _CGL of the curved graphene layer can be expressed as follows.

Here, [rho _G is the density of single crystal graphite ^{_{(ρ G = 2.25g / cm 3}} ), L G is the interlayer distance of the single crystal graphite _(L G = 0.34nm), _β is reflected curvature of graphene curved Β = 1 represents a flat graphene layer, and β = π / 2 corresponds to a half circle. FIG. 41c is a schematic diagram showing a rational laminated structure of two curved graphene layers. Using this geometry as a guide, the density of pyrolytic carbon can be estimated by equation (2) below (Reference 26).

Here, θ is the orientation angle between two graphene layers in a typical unit cell (FIG. 41c), and θ = 45 ° corresponds to isotropic pyrolytic carbon (Reference 26) and is curved graphene. The layers are randomly distributed. Based on TEM observations (FIGS. 33E-33F), we obtained β = 1 or π / 2, θ = π / 4, L _s = 0.4 nm, and L = 1.0 to 1.5 nm. By substituting these parameters into equations (1) and (2). ρ _PC = 1.0 to 1.8 g / cm ³ is obtained. FIG. 41d compares experimental data of bulk pyrolytic carbon with our modified model, previous geometric model. The predictions from this modified model are in agreement with the experimental data (References 25, 39).

Estimating Carbon Fragment Size Based on Raman Spectrum: Raman spectroscopy is widely used to study defects and disorders of carbon materials at the nanoscale level, including graphene, carbon nanotubes and glassy carbon. 31, 47). The ratio I _{D /} I _G of the integrated area under the integrated area and G peaks below the D peak in the Raman spectrum is related to the plane crystallite size of the carbon material (L) by Equation (1) (Reference 31). First, four Lorentzian band Raman shift of about ^{_{1580,1350,1620,1200cm -1 (G, D i,}} D 2, D 4) Gaussian band at 1500 cm ^-1 in and literature 47 _{(D 3)} It was used to fit the Raman spectrum of pyrolytic carbon micropillars. Raman spectra shown in FIG. 33G _{is, I D / I G = 6.937} , having a laser wavelength lambda _l = 514.5 nm, which gives L = 2.4 nm by equation (1). This result is consistent with the characteristic size of the curved graphene layer of 1.0-1.5 nm estimated based on HRTEM analysis.

Estimating the proportion of sp ² bonds based on EELS: EELS spectra provide quantitative information about the electronic structure of carbon materials (References 27, 32). A common two-window method (Reference 32) was used to estimate the percentage of sp ² bonds in the pyrolytic carbon micropillars, where the EELS data of the fully sp ² hybridized glassy carbon feedstock was used as a reference. .. Calculate the areas l _π and l _σ under the two windows around the pyrolytic carbon and glassy carbon raw materials π ^* and σ ^* from the pyrolytic carbon EELS data in FIG. 33H and using the glassy carbon raw material. did. Next, the normalized ratio N _int is calculated by the following equation (3) (References 27 and 32).

Here, the superscripts "PC" and "RG" represent pyrolytic carbon and glassy carbon raw materials, respectively. As shown in the following equation (4), the normalized ratio N _int is also a function f of the ratio of sp ² bonds (References 27 and 32).

Assuming that the above formula (3) and the formula (4) are equal to each other, it was found that the ratio of sp ² bonds in the pyrolytic carbon micropillar was 96.5%.

Atomistic simulations: LAMMPS (37) was used to perform a series of large-scale atomistic simulations emulating the uniaxial compression and tension of pyrolytic carbon nanopillars. In order to explain the interatomic interaction, an adaptive intermolecular reaction empirical bond order force field (Reference 38) was used in all simulations. This force field describes bond interactions based on bond order (bond order), unbonded interactions (ie, van der Waals), and twisted interactions. This makes it possible to capture the formation and destruction of carbon bonds (Reference 38). First, samples were made from HRTEM images using experimentally determined microstructures, which contained a large number of curved graphene fragments with an average size of 1 nm. These graphene fragments were extracted from C ₈₄ fullerenes. Numerous such graphene fragments with random orientation were initially hexagonally close-packed into a simulation box sized at 27.5 x 27.5 x 54.3 nm ³ . The system was then equilibrated under an NPT ensemble with energy minimization and 50 ps free relaxation at 300 K. After equilibration, the simulated system density is condensed to 1.40 g / cm ³ (estimated median density of pyrolytic carbon micropillars based on microstructural features) at 550 ps at 300 K via an NVT ensemble, 10 It was hydrostatically compressed at a constant strain rate of ⁹ / sec. After compression, the hydrostatic pressure increased to 10 GPa. Subsequently, a melting-quenching process was performed while keeping the volume constant by limiting all dimensions of the simulation box. During this process, the temperature is first gradually raised from 300K to 1200K within 50ps, then the temperature is maintained at 1200K for 300ps to melt the graphene flakes at high temperature and high pressure, and finally the temperature is lowered from 1200K to 300K at 50ps. It was. Next, under the NPT ensemble, the 200 ps simulation sample at 300 K was relaxed to completely zero the pressure. After relaxation, the size of the simulated sample was 20.5 × 20.4 × 40.8 nm ³ and the density was about 1.4 g / cm ³ . Throughout these processes, periodic boundary conditions were imposed on all three directions of the simulated sample.

Next, nanopillars with diameters of 10 nm and 20 nm were extracted from the relaxed cube sample to perform a uniaxial deformation simulation. To mimic the experiment, the aspect ratio of all nanopillars was maintained near 2. After equilibration, the NVT ensemble compressed or stretched the nanopillars along the axis at a constant strain rate of 5 × 108 / sec and a constant temperature of 300 K. During the simulation, the stress of each atom was calculated based on the virial stress theorem. The compressive and tensile stresses were obtained by averaging the axial stresses of all the atoms in the nanopillar.

We also investigated the effect of defects such as nanoscale cracks on the tensile response of the simulated sample. By removing some atoms from the "as-built" sample, some nanoscale cracks with a length of 4 nm or 8 nm were introduced. After equilibration, the same tensile loads were applied to "as-built" and nanocracked samples and their stress-strain response and fracture were compared. Periodic boundary conditions were imposed along the axial direction of the simulated nanopillars throughout the simulation. By counting the coordination number of each atom, the sp, sp ² , and sp ³ bonds of the simulated sample were identified. It was found that the sp bonds are mainly distributed at the ends of the curved graphene layer, and the sp ³ bonds connect the adjacent graphene layers to each other or are formed on the high energy curved surface of the graphene layer (Fig.). 44). The proportions of sp, sp ² , and sp ³ hybrid bonds in the "as constructed" sample were 8.8%, 89.1%, and 1.8%, respectively, and the simulated sample had sp ² bonds. It was dominant and consistent with the experimental results (Fig. 33H). The remaining 0.3% of bonds were dangling bonds.

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Example 9: Lightweight, defect-free, strong nanostructured carbon

Summary of this exemplary embodiment: A long-standing challenge in modern material design is to produce low density materials that are robust against defects and can withstand extreme thermomechanical environments. This is because these properties are usually mutually exclusive, and the lower the density, the weaker and more fragile the material. We, this until all nano and micro construction of material Nano construct carbon achievable specific strength (strength to density ratio) of 1.90GPa / g · cm ³ which is also superior to 3 orders of magnitude higher than the Developed a simple process to make. We used two-photon lithography and subsequently pyrolyzed in a vacuum of 900 ° C. to create two prototype topologies for pyrolytic carbon: octets and isotras. The size of the unit cell was about 2 μm, the beam diameter was 261 nm to 679 nm, and the density was 0.24 to 1.0 g / cm ³ . Micromechanical experiments have shown that Young's modulus is 0.34 to 18.6 GPa, intensity is 0.05 to 1.9 GPa, and average strain vs. fracture is 14% to 17%. Experiments and simulations have shown that when the density exceeds 0.95 g / cm ³ , these nanolattices become insensitive to manufacturing-induced defects, resulting in near-reaching the theoretical strength of the constituent materials, nano- and micro-constructed carbon. Makes it a particularly promising candidate for structural applications in harsh thermomechanical environments. Describes this combination of high specific strength before fracture, low density, and wide range of deformability in terms of atomic-level microstructure of pyrolytic carbon, nano-sized beam dimensions, and interactions between optimized lattice topologies. To do.

significance:

The strength and density of the porous material typically scale together. A long-standing challenge in modern material design has been to produce lightweight, strong and rigid porous materials. Here, we have shown the fabrication of pyrolytic carbon nanolattices with a designable topology by a two-step procedure of direct laser writing and thermal decomposition at high temperatures. The smallest characteristic size of the stanchions in the nanolattice has approached the resolution limits of available 3D lithograph techniques. We have demonstrated that these pyrolytic carbon nanolattices are 1-3 orders of magnitude stronger than almost all micro / nanostructured materials reported so far.

Lightweight porous materials such as wood, bone, Euplectella sponges, diatoms and bamboo are ubiquitous in nature. These natural structural materials have been extensively investigated (References 1-5) and have been shown to be resilient to externally loaded loads and to be potent in absorbing and dissipating collision energy. Such mechanical elasticity is two main design principles: (i) multi-scale hierarchy in constituent materials and length scales of natural materials, which usually have characteristic dimensions from nano to macro scale. This is made possible by the complex multi-level architectures (5) and (ii) their resistance to defects when the characteristic material length scale is below the critical value (4). Both principles have been applied to the engineering of advanced materials and have been successful to varying degrees (References 5 and 6).

A general guideline for materials that are considered "lightweight" is that their density is lower than that of water (ie, ρ ≤ 1.0 g / cm ³ ) (Reference 16). In particular, recent advances in material processing techniques in three-dimensional (3D) microfabrication and laminate molding often have other beneficial properties such as high specific stiffness, high specific strength, and excellent restoring / recovery. Provided are particularly promising paths for producing materials (References 7-27). The disadvantage of ultra-lightweight in these nano and micro designed materials is the significant reduction in rigidity and strength due to power law scaling: ρ _y ~ (ρ / ρ _s ) ^m , E ~ (ρ / ρ _s ) ⁿ Here, ρ _y is the yield strength, E is Young's modulus, ρ is the density, and ρ _s is the density of a completely dense constituent solid (Reference 1). Since the indices m and n are generally greater than 1, it is possible to develop a methodology for simultaneously producing lightweight and strong materials while maintaining other properties (thermal stability, conductivity, magnetism, recoverability, etc.). .. This is a major challenge due to limited material selection and architecture.

Most research into micro / nano constructive materials to date has focused on hollow beam-based architectures, which are very lightweight with the associated high compliance. The hollow beam gace architecture is, for example, a nickel-based hollow tubular microlattice having an elastic modulus of 529 kPa, a density of about 0.010 g / cm ³ and a compressive strength of about 10 kPa (Reference 7), 0.003 to 1.4 GPa. It is a ceramic hollow tubular nanolattice (References 10 to 14) having a Young's modulus and a density of 0.006 to 0.25 g / cm ³ and a compressive strength of 0.07 to 30 MPa. These micro / nano constructive materials share the common features of the length scale hierarchy, i.e., the relevant dimensions of their structural elements range from 3 to 5 orders of magnitude and tens of nanometers to hundreds of micrometers. Meters and even more. The structural features of nickel-based alloy hollow tubular nanolattices manufactured using large-area projection microsteretic lithography range from tens of nanometers to tens of centimeters in spatial dimensions, ranging from seven digits to over 20% tensile strain. It has a low elastic modulus of 125 kPa and a low tensile strength of about 80 kPa at a density of about 0.20 g / cm ³ corresponding to a relative density of 0.15% (Reference 17). The deformability of these nanolattices is due to the combination of a series of layers and a layered structure of bending and stretching domination over shell buckling, and the elastic instability characteristic of thin-walled hollow cylinders (Reference 17). In a thin-walled architecture, 3D periodic graphene airgel microlattices are synthesized by direct ink writing, these materials are very lightweight, have a density of 0.031 to 0.123 g / cm ³ , and are very. It is adaptable, has a modulus of elasticity of 1-10 MPa, is weak, has a low strength of 0.10 to 1.6 MPa, and exhibits almost complete recovery after compression to 90% strain (Reference 23).

Some efforts have also been devoted to the synthesis and mechanical properties of micro and nano-constructed materials consisting of non-hollow beams of various materials, which are more rigid and more rigid compared to hollow beam counterparts. Provides high density. Most of these studies are based on an architecture consisting of a core-shell beam, usually with an acrylic polymer core and a thin, solid outer coating of tens to hundreds of nanometers. For example, a triangular truss microlattice with a polymer core alumina shell beam is synthesized by combining TPL and ALD and has a low fracture strain of about 4-6% and an elastic modulus of about 30 MPa at a density of 0.42 g / cm ^3. The rate was maintained (Reference 16). An octet truss nanolattice composed of a 262 to 774 nm diameter polymer beam with a sputtered 14-126 nm thick high entropy alloy (HEA) coating has a density of 0.087 to 0.865 g / cm ³ and a Young's modulus. It has been reported that the compression strength is 16 to 95 MPa and the compression strength is 1 to 10 MPa (Reference 20). Samples with a HEA thickness of less than 50 nm were completely recovered after compression of more than 50% (Reference 20). Beyond core-shell beam nano and micro-constructed materials, there are several reports on the production and deformation of 3D structural metamaterials using monolithic beams. For example, a nanocrystalline nickel octet truss nanolattice with a monolithic beam with a diameter of 300-400 nm and a unit cell of 2 μm, made by TPL and subsequent thermal decomposition on a custom synthetic resin, has a density of 2.5 g / cm ³ . It showed an elastic modulus of about 90 MPa, a compressive strength of 189 MPa, and a high fracture strain exceeding 20% (Reference 20). A report on a glassy carbon octet truss microlattice with a beam diameter of about 100 μm, which was produced by thermally decomposing a UV mask patterned polymer template, has an elastic modulus of 1.1 GPa at a density of 0.19 g / cm ³ . A compressive strength of 10.2 MPa and a fracture strain of only about 3% have been reported (Reference 24). A rhombic dodecahedron unit cell and a glassy carbon microlattice with a beam diameter of 50-150 μm, made using stereolithography and pyrolysis, have a density of 0.03-0.05 g / cm ³ and a modulus of elasticity. It was shown that the compression strength was 5 to 25 MPa and the compression strength was 0.08 to 0.35 MPa, and the strain was broken by about 5% (Reference 25). A glassy carbon nanolattice with tetrahedral unit cells made by TPL and pyrolysis has smaller dimensions of 0.97 to 2.02 μm unit cells and a beam diameter of about 200 nm, 0.35 g / cm. ^It had an elastic modulus of 3.2 GPa and a compressive strength of about 280 MPa at a density of about ³ (Reference 18). This brief overview emphasizes the bond between density and adaptability of the designed material, i.e., the lower the density, the softer and weaker the material.

We have developed an approach for producing nanoconstructed pyrolytic carbon, using two-photon lithography and pyrolysis to use the two prototype unit cell geometries shown in Figures 48A-48F, namely octets and isotras. Demonstrated. The octet truss architecture has cubic anisotropy and superior overall properties compared to other conventional lattices (Reference 28) such as triangular, tetrahedral, or cubic trusses and forms, and isotras. The structure is isotropic and has been theorized to retain optimal composition compared to conventional lattice topologies (Reference 29). In the uniaxial compression experiment, the density was in the range of 0.24 to 1.0 g / cm ³ , Young's modulus was 0.34 to 18.6 GPa, fracture strength was 0.05 to 1.9 GPa, and the deformability before breakage was It was revealed that it was 14 to 17%. The highest specific strength is up to 1.90 GPa / g · cm ³ , which is superior to all other mechanically robust lightweight meso / micro / nano lattices reported (References 7-27). This difference is due to the optimized unit cell geometry, reduced feature size, and high quality pyrolytic carbon.

result:

FIG. 48A illustrates the manufacturing process, starting with printing a microgrid of 5x5x5 unit cells from an IP-Dip photoresist using TPL. A 7-10 μm long strut with a circular cross section 0.8-3.0 μm in diameter was printed using high speed galvo mode in a layer-by-layer fashion. The polymer sample was then heated in a vacuum furnace to 900 ° C. at a heating rate of 7.5 ° C./min, pyrolyzed for 5 hours and cooled to room temperature at natural rate (see Method for details). .. Figures 48B and 48D show the CAD design of 10 μm sized octets and isotrus unit cells. The diameter d of the octet truss column is designed to be 0.8 to 2.4 μm. In Isotras, the diameter d ₁ of the vertical struts is 1.4-3.0 μm, and the diameter d ₂ of the predetermined inclined struts is

Maintained as, the _d 2 / _{d 1} ratio based on the optimization of topology was about 1.14 (Reference 29). After thermal decomposition, the polymer changed to carbon form and underwent significant volume shrinkage and mass loss (Reference 30). Each strut shrank to about 20% to 25% of its initial dimensions (FIGS. 48C and 48E), and the associated unit cell size shrank from about 10 μm to about 2 μm. The diameter of the columns, about 261 to 679 nm, obtained after pyrolysis is well below the resolution limits (0,0,0) of most 3D lithography techniques available. The relative density of the pyrolytic carbon nanolattice is estimated to be 17% to 72% by calculating the volume fraction of the solid material in the nanolattice based on the dimensions measured by the 3D CAD model and scanning electron microscope (SEM). It was. The enlarged view in FIG. 48E shows that d ₂ / d ₁ is maintained at about 1.14 after thermal decomposition, which suggests uniform volume shrinkage. FIG. 48F shows a high resolution transmission electron microscope (HRTEM) image of a typical sample extracted from a nanolattice by focused ion beam (FIB) milling, showing a glassy / amorphous microstructure. In our previous studies, using a combination of atomic models and experimental measurements, we estimated the density of pyrolytic carbon produced under these conditions to be around 1.40 g / cm ³ (Reference 31). It is consistent with the density of type I glassy carbon produced at a pyrolysis temperature of less than 2000 ° C. (Reference 32). By multiplying the relative density of the nano-grating this absolute density ^was calculated the density of the nano-lattice varying from ^{0.24g / cm 3 ~1.0g / cm 3} . This is within the lightweight range.

We performed uniaxial compression on all polymer microlattices and pyrolytic carbon nanolattices (see method details). The engineering stresses and strains were determined by normalizing the load-displacement data from the compression experiments by the cross-sectional area and initial height of all samples. Figures 49A and 49B convey the compressive stress-strain response of some typical octet and isotras pyrolytic carbon nanolattices, which appear to be similar in all samples. Since the relative density of the octet truss nanolattice is in the range of 24% to 68%, its Young's modulus increased from 2.57 GPa to 1.73 GPa and its compressive strength increased from 0.21 GPa to 1.73 GPa (Fig.). 49A). The relative density of the isotras nanolattice was rather high, 28-72%, Young's modulus increased from 2.28 GPa to 9.67 GPa, and its compressive strength increased from 0.14 GPa to 1.90 GPa (Fig. 49B). .. The initial non-linearity of stress-strain data results from incomplete initial contact and slight initial inconsistency between coarse lattice planes and flat punches (16). The linear elastic load lasted up to a strain of about 10-20%, after which all pyrolysis samples were catastrophically broken by brittle fracture (FIGS. 49C-49F). The average fracture strain is 14.0% in the case of the octet nanolattice and 16.7% in the case of the isotras nanolattice, which is about 10% of the reported glassy carbon nanolattice (Reference 18), glass. It exceeds about 3 to 5% of the glassy carbon microlattice (References 24 and 25). This enhanced deformability is made possible by the better mechanical stability of the circular struts, which can transfer loads more evenly than the elliptical struts (Reference 33) and have a longer thermal decomposition time. Guarantees and enables sufficient carbonation. Figures 53A-53H show compressive stress-strain data for typical polymer microlattices with octet truss (FIG. 53A) and isotras (FIG. 53E) unit cells for comparison and integrity. This data has an initial nonlinear region exceeding about 2.5% strain caused by slight incomplete initial contact and inconsistency between the coarse lattice surface and the flat punch (Reference 16). The linear elastic load begins with a strain range of about 2.5 to 7.5%, followed by plastic deformation, followed by a stress plateau that extends above 5 to 7.5%. Such stress plateaus correspond to buckling of struts, as evidenced by SEM images (FIGS. 53C and 53G). Table 1 summarizes the Young's modulus and strength of the polymer microlattices tested with different relative densities. At similar relative densities, the Young's modulus of the Isotras microlattice is about twice that of the Octet truss microlattice. The intensity is 1.3 times higher, which is consistent with the prediction (Reference 29).
Table 1. Mechanical properties of polymer microlattices under compression

Figures 50A-50B show all reported micro / nano-constructed materials of carbon, ceramics, or ceramics-polymer composites (References 11, 16, 18, 22-26) vs. pyrolytic carbon nanolattices in this study. The material property space of Young's modulus (FIG. 50A) and compressive strength (FIG. 50B) with respect to the density of These plots show that their elastic moduli and intensities are comparable in density to carbon aerogels (Reference 22), glassy carbon microlattices (Reference 24), and alumina polymer nanostructures (Reference 16). It shows that it is two orders of magnitude larger. The mechanical attributes of the pyrolytic carbon nanolattices in this study range from 0.24 to 1.0 g / cm ³ and between mechanical attributes and densities compared to glassy carbon nanolattices. It shows a scaling index 40% higher than that of (Reference 18). These results, in a density of greater than 0.4 g / cm ³ degrees, the strength and rigidity of the nanostructured carbon in this study, that is greater than the strength and rigidity of all constructs of materials which have been reported to date Suggests. Pyrolytic carbon nanolatts with a density of 1.0 g / cm ³ have a strength of 1.90 GPa and octet trusses with a density of 0.95 g / cm ³ have a strength of 1.73 GPa. theoretical strength _E s / 10 about Jo carbon, that is, comparable to 2～3GPa, where _{E s} is the elastic modulus of the glass-like carbon (0,0,0). FIG. 50A shows the theoretical limit of Young's modulus as a function of density, represented by E = 250ρ (Reference 11), FIG. 50B includes the theoretical limit of strength vs. density, the lower limit of which is defined by diamond. The upper limit corresponds to graphene (Reference 18). Details regarding the determination of these theoretical limits are given in the above method.

FIG. 30. FIGS. 54A and 54B show changes in Young's modulus and compressive strength with respect to relative density, respectively. Due to the relative densities ranging from 17% to 72%, our thermally decomposable carbon nanolattices are E to ρ ^2.25 for octet truss (ρ is ρ bar) and E to ρ ^1.95 for isotras. It has a scaling relationship of Young's ratio, and has a scaling relationship of compressive strength, which is σ _{y to} ρ ^2.41 for octet truss and σ _{y to} ρ ^2.50 (ρ is ρ bar) for isotras. These scaling relationships deviate from the ideal prediction for the stretch-dominated structure (Reference 1), that is, E to ρ (ρ is the ρ bar) and σ _{y to} ρ (ρ is the ρ bar) are mainly due to manufacturing. It is due to structural defects and a non-elongated beam. The SEM images in FIGS. 49C and 49E show typical detectable manufacturing-induced defects found to be present in almost all samples, such as beam junction offsets and bulges, slight curvature of struts, micropits and voids. Some are shown. During compression, these defects cause local deformation and microcracks around the node as well as buckling / bending of the struts, which causes premature structural failure (11). When such local deformations and fractures occur in stretch-dominated lattices, the modulus of elasticity and strength scaling index of the lattice exceeds theoretical expectations and is usually 1.4-2, as exemplified in previous studies. It is in the range of 5 (References 8, 11, 12, 18). Assuming that R is the beam radius and L is the beam length, it has been shown that the slenderness ratio defined as R / L has a great influence on the rigidity and strength of the lattice as well as the geometry of the node (Reference 9, Reference 9). 12, 37). The nodes usually form a solid junction that impedes the rotation of the beam, shortening the effective length of the adjacent beam to some extent, leading to hardening of the entire lattice (Reference 12). Recent calculations and experimental studies have shown that for solid beam octet truss lattices with a beam slenderness ratio greater than 0.06 and a corresponding relative density greater than 10%, the modulus and intensity scaling relationships deviate from existing analytical theories. It turns out that it has an index of 2.20 and 1.88 instead of 1.0 (Reference 12). The beam slenderness ratio R / L of the octet truss nanolattice in this study was 0.07 to 0.24, and 0.07 to 0.12 of the monoric polymer octet truss nanolattice (Reference 12), tetrahedral unit cells. It is the same as 0.06 to 0.20 of the glassy carbon (Reference 18) having. The Young's modulus of nanoconstructed carbon with a relative density of 15% to 80% in this study was 2.25 (octet truss) and 1.90 (isotras) scaling index, and the intensity was 2.41 (octet truss). And the scaling index of 2.50 (isotras) is consistent with existing reports (Reference 12). Figures 50A-50B report that these relatively high scaling indices for the mechanical attributes of pyrolytic carbon nanolattices provide the highest stiffness and strength reported so far (References 11 and 18).

To further investigate the effect of initial defects on the mechanical properties of the pyrolytic carbon nanolattice, a series of finite element (FE) simulations were performed to compress samples with relative densities varying from 15.9% to 70%. .. Details of the FE simulation are described in the above method. The simulated nanolattice has three unit cell geometries, namely octet trusses and isotras for comparison with experiments, and tetrahedral trusses for comparison with previous literature (Reference 18). It was found that the initial deflection of the struts could reduce the compressive strength of the nanolattice at lower relative densities. Figures 51A-51C show nanolattices simulated using different unit cells, with existing defects defining maximum strut deflection as 5%, 10%, 15% of edge length. It was formed by imposing a corresponding buckling-specific mode (Reference 18). After introducing these initial deflections, some struts remain pre-bent prior to compression, similar to structural defects in experimental samples (FIGS. 49C and 49E). Also, for reference, a complete nanolattice compression was simulated. Figures 51D-51F show the compressive stress-strain response of up to 12% strain of simulated nanolattices, indicating that the strength of nanolattices with initial deflection is always lower than the strength of their full nanolattices. It is clear. FIG. 55 shows that the FE simulation reveals a tendency similar to the experimental measurement of the dependence of modulus of elasticity and strength on relative density. Figures 56A-56B quantify the variation in strength reduction as a function of initial deflection with respect to the strength of the complete nanolattice, (i) for a given relative density and architecture, the relative reduction in strength is the initial deflection. (Ii) For a given architecture, dense nanolattices have less relative weakening due to defects, and (iii) nanolattices with tetrahedral truss unit cells have defects. It has been shown to be most sensitive to octet trusses and isotras for all densities. For example, for a relative density of 15.9%, the relative reduction in strength is 15% for maximum deflection, 2% for isotras architectures, and 15% for octet truss architectures. A similar relative weakening for a relative density of 70% is less than 1%.

Results from our current experimental and computational studies show that carbon nanolattices with inherently fragile isotras and octet truss architectures exhibit reduced sensitivity to defects at high densities. This can be explained by the local failure of the individual struts that redistribute the elastic energy stored between the other load-bearing truss members instead of causing catastrophic structural failure. This is consistent with the achievement of near-theoretical strength of carbon nanolattices with densities above 0.95 g / cm ³ . When the strut diameter is reduced by hundreds of nanometers to a size comparable to the defect-insensitive critical size of the component, the strut exhibits high strength and good defect resistance, which to some extent the high strength of the carbon nanolattice. This is due to the local stress and the volume fraction of the struts (Reference 4). The low density nanolattice has thinner and elongated struts, and due to its small cross-sectional area, the local stress is higher during compression and the contribution of the nodes is negligible (References 12 and 37). In this case, the higher local stress causes some struts to buckle faster and the stress concentration around the node to increase. Due to the low volume fraction of thinner struts, lower density nano struts (ie, thinner struts) can be destroyed with less overall stress. In contrast, higher density (ie, thicker strut) nanolattices have a greater contribution to the load bearing capacity of the node due to the larger cross-sectional area of each strut and thus have lower local stresses. , The load applied to the entire nanolattice is distributed relatively uniformly (References 12 and 37). Under these conditions, the nanolattice breaks when the local stress of the strut approaches the theoretical strength of the constituent carbons. Such local stresses and higher volume fractions of the struts ultimately result in higher densities and higher strength nanolattices. Optimized unit cell geometries such as octet trusses and isot trusses make it easier to achieve high strength as well as better scratch resistance.

Figure 52 is a specific strength of pyrolytic carbon nano grating is from 0.146GPa / g · cm ³ in the range of 1.90GPa / g · cm ^3, which is hollow tubular nickel (Document 7) and NiP (literature 8 ), Copper (Reference 19), TiAl ₆ V ₄ (Reference 27) microlattice, and hollow beam alumina (Reference 11), alumina polymer (Reference 16), and metallic glass Zr ₅₄ Ni ₂₈ Al ₁₈ nanolattice (Reference 33). ) Has been shown to improve by a few orders of magnitude over all previously reported nano- and micro-constructed periodic lattices. The maximum specific strength of the carbon nanolattice in this study was 1.0 g / cm ³ and was 2.4 times the 0.80 GPa / g · cm ³ reported for the glassy carbon nanolattice (Reference 18). ), 35% of 5.60 GPa / g · cm ³ of dense diamond having the highest specific strength (Reference 18) of all bulk materials. Such ultra-high specific strengths of our pyrolytic carbon nanolattices result from both nanosized beam diameters and optimized lattice topologies.

In this exemplary example, we create micro and nano-constructed pyrolytic carbon with densities less than 1.0 g / cm ³ , GPa level strength, and greater than 10% deformability before fracture. Developed a laminated molding method. As a starting point from all existing studies on micro / nano lattices (References 11, 16, 18, 22-26), the elastic modulus and strength of nanostructured carbons in this study approaches theoretical limits. Due to the rational design of the lattice topology with suitable microstructure and characteristic material dimensions of nano and microscale, Young's modulus of 0.34 to 18.6 GPa at a density of 0.24 to 1.0 g / cm ³ . Prototype architectures of octets and isotras pyrolytic carbon nanolattices with densities of 0.05 to 1.90 GPa can be created, with specific strengths of 0.146 to 1. which were not achieved with carbon-based or building materials. It is converted to 90 GPa / g · cm ³ . This nanoconstructed carbon also exhibits an average fracture strain of 14.0 to 16.7%, exceeding the average fracture strain of all other reported brittle constructive materials. Experiments and simulations have shown that for densities above 0.95 g / cm ³ , these samples are substantially insensitive to manufacturing-induced defects, thereby achieving a theoretical strength of 1.90 GPa. Make it a particularly advantageous candidate for structural applications. This study provides insights into the basic scientific principles governing the design and properties of nanoconstructed materials, for their use in scalable manufacturing due to their resistance to defects, ultra-lightweight, and excellent strength. Provides a viable route for.

Materials and methods:

Production of Pyrolytic Carbon Nano Lattice: First, a polymer microlattice was produced from an IP-Dip photoresist using a TPL DLW (manufactured by Nanoscribe) at a speed of 10,000 μm / sec and a laser output of 17.5 mW. During the DLW process, struts with a circular cross section with a diameter of 0.8-3.0 μm were printed in a layer-by-layer fashion in high speed galvo mode. All printed polymeric microlattices have two typical unit cell geometries, one octet truss (Fig. 48B) and the other isotrus (Fig. 48D). The unit cell size of the polymer microlattice is about 10 μm. Next, the polymer microlattice was heated at a heating rate of 7.5 ° C./min, thermally decomposed in vacuum at 900 ° C. for 5 hours, and then cooled to room temperature at a natural rate. After thermal decomposition, the polymer microlattice was transformed into a pyrolytic carbon nanolattice by chlorination induced by mass loss of the polymer at high temperature (Reference 30). The diameter of all struts in the pyrolytic carbon nanolattice shrinks isotropically to about 261-679 nm, which is about 20% to 25% of their initial dimensions (FIGS. 48C and 48E). The unit cell size of all pyrolytic carbon nanolattices is about 2 μm.

Mechanical test: Uniaxial compression experiments were performed on all the polymer microlattices and pyrolytic carbon nanolattices produced. Some of these experiments are in-situ devices (InSEM, nanomechanics) using flat diamond punches with a diameter of 170 μm at a constant strain rate of 10 ³ / sec to reveal deformation morphology at the same time as acquiring mechanical data. Made by the company). Another experiment was performed on a nanoindenter (G200, Agilent / Keysight Technologies) with a constant loading factor of 0.2 mN / sec.

Finite element modeling: A series of FE modeling was performed by Abaqus to compress the pyrolytic carbon nanolattice. An isotropic linear elastic material was used for modeling. All nanolattices were modeled with beam elements. The Young's modulus of the material was 20 GPa (Reference 34), and the Poisson's ratio was 0.15 (Reference 18). The simulated nanolattice has three types of unit cell geometries, such as octet truss, isotras, and tetrahedron truss. For each type of nanolattice, the unit cell size is set to 2 μm and the relative density varies from 15.9% to 70% by varying the diameter of the struts. Initial deflection is introduced into the simulated nanolattice struts by imposing the corresponding nanolattice buckling-specific mode prior to compression (eg, FIGS. 51A-51C). The maximum deflection of the strut is set to 5%, 10%, 15% of the edge length. After introducing the initial deflection, some struts remain pre-bent before compression, which is very similar to some structural defects in the experimental sample (FIGS. 49C and 49E). During compression, the lower part of the nanolattice is fixed and the upper part is loaded with a displacement load. As a reference for addressing the effects of defects / imperfections on the mechanical properties and the response of the nanolattice, we simulated the compression of the nanolattice using perfectly straight struts.

Determination of Young's Modulus and Strength Theoretical Limits for Density: Modulus-Density Theoretical Limits are obtained from Ref. 11 and are based on Granta Design, the standard software for material selection and graphic analysis of material properties. It is determined by the boundaries of many data of the actual material. For more information on Granta design, please refer to the web page (https://www.arantadesian.com/) and related software documentation. The strength-density limit is defined in Ref. 18 for a particular range based on measurements of all materials to date. The lower limit of this range is determined by diamond, which has the highest specific strength of all bulk materials, and the upper limit is determined by graphene, which retains the highest strength of all materials to date.

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Example 10: A scalable manufacturing method of a 3D constructed structure using laminated molding and its thermal decomposition.

Figures 57A-57D show the features of the microstructure of 3D constructed carbon. The cross-sectional image of the 3D constructed carbon shown in FIG. 57A shows a monolithic structure without micropores. EDS analysis of the cross section shows that the 3D constructed carbon is composed of an average of 98.4% carbon and a small amount of oxygen. Line analysis shows a uniform elemental composition over the entire cross section (FIG. 58A). FIG. 57C shows an XRD pattern with three wide peaks at 23.5 °, 44.3 ° and 79.8 ° at 2θ corresponding to (002), (100) / (101) and (110) of graphite. Shown. The average interlayer spacing and crystallite size of graphene sheets along (002) (ie, d ₀₀₂ and L _c ) are estimated to be 3.78 A and 9.3 A using Bragg's law and Scheller's equation, respectively. It suggests that there are several laminated graphite layers on average. Raman spectra shown in FIG. 57B, five peaks, i.e. D1 weak peaks strong peaks (at 1355 cm ^-1) and G (at 1603cm ^-1) and D3 weak peak (at 1613cm ^-1), D3 It was degraded (at 1539cm ^-1) and the D4 (at 1225 cm ^-1). The G peak corresponds to the in-plane bond stretching motion of a pair of sp ² carbon atoms having E _2g symmetry (Reference 24). The D1 peak is displayed only in the presence of graphite disorder, and the _A1g symmetry is obtained. Corresponds to the graphite lattice vibration mode having (Reference 24). The D2 peak is thought to be due to the lattice vibration of graphite, and the D3 and D4 peaks were found in amorphous or glassy carbon in other studies (References 25 and 26). The high resolution image of the TEM in FIG. 57D confirmed an entangled microstructure containing several stacked graphite layers. The diffusion diffraction rings at (002), (100) / (101) and (110) in the inset show a chaotic carbon microstructure that coincides with the broad peaks of the XRD pattern.

The mechanical behavior of 3D constructed carbon was evaluated by using a uniaxial compression test with a microcamera. Figures 59A-59B show typical mechanical responses using in-situ compression side views of some of the architecture near each collapse event (supplementary access to complete video). Note that the bottom side of the 3D architecture was missing when the 3D architecture was removed from the material of the DLP 3D printer.

Following the initial stress release, the load with local fracture events gradually decreased, as shown by the red circles in Figures 59B II-a, II-b, II-c. The second stress release event occurred when the contact area of the 3D constructed carbon on the substrate was fractured by the half layer (Fig. 59BIV). At the third stress release, the 3D constructed carbons almost completely contacted on the substrate and collapsed in half a layer, showing the maximum yield stress (29.9 MPa). These two collapse events with low yield intensity followed by a third collapse with high intensity were repeatedly observed (Fig. 60). Table 2 summarizes the average yield strength at each stress release event.
Table 2: Mean and standard deviation of first, second and third yield stresses

Example 11: Geometry without nodes

61A and 61B. An image showing constructed three-dimensional structure with nodeless geometry according to a particular embodiment of the invention. Geometry without additional exemplary nodes can be found in Abueidda et al. ("Effective Conductivity and Modulus of New Foam with Triple Periodic Minimal Surfaces", Mechanics of Materials, vol. 95, April 2016, pp. 102-115), which is incorporated herein by reference.

Example 12: Impregnation of carbon-reinforced phase

The constructed carbon-reinforced phase was impregnated with a low-viscosity epoxy and cured at room temperature for 8 hours.

Spinodal cube shown in FIG. 62A and 62D are, ρ = 123.7kg / ^{m 3} has been experienced density transition to 1152kg / ^{m 3} from octet cube of FIG. 62C and 62F are, 1,160 kg / ^{m 3} from 255 kg / ^{m 3} Increased to. The resulting composite is less than half the average density of light metals (ie, ρ _Al = 2700 kg / m ³ ) and is more than some carbon fiber reinforced polymers with densities typically around 1600 to 1800 kg / m ^3. Also had an average density that was at least 40% lower.

Example 13: Mechanical test

Octet phase and uniaxial compression of composites

We performed quasi-static uniaxial compression on the octet carbon-reinforced phase with and without epoxy impregnation. The experiment was performed at a strain rate of 10 ^-3 / sec. The underlying tessellation consisted of 17 x 17 x 17 unit cells (having a characteristic unit cell size of about 590 μm), so scale sufficient to discuss effective material properties (rather than structure). Is assumed to be separated.

The experiments shown in FIGS. 63A-63D were used to obtain the effective modulus and yield strength of this particular carbon material with a relative density of about 15%. Young's modulus was calculated to be 669.7 MPa and 495 MPa, and yield strength (defined as stress in the initial catastrophic fracture event) was calculated to be 11.33 MPa and 8.67 MPa. These carbon phases were susceptible to sample defects due to the lack of a matrix phase.

Epoxy impregnation significantly changed the mechanical behavior of these materials. Most notably, the material did not experience catastrophic events, and significant strain hardening occurred above ε> 0.1. Figures 64A-64C show the response of the material under compression with significant densification after ε = 0.5, with no fatal breakage or cracking of the entire sample.

These experiments resulted in Young's modulus of 1.82 and 2.24 GPa and yield strength of 59.8 and 69.6 MPa. At these strength values, the material is as strong as some metals (eg, copper), but 7.7 times less dense. The yield strength σ _y of these materials is defined using a 0.2% strain offset from the linear regime, and also the deformation stress σ _f corresponding to the maximum stress before the negative tangential coefficient is observed. The deformation stresses of these samples were calculated to be 71.9 and 78.2 MPa. Table 3 summarizes the comparison between these values and some metals.
Table 3: Mechanical parameters of tested carbon octet materials, including comparisons with several metals.

The non-destructive and strain hardening response of these composites makes them very suitable for energy absorption applications. For these quasi-static experiments, specific energy absorption (SEA) can be defined as follows.

Here, W is the strain energy density, which is defined by W = ∫σdε, and ρ is the material density. When this formula was calculated for the experiment in FIG. 64C, it was W = 30.9 MJ / m ³ and Ψ = 26.7 J / g at a density of ρ = 1159 kg / m ³ . Comparing these values with the stainless steel 316L octet (Reference 1) whose reported values are Ψ = 10.1J / g and ρ = 2160kg / m ³ , the carbon octet composite material has twice the SEA at half the density. It can be seen that it has a capacity.

4-point bending of octet phase and composite

We also investigated the bending behavior of octet carbon materials using a 4-point bending setup according to ASTM standard D6272.

Bending of the carbon phase without matrix showed catastrophic failure, as observed in compression experiments. FIG 65A~65D shows the data corresponding to the experiment, from which approximate the bending strength 10.34MPa, the flexural modulus becomes _E B = 1.4 GPa.

The same experiment was performed on the epoxy impregnated material and flexural moduli of 3.3 and 3.9 GPa were obtained. No bending strength values were calculated because no breakage was observed within the tolerance limits of this ASTM standard. After undergoing major bending, the sample returned to its original shape with no apparent permanent deformation or cracking.

References 1: T. Tancogne-Dejean, AB Spierings, and D. Mohr, “Additively-manufactured metallic micro-lattice materials for high specific energy absorption under static and dynamic loading,” Acta Materialia, Corresponding to Examples 11-12. vol. 116, pp. 14-28, 2016. Available online: http://linkinghub. elsevier.com/retrieve/pii/S1359645416304153

Statements and Modifications Concerning Incorporation by References All references throughout this application, such as patent documents containing patents or equivalents issued or granted; patent application publications; and non-patent literature documents or other source material. The whole is incorporated as an individual reference to the extent consistent with the disclosure of the present application (for example, a partially contradictory document is incorporated as a reference except for a partially contradictory part of the document).

The terms and expressions used herein are used as descriptive terms, but not in limitation, with the intent of excluding the features shown and described or their equivalents in the use of such terms and expressions. There is no. However, it is recognized that various modifications are possible within the scope of the claimed invention. Accordingly, although the present invention has been specifically disclosed by preferred embodiments, exemplary embodiments and any features, modifications and modifications of the concepts disclosed herein can be made by one of ordinary skill in the art, and the like. It should be understood that such modifications and modifications are considered to be within the scope of the invention as defined by the appended claims. The particular embodiments provided herein are examples of useful embodiments of the present invention that can be performed using a number of modifications of the devices, device components, and method steps described. It will be obvious to those skilled in the art. As will be apparent to those skilled in the art in the present description, methods and devices useful in this method may include a number of arbitrary compositions and processing elements and steps.

As used herein and in the appended claims, the singular forms "a", "an", and "the" include multiple references unless the context clearly dictates otherwise. Thus, for example, reference to a "cell" includes a plurality of such cells and their equivalents known to those of skill in the art. Similarly, the terms "a" (or "an"), "one or more" and "at least one" can be used interchangeably herein. It should also be noted that the terms "prepare," "include," and "have" can be used interchangeably. The expression "any of claims XX to YY" (where XX and YY refer to claim numbers) is intended to provide a plurality of dependent claims in an alternative form, and in some embodiments, "claims". It is replaced with "any one of XX to YY".

It is understood that when a group of substituents is disclosed herein, all individual members of that group and all subgroups are disclosed separately. When Markush groups or other groupings are used herein, all individual members of the group and all possible combinations and sub-combinations of the group are intended to be individually included in the disclosure. .. If a particular isomer, enantiomer or diastereomer of a compound is described herein individually or in any combination, for example as not specified by a formula or chemical name, the description thereof will be described. , Each isomer and enantiomer of the described compounds are intended to be included. Furthermore, unless otherwise specified, all isotopic variants of the compounds disclosed herein are intended to be included in this disclosure. For example, it is understood that any one or more hydrogens in the disclosed molecule can be replaced with deuterium or tritium. Isotopic variants of the molecule are generally useful as standards for analytical evaluation of the molecule and for chemical and biological studies related to the molecule or its use. Methods for making such isotope variants are known in the art. The specific names of the compounds are intended to be exemplary, as it is known that those skilled in the art can give different names to the same compounds.

All systems, structures, shapes, features, combinations or methods thereof described or exemplified herein can be used to carry out the present invention unless otherwise specified.

Whenever a range, eg, a temperature range, a time range, or a composition or concentration range, is given herein, all intermediate and subranges and all individual values contained within the given range are disclosed. include. It will be appreciated that any subrange or range or individual values within a subrange contained in the description herein can be excluded from the scope of the claims herein.

All patents and publications referred to herein indicate the level of skill of one of ordinary skill in the art to which the invention relates. The references cited herein are incorporated herein by reference in their entirety, indicating their prior art on the date of publication or filing, and where appropriate, this information is used herein. Certain embodiments in the prior art can be excluded. For example, if the composition of the material is claimed, it is known and available in the art prior to the applicant's invention, including compounds for which feasible disclosure is provided in the literature cited herein. Compounds are not intended to be included in the claims of the compositions herein.

As used herein, "comprising" is synonymous with "including," "containing," or "characterized by," and is inclusive or open. It is an end and does not exclude additional unlisted elements or method steps. As used herein, "consisting of" excludes any element, step, or component not specified in the claim element. As used herein, "consisting essentially of" does not exclude materials or steps that do not substantially affect the basic and novel features of the claim. In each of the examples herein, the terms "contains," "consisting essentially of," and "consisting of" can all be replaced with any of the other two terms. The invention exemplified herein can be practiced without any one or more elements, limitations or limitations not specifically disclosed herein.

Those skilled in the art can use starting materials, biological materials, reagents, synthetic methods, purification methods, analytical methods, analytical evaluation methods, and biological methods other than those specifically exemplified in the practice of the present invention. You will understand. You don't have to rely on undue experimentation. All functional equivalents known in the art of such materials and methods are intended to be included in the present invention. The terms and expressions used are used as descriptive terms rather than limitations, and the use of such terms and expressions is not intended to exclude the features shown and described or their equivalents, but is claimed. We recognize that various modifications can be made within the scope of the present invention. Accordingly, although the present invention has been specifically disclosed by the features of preferred embodiments and options, modifications and variations of the concepts disclosed herein can be made by one of ordinary skill in the art, and such modifications and modifications. Modifications are within the scope of the invention as defined by the appended claims.

Claims (99)

It ’s a composite material system.
Structures with constructed 3D geometry that are monolithic and deterministic,
With a matrix phase that at least partially impregnates the structure,
A composite material system characterized by comprising. The composite material system of claim 1, wherein the three-dimensional geometry is a nano or micro-constructed three-dimensional geometry. The composite according to claim 1 or 2, wherein the structure is characterized by an area normalized energy relaxation metric (Ψ) selected from the range 2 × 10 ⁴ J / m ² to 4 × 10 ⁵ J / m ^2. Material system. The composite according to claim 1 or 2, wherein the structure is characterized by a density normalized energy relaxation metric (Ψ) selected from the range of 1.9 × 10 ⁶ J / kg to 4 × 10 ⁶ J / kg. Material system. The composite material system according to any one of claims 1 to 4, wherein the structure is characterized by a bounce coefficient selected from the range of 0.8 to 0.7. The composite material system according to any one of claims 1 to 4, wherein the structure is characterized by a bounce coefficient selected from the range of 0.8 to 0.3. The composite material system according to any one of claims 1 to 6, wherein the composite material system is configured such that collision energy is substantially absorbed by the structure. The composite material system according to any one of claims 1 to 7, wherein the structure is characterized by at least one vibration frequency bandgap. The composite material system of claim 8, wherein the at least one vibration frequency bandgap is deterministic. The composite material system according to claim 8 or 9, wherein the at least one vibration frequency bandgap is in the range of 0.1 MHz to 200 MHz. The composite material system according to any one of claims 1 to 10, characterized by an attenuation ratio of at least 1.2%. The composite material system according to any one of claims 1 to 11, wherein the three-dimensional geometry has at least one surface feature. The composite material system according to any one of claims 1 to 12, wherein the three-dimensional geometry has at least one surface feature and at least a portion of the at least one surface feature is characterized by a non-zero Gaussian curvature. The composite material system according to any one of claims 1 to 13, wherein the three-dimensional geometry has at least one surface feature and at least a portion of the at least one surface feature is characterized by a non-zero mean curvature. The composite material system according to any one of claims 1 to 14, wherein the three-dimensional geometry has at least one surface feature and at least a portion of the at least one surface feature is characterized by zero mean curvature. 10. One of claims 1-15, wherein the three-dimensional geometry has at least one surface feature and at least a portion of the at least one surface feature is characterized by a non-uniform Gaussian curvature or a non-uniform mean curvature. The composite material system described. The third aspect of any one of claims 1-16, wherein the three-dimensional geometry has at least one surface feature and at least a portion of the at least one surface feature is characterized by a uniform Gaussian curvature or a uniform mean curvature. Composite material system. The three-dimensional geometry has at least one surface feature and
The composite material system according to any one of claims 1 to 17, wherein the thickness dimension of the at least one surface feature is not uniform over the entire at least one surface feature. The three-dimensional geometry has at least one surface feature and
The composite material system according to any one of claims 1 to 18, wherein the thickness dimension of the at least one surface feature is uniform over the entire at least one surface feature. The composite material system according to any one of claims 1 to 19, wherein the three-dimensional geometry is characterized as a spinodal geometry. The composite material system according to any one of claims 1 to 20, wherein the structure is characterized by a gradient of normalized modulus vs. relative density, selected from the range 1-1.3. The composite material system according to any one of claims 1 to 21, wherein the three-dimensional geometry includes a resonator. The composite material system according to any one of claims 1 to 22, wherein the three-dimensional geometry is characterized by a unit cell geometry having a resonator. The composite material system of claim 22 or 23, wherein the resonator has micro-inertial features connected to at least one other feature of the three-dimensional geometry. The composite material system according to any one of claims 22 to 24, wherein the resonator has a micro-inertia feature, and the other feature of the three-dimensional geometry has the micro-inertia feature. The composite material system according to any one of claims 22 to 25, wherein the resonator comprises a cantilever beam feature and a micro-inertial feature connected to the end of the cantilever beam feature. Any of claims 1-26, wherein the structure is characterized by deterministic anisotropic damping characterized by damping at least 1% more along the first direction than along the second direction. The composite material system according to one item. The composite material system according to any one of claims 1 to 27, wherein the structure exhibits oscillating black scattering and the structure does not exhibit oscillating local resonance. The composite material system according to any one of claims 1 to 28, wherein the structure comprises a carbon allotrope material, a polymer ceramic material, a metal material or a combination thereof. The composite material system according to any one of claims 1 to 29, wherein the structure comprises at least one of a carbon allotrope material, a polymer, a ceramic material, or a combination thereof. The composite material system according to any one of claims 1 to 30, wherein the structure comprises one or more carbon allotrope materials. The composite material system according to any one of claims 29 to 31, wherein the structure contains at least 50% by volume of one or more carbon materials. 32. The composite material system of claim 32, wherein the structure has a plurality of features characterized by a core of at least 50% by volume of one or more carbon materials. The composite material system according to any one of claims 1 to 33, wherein the three-dimensional geometry is a nodeless geometry. The composite material system according to any one of claims 1 to 34, wherein the structure has at least one hollow feature. The composite material system according to any one of claims 1 to 35, wherein the structure is formed by laminated molding. Claims 1-36, wherein the three-dimensional geometry has at least one longitudinal feature and at least a portion of the at least one longitudinal feature is characterized by a non-zero curvature along the longitudinal direction of the feature. The composite material system according to any one item. The three-dimensional geometry has at least one longitudinal feature, and at least a portion of the at least one longitudinal feature is characterized by a non-uniform curvature along the longitudinal direction of the feature. The composite material system according to any one of 37. One of claims 1-38, wherein the three-dimensional geometry has at least one longitudinal feature having at least one cross-sectional dimension, the at least one cross-sectional dimension being non-uniform along the longitudinal direction of the feature. The composite material system according to one item. The composite material system according to any one of claims 1 to 39, wherein the three-dimensional geometry has at least one feature having a non-uniform cross-sectional shape. The composite material system according to any one of claims 1 to 40, wherein the three-dimensional geometry has at least one longitudinal feature having a longitudinal axis oriented perpendicular to the thickness direction of the structure. The structure defines a three-dimensional outer boundary shape, and the three-dimensional geometry has at least one feature that intersects the boundary at only one point or at zero point, according to any one of claims 1-41. The composite material system described. The composite material system according to any one of claims 1 to 42, wherein the three-dimensional outer boundary shape of the structure corresponds to the shape of the composite material system. The composite material system according to any one of claims 1 to 43, wherein the three-dimensional outer boundary shape defined by the structure is hollow. The three-dimensional geometry is an overall three-dimensional geometry including at least a primary three-dimensional geometry and a secondary three-dimensional geometry, and any of claims 1 to 44, wherein the primary three-dimensional geometry and the secondary three-dimensional geometry are different. The composite material system described in item 1. The composite material system according to any one of claims 1 to 45, wherein the matrix phase comprises at least a primary matrix phase and a secondary matrix phase. The composite material system according to any one of claims 1 to 46, wherein the structure has a closed region without the matrix region. The composite material system according to any one of claims 1 to 47, wherein the structure is contained in the matrix phase. The composite material system according to any one of claims 1 to 48, wherein the structure is encapsulated in the matrix phase, and no part of the structure exists beyond the outer boundary of the matrix layer. At least a portion of the three-dimensional geometry is a tetrahedral tetrahedron, Weaire-Phelan shape, honeycomb shape, auxetic shape, octet truss shape, octahedron, diamond lattice, 3D cage shape, square shape, cube shape, tetrahedron shape, space. The composite material system according to any one of claims 1-49, characterized as a filled polyhedron, a periodic minimum surface, a triple periodic minimum surface shape, a spinodal shape, a chiral shape, or a combination thereof. One of claims 1-50, wherein the three-dimensional geometry is characterized by a beam or shell-based geometry, and the beam or shell-based geometry is not symmetrical, periodic, or periodically tessellated. The composite material system described in. The composite material system according to any one of claims 1 to 51, wherein the feature comprises one or more struts, beams, ties, trusses, sheets, surfaces, spheres, ellipsoids, and shells. The composite material system according to any one of claims 1 to 52, wherein the structure is characterized by a relative density selected from the range of 5% to 99.9%. The composite material system according to any one of claims 1 to 53, wherein the structure is characterized by a relative density selected from the range of 8% to 60%. The three-dimensional geometry is characterized by a plurality of features, wherein at least a portion of the feature independently has one or more cross-sectional physical dimensions selected from the range of 50 nm to 200 μm. The composite material system according to any one of 54. The composite material system according to any one of claims 1 to 55, wherein at least a portion of the features is characterized by one or more average longitudinal physical dimensions selected from the range of 10 nm to 2000 μm. The composite material system according to any one of claims 1 to 56, wherein the structure is characterized by elasticity and the elasticity of the structure is deterministic. Any of claims 1-57, wherein the structure is substantially undamaged by collisions from SiO ₂ particles having a diameter selected from 7 μm to 14 μm and a velocity selected from 500 m / sec to 1100 m / sec. The composite material system according to one item. The composite material system according to any one of claims 1 to 58, wherein the structure is characterized as having a bending control mode. The composite material system according to any one of claims 1 to 59, wherein the structure is characterized as having a stretch control mode. 29. The composite material system of claim 29, wherein the carbon isotope material is selected from the group consisting of carbon on glass, graphite carbon, amorphous carbon, pyrolytic carbon, fluffite, carbon black and combinations thereof. The composite material system of claim 61, wherein the carbon allotrope material comprises pyrolytic carbon. The composite material system according to any one of claims 1 to 62, wherein the structure comprises a coating. The composite material system of claim 63, wherein the coating comprises a metal, ceramic, or a combination thereof. The matrix phase according to any one of claims 1 to 64, wherein the matrix phase comprises one or more materials selected from the group consisting of polymers, epoxies, carbon allotropes, ceramics, metals, viscous fluids or combinations thereof. Composite material system. A method of making a composite material system, the method of which is
A step of preparing a structure by a laminating process, wherein the structure has a constructed 3D geometry, and the 3D geometry is monolithic and deterministic.
A step of impregnating the structure with the matrix phase, wherein the structure is at least partially impregnated with the matrix phase.
The steps to make the composite material system thereby
How to include. The method of claim 66, wherein the three-dimensional geometry is a nano or micro-constructed three-dimensional geometry. The method of claim 66 or 67, further comprising designing the three-dimensional geometry using computer-aided design technology. The method of claim 68, wherein the designing step comprises determining the three-dimensional geometry based on a calculated Pinodar decomposition. The step of designing is any of claims 66-69, comprising a step of determining at least one vibration frequency bandgap of the structure, wherein the at least one vibration frequency bandgap is deterministic. The method described in item 1. The method according to any one of claims 66 to 70, wherein the designing step is a step of determining the elasticity of the structure, and the elasticity of the structure is deterministic. The method according to any one of claims 66 to 71, wherein the designing step is a step of determining a bounce coefficient of the structure, and the bounce coefficient of the structure is deterministic. The designing step is a step of determining the area-normalized energy relaxation metric (Ψ) of the structure, and the area-normalized energy relaxation metric (Ψ) of the structure is deterministic, claim 66 to The method according to any one of 72. The laminated modeling process
Stereolithography technology (SLA)
Digital Optical Processing (DLP) Technology,
Continuous liquid interface manufacturing technology,
Micro Stereolithography (μ-SLA) Technology,
Two-photon polymerization lithography technology,
Interface lithography technology,
Holographic lithography technology,
Stimulated emission impairment (STED) lithography technology,
Other bat photopolymerization techniques,
Material extrusion technology,
Powder bed fusion technology,
Material injection technology, and combinations of these,
The method according to any one of claims 66 to 73, which is selected from the group consisting of. The method according to claim 75, wherein the laminated molding process is a three-dimensional lithography technique. The step of preparing the structure includes a step of forming a three-dimensional framework.
The step of preparing the structure further comprises a step of handling the three-dimensional framework to prepare the structure, and the step of handling comprises a thermal decomposition process according to any one of claims 66 to 75. the method of. The method of any one of claims 66-76, further comprising the step of applying a coating onto the structure. 77. The step of etching a portion of the structure such that the structure has at least one hollow feature, further comprising a step in which at least a portion of the at least one hollow feature comprises the coating. Method. The method of any one of claims 66-78, wherein the impregnating step comprises a process selected from the group consisting of sonication, vacuum exposure, and any combination thereof. The method of any one of claims 66-79, further comprising a step of post-treating the composite material system, wherein the post-treatment step comprises a curing process. The step of any one of claims 66-80, wherein the prepared step comprises selecting a precursor material selected from one or more resins, one or more metals, one or more ceramics and combinations thereof. the method of. 7. The method of claim 76, wherein the pyrolysis process is performed at a temperature selected from the range of 500 ° C. to 3000 ° C. and a duration selected from the range of 1 hour to 336 hours. The method of claim 76, wherein the pyrolysis process provides an isotropic shrinkage ratio selected from the range of 15% to 80% to the structure of the three-dimensional framework. The method according to any one of claims 66 to 83, further comprising the step of applying an external stimulus to the structure to change at least one vibration frequency bandgap of the structure. The method of any one of claims 66-84, which does not include the step of etching the template. Any of claims 66-85, wherein the structure is characterized by an area normalized collision energy relaxation metric (Ψ) selected from the range 2 × 10 ⁴ J / m ² to 4 × 10 ⁵ J / m ² . The method described in paragraph 1. One of claims 66-86, wherein the structure is characterized by a density normalized collision energy relaxation metric (Ψ) selected from the range 1.9 × 10 ⁶ J / kg to 4 × 10 ⁶ J / kg. The method described in paragraph 1. The method of any one of claims 66-87, wherein the structure is characterized by a bounce coefficient selected from the range 0.8-0.7. The method of any one of claims 66-88, wherein the structure is characterized by a bounce coefficient selected from the range 0.8-0.3. The method of any one of claims 66-89, wherein the composite material system is configured to substantially absorb collision energy by said structure. The method of any one of claims 66-80, wherein the structure is characterized by at least one vibration frequency bandgap. The method of claim 91, wherein the at least one vibration frequency bandgap is deterministic. The method of claim 91 or 92, wherein the at least one vibration frequency bandgap is in the range of 0.1 MHz to 200 MHz. The method of any one of claims 66-93, characterized by an attenuation ratio of at least 1.2%. The method of any one of claims 66-94, wherein the three-dimensional geometry has at least one surface feature. The method of any one of claims 66-95, wherein the three-dimensional geometry is characterized as spinodal geometry. The method of any one of claims 66-96, wherein the three-dimensional geometry comprises a resonator. The method of any one of claims 66-97, wherein the three-dimensional geometry is characterized by a unit cell geometry, wherein the unit cell comprises a resonator. The method according to any one of claims 66 to 98, wherein the three-dimensional geometry is a geometry without nodes.