JP7473126B2 - Aluminum-aluminum carbide composite molding and its manufacturing method - Google Patents

Description

The present invention relates to an aluminum-aluminum carbide composite molded body and a manufacturing method thereof. More specifically, the present invention relates to an aluminum-aluminum carbide composite molded body that can be suitably used for electronic parts and electric wires for aircraft and automobiles, and a manufacturing method thereof.

To realize a sustainable society, there is a need to reduce the environmental burden. In this context, light metal materials are attracting attention, and the use of light metal materials as conductive materials is no exception. One light metal material that is attracting attention is aluminum. However, the mechanical reliability of pure aluminum is low, and improvements are required. Previous research into improving the strength of aluminum has mainly focused on alloying and adding reinforcing materials.

Carbon materials are lightweight and have excellent properties such as strength, thermal conductivity, and electrical conductivity, and are attracting attention as a reinforcing material (composite material) for other materials. In addition, carbon materials with a graphene skeleton have been the subject of much research due to their unique structure and physical properties.

As an example of a composite of aluminum and a carbon material, a carbon-metal composite compact has been disclosed, which is made of a flaky carbon material having a graphene skeleton and metal particles, and which has a relative density of 90% or more and contains 25% or less by volume of carbon (see, for example, Patent Document 1 and Non-Patent Document 1). In the examples of Patent Document 1, a composite compact of aluminum and graphene oxide is disclosed.

Also disclosed is an aluminum-based composite material having an aluminum matrix and a rod-shaped or needle-shaped dispersion of aluminum carbide dispersed within the aluminum matrix (see, for example, Patent Document 2). Furthermore, disclosed is an aluminum-based composite material having an aluminum matrix and a dispersion dispersed within the aluminum matrix and formed by reaction of some or all of the additives with the aluminum in the aluminum matrix, the average particle size of the dispersion being 20 nm or less, the content of the dispersion being 0.25 mass % or more and 0.72 mass % or less in terms of carbon amount, and the distance between adjacent dispersions being 210 nm or less (see, for example, Patent Document 3).

In addition to the above, a composite of aluminum and graphene oxide has been disclosed (see, for example, Patent Document 4 and Non-Patent Document 2).

It has also been disclosed that aluminum powder and carbon powder are mixed to synthesize Al ₄ C ₃ , which is then subjected to heat treatment and used to reinforce an aluminum matrix (see, for example, Non-Patent Document 3).

Furthermore, an Al matrix composite reinforced with several layers of graphene has been disclosed, and hysteresis has been reported in the expansion behavior of the composite under repeated thermal loading in the temperature range of 323 to 573 K (see, for example, Non-Patent Document 4).

JP 2019-070186 A JP 2015-227498 A JP 2019-077901 A International Publication No. 2017/031403

Weiwei Zhou et al., Composites Part A: Applied Science and Manufacturing, 112 (2018) 168-177 Yunya Zhang et al., Nano Lett, 2017, 17, 6907-6915 A. Santos-Beltran et al., Materials Characterization, 106 (2015) 368-374 Weiwei Zhou et al., Journal of Alloys and Compounds, 792 (2019) 988-993

As mentioned above, there are various reports on composites of aluminum and carbon materials, but there is room for improvement in order to make the composites obtained in this way have excellent mechanical strength while maintaining sufficient electrical conductivity. The reason for this is thought to be that, for example, in the inventions described in Patent Document 4 and Non-Patent Documents 2 and 4, graphene oxide or graphene remains in the composite, but strong interfacial adhesion between the graphene oxide or graphene and aluminum is not achieved, load transmission at the interface is insufficient, and the inherent strength of graphene is not fully utilized.

The present invention was made in consideration of the above-mentioned current situation, and aims to provide a material that is lightweight yet has both excellent electrical conductivity and excellent mechanical strength, making it suitable for use, for example, in electronic components and electrical wires in aircraft and automobiles.

The inventors conducted extensive research into a material that is lightweight yet has both electrical conductivity and mechanical strength, making it suitable for use in electronic components and electrical wires in aircraft and automobiles. They came to the conclusion that the material can be obtained by pressurizing a lightweight carbon material and particles of aluminum, a light metal, at a temperature exceeding 600°C, thereby successfully resolving the above-mentioned problems, and thus arrived at the present invention.

That is, the present invention is a method for producing an aluminum-aluminum carbide composite molded body, which is characterized by including a step of pressing a flaky carbon material having a graphene skeleton and aluminum particles at a temperature exceeding 600°C.

In the examples of Patent Documents 2 and 3, a pressure of 600 MPa is applied during molding to prepare a powder compact, but there is no disclosure of applying pressure when subsequently heating the powder compact. Furthermore, Patent Document 4 and Non-Patent Document 2 disclose that heating is performed after applying a load of 50 kN, and again there is no disclosure of heating under pressure. In the inventions described in Patent Document 4 and Non-Patent Document 2, graphene oxide remains in the composite, as described above.

The manufacturing method of the present invention for aluminum-aluminum carbide composite moldings makes it possible to obtain moldings that are lightweight yet have both excellent electrical conductivity and excellent mechanical strength.

Figure 1(a) shows a TEM image of raw graphene oxide (GO), Figure 1(b) shows the AFM analysis of raw GO, Figure 1(c) shows an FESEM image of GO placed on silica glass, and Figure 1(d) shows the statistical results of the lateral size distribution of GO observed by FESEM. Figure 2(a) shows the Raman spectra of raw GO, GO/Al mixed powder, and PCPS composite (sintered or compact) sintered at 913 K. Figure 2(b) shows the weight loss of GO powder from 288 K to 913 K measured by TGA. 3(a) and 3(b) are FESEM images showing 0.2 wt% GO/Al mixed powder, and (c) and (d) are FESEM images showing 0.4 wt% GO/Al mixed powder. 4(a)-(d) are TEM images of the cross-sections of the PCPS-treated 0.4 wt% GO/Al composites. Figure 5(a) is a graph showing the statistical results of the diameter and length of _Al4C3 nanorods measured from 100 TEM images _{, respectively. Figure 5(b) is a graph showing the statistical results of the length and diameter of Al4C3 nanorods measured} from 100 TEM images. 6(a) is an EBSD-IPF map of pure Al. FIG. 6(b) is an EBSD-IPF map of a 0.32 vol.% Al ₄ C ₃ /Al composite. FIG. 6(c) is an EBSD-IPF map of a 0.64 vol.% Al ₄ C ₃ /Al composite. FIG. 6(d) is a graph showing the grain size distribution of pure Al. FIG. 6(e) is a graph showing the grain size distribution of a 0.32 vol.% Al ₄ C ₃ /Al composite. FIG. 6(f) is a graph showing the grain size distribution of a 0.64 vol.% Al ₄ C ₃ /Al composite. 7(a)-(c) are TEM images of the longitudinal section of a hot extruded 0.64 vol.% Al ₄ C ₃ /Al composite. Fig. 8(a) is a graph showing the nominal tensile stress-tensile strain curves of pure Al, 0.32 vol% Al ₄ C ₃ /Al composite, and 0.64 vol% Al ₄ C ₃ /Al composite, while Fig. 8(b) is a graph showing the theoretical calculation results of the composite strength according to each strengthening mechanism of pure Al and Al ₄ C ₃ /Al composite. FIG. 9 is a graph showing the cyclic thermal expansion behavior of a 0.64 vol. % Al ₄ C ₃ /Al composite.

The present invention will be described in detail below.
In addition, a combination of two or more of the individual preferred embodiments of the present invention described below is also a preferred embodiment of the present invention.

<Method of manufacturing aluminum-aluminum carbide composite molded body of the present invention>
The method for producing an aluminum-aluminum carbide composite molded product of the present invention includes a step of pressing a flaky carbon material having a graphene skeleton (hereinafter also simply referred to as a carbon material) and aluminum particles at a temperature exceeding 600°C.
As a result, the raw material, a flaky carbon material having a graphene skeleton, is converted into aluminum carbide, and the interface between the aluminum carbide and the aluminum matrix within the compact is very stable, making it possible to obtain an aluminum-aluminum carbide composite compact that has sufficient electrical conductivity and excellent mechanical strength.

The pressure in the pressurizing step may be any pressure that exceeds normal pressure (atmospheric pressure). The pressure is preferably 0.2 MPa or more, more preferably 0.5 MPa or more, even more preferably 1 MPa or more, even more preferably 5 MPa or more, even more preferably 10 MPa or more, even more preferably 20 MPa or more, and particularly preferably 30 MPa or more.
The upper limit of the pressure is not particularly limited, but is preferably 500 MPa or less, more preferably 300 MPa or less, further preferably 200 MPa or less, and particularly preferably 100 MPa or less.

The temperature in the above-mentioned pressurizing step is preferably 610°C or higher, more preferably 620°C or higher, even more preferably 630°C or higher, and particularly preferably 640°C or higher, from the viewpoint of the reaction of producing _Al3C4 by the reaction between the flaky carbon material having a graphene skeleton and the _aluminum particles.
The upper limit of the temperature is not particularly limited, but is preferably 670° C. or less, and more preferably 660° C. or less.
That is, the particularly preferred range of the temperature in the pressurizing step is 640 to 660°C.

Furthermore, the above temperature is preferably lower than the melting point (Tm) of the aluminum particles used. More preferably, it is at least 10°C lower than the melting point of the aluminum particles used ([Tm-10]°C or lower).

The above-mentioned pressurizing step is more preferably a step of sintering by pulse current pressure sintering (PCPS).
In particular, pulse current pressure sintering makes it possible to uniformly sinter the entire material, and even if the pressure process is performed at a high temperature that makes the carbon material more mobile, it is possible to produce a denser molded body while fully preventing agglomeration, resulting in a remarkable effect of excellent strength.

The time for carrying out the pressurizing step is preferably 1 minute to 1 hour, more preferably 2 minutes to 30 minutes, and further preferably 3 minutes to 20 minutes.
The pressurizing step may be carried out in air or in an atmosphere of an inert gas such as nitrogen, helium, or argon.

(Flake-like carbon material having a graphene skeleton)
The flaky carbon material having a graphene skeleton used in the production method of the present invention is not particularly limited as long as it has carbon (C) bonded by ^sp2 bonds and has a layered structure in which the carbon is arranged in a flaky shape (planarly). Specifically, the carbon material may have a structure consisting of only a layer (single layer) in which carbon atoms bonded by ^sp2 bonds are arranged in a planar shape, or may have a structure in which two or more layers are stacked. The carbon material preferably has, for example, a single layer structure consisting of only one layer, or a structure in which about 2 to 100 layers are stacked.

The carbon material preferably has carbon bonded to oxygen (O). For example, the carbon material is more preferably graphite oxide. More preferably, the carbon material is graphene oxide (also referred to as GO in this specification) in which oxygen is bonded to the carbon of graphene.
That is, in the production method of the present invention, the flaky carbon material having a graphene skeleton is preferably graphene oxide.
In general, graphene refers to a sheet consisting of one layer in which carbon atoms bonded by ^sp2 bonds are arranged in a plane, and a stack of many graphene sheets is called graphite. However, the above-mentioned graphene oxide includes not only a sheet consisting of a single layer, but also a sheet having a structure in which about 2 to 100 layers are stacked.

It is preferable that the carbon material is, for example, graphite oxide, which is a raw material for the composite, is heated or otherwise reduced during the composite formation, and becomes reduced graphite oxide in the aluminum-aluminum carbide composite molded product obtained by the production method of the present invention. That is, the aluminum-aluminum carbide composite molded product obtained by the production method of the present invention preferably contains reduced graphite oxide. More preferably, it contains reduced graphene oxide.
The carbon material may further have a functional group such as a carboxyl group, a hydroxyl group, a sulfur-containing group, or an alicyclic epoxy group.

Among the above carbon materials, those having a structure in which about 2 to 50 layers are laminated are preferred, those having a structure in which about 2 to 15 layers are more preferred, and those having a structure in which about 2 to 9 layers are further preferred. Those having a structure in which about 2 to 9 layers are laminated are also called FLG (few layers of graphene). Note that FLG may be a few layers of graphene, a few layers of graphene oxide, or a few layers of reduced graphene oxide, but it is preferably a few layers of graphene oxide or a few layers of reduced graphene oxide.
The number of layers of sheets in which carbon atoms bonded with ^sp2 bonds are arranged in a plane is expressed as the number of intensity peaks with respect to distance in a cross-sectional view of a TEM image.

The carbon material preferably has a specific surface area of 1 m ² /g or more, more preferably 10 m ² /g or more. A carbon material with such a specific surface area can exhibit higher dispersibility when composited with aluminum particles. The upper limit of the specific surface area is not particularly limited, but it can be, for example, 2000 m ² /g or less.
The specific surface area can be measured by a nitrogen adsorption BET method using a specific surface area measuring device.

The carbon material preferably has an average particle size of 1000 μm or less, and more preferably has an average particle size of 5 nm or more.
The average particle size can be measured by a particle size distribution measuring device.
Particles having an average particle size in the above-mentioned range can be produced, for example, by a method in which particles are pulverized using a ball mill or the like, the resulting coarse particles are dispersed in a dispersant to obtain a desired particle size, and then dried and solidified, or by a method in which the coarse particles are sieved or the like to select the particle size, or by a method in which preparation conditions are optimized in the particle production stage to obtain particles having a desired particle size.

The carbon material preferably has an average diameter of the flake surface (thin flake main surface) of 0.1 μm or more, more preferably 0.3 μm or more, further preferably 0.5 μm or more, and particularly preferably 1.0 μm or more.
The average diameter of the above-mentioned thin surface is preferably 15 μm or less, more preferably 10 μm or less, further preferably 8 μm or less, and particularly preferably 5 μm or less.
The average diameter of the above-mentioned flake surface can be calculated from an image of a field emission scanning electron microscope (FESEM). It is preferable to measure 20 or more points at an appropriate magnification and use the number-based average value obtained from the measurement results.

The above carbon material can be obtained by treating graphite with a known oxidizing agent, for example, by treating it with potassium permanganate using the modified Hummers method, and, if necessary, by performing a solid-liquid separation process such as ultrasonic treatment or centrifugation in a solvent.

In the manufacturing method of the present invention, the amount of the flaky carbon material having a graphene skeleton used (preferably, the amount added) is preferably 10% by mass or less relative to 100% by mass of aluminum particles. If the amount of the carbon material used exceeds 10% by mass, the crystal growth of aluminum is suppressed, and the electrical conductivity may not be very good due to the influence of the grain boundaries. In addition, the mechanical strength and elongation of the molded body obtained by the manufacturing method of the present invention may not be very good. The amount used is more preferably 5% by mass or less, even more preferably 3% by mass or less, even more preferably 1% by mass or less, even more preferably 0.8% by mass or less, and particularly preferably 0.5% by mass or less. As a result, the molded body obtained by the manufacturing method of the present invention has better electrical conductivity in that the influence of the grain boundaries is reduced.
The amount of use is not particularly limited to a lower limit, but from the viewpoint of exerting mechanical strength by the presence of carbon and by appropriately reducing the crystal grain size of aluminum due to the presence of carbon in the molded body obtained by the manufacturing method of the present invention, and from the viewpoint of making the electrical conductivity excellent due to the carbon material reduced during molding, for example, it is preferably 0.01 mass% or more, more preferably 0.05 mass% or more, and even more preferably 0.1 mass% or more. In addition, the lubricity of the carbon material can also make the sliding properties of the molded body excellent.

(Aluminum particles)
The aluminum particles used in the production method of the present invention may be particles containing aluminum as a primary component (the largest component in terms of mass proportion) and/or an alloy.
In the production method of the present invention, the aluminum particles are preferably composed of elemental aluminum and/or an alloy.

When the aluminum particles are an alloy, the other elements used together with the aluminum are not particularly limited, but are preferably at least one selected from the group consisting of magnesium, silicon, titanium, nickel, copper, iron, zirconium, silver, and tin, and more preferably at least one selected from the group consisting of magnesium, titanium, nickel, copper, iron, silver, and tin.
The aluminum particles may further contain silicon or zirconium as an additive. The other elements constituting the alloy particles according to the present invention are preferably crystalline.
The aluminum particles are more preferably composed of simple aluminum.

The melting point of the aluminum particles varies depending on the type and purity of impurities, but is preferably 630° C. or higher from the viewpoint of the reaction with the carbon material to produce Al ₄ C ₃ .
The melting point of the aluminum particles is more preferably 640° C. or higher.
The melting point of the aluminum particles is not particularly limited, but is usually 670° C. or lower.
That is, it is particularly preferable that the aluminum particles have a melting point of 640 to 670°C.

The aluminum particles preferably have an average particle size of 1.0 μm or more. More preferably, the average particle size is 3.0 μm or more. If the average particle size is less than 1.0 μm, the specific surface area increases, and the amount of Al ₂ O ₃ on the particle surface increases, which is undesirable from the viewpoint of making the effects of the present invention significant.
The upper limit of the average particle size is not particularly limited, but is usually 100 μm or less, and preferably 20 μm or less.
In this specification, the average particle size of the aluminum particles is a volume-based average particle size measured by a laser diffraction scattering method.

As described above, the aluminum particles may contain impurities other than elemental aluminum and/or aluminum alloys.
The impurities include those typically contained in commercially available aluminum particles.
The content of impurities is preferably 5% by mass or less, more preferably 2% by mass or less, further preferably 1% by mass or less, and particularly preferably 0.3% by mass or less, based on 100% by mass of the particles.

The manufacturing method of the present invention may be a composite using one type each of carbon material and aluminum particles, or may be a composite using two or more types. In addition, other components other than the carbon material and aluminum particles may be included, but the content of other components is preferably 3 mass% or less, more preferably 1 mass% or less, and even more preferably 0.1 mass% or less, based on 100 mass% of the raw material for the molded body. In particular, it is particularly preferable that the raw material for the carbon metal composite molded body of the present invention does not substantially contain other components.

The carbon metal composite molding of the present invention can be used in a variety of applications, but from the viewpoint of achieving weight reduction of the material, it is preferable to use it as, for example, electronic parts and electric wires for aircraft and automobiles.

The method of the present invention preferably further comprises a step of mixing the flaky carbon material having a graphene skeleton with aluminum particles before the pressurizing step, and more preferably obtains a composite powder of the carbon material and the aluminum particles by the mixing step.
The mixing step can be carried out, for example, by dropping the aqueous dispersion of the carbon material into water in which aluminum particles are dispersed, and filtering and drying the composite powder of the carbon material and aluminum particles obtained. Note that the water in which aluminum particles are dispersed is preferably prepared by stirring the aluminum particles and water, subjecting them to ultrasonic treatment, or the like, to more uniform mixing.

In the process of preparing a composite powder of aluminum particles and the carbon material by the above-mentioned mixing step, methods for bonding the aluminum particles and the carbon material include ball mill mixing, the use of a surface modifier, and the use of electrostatic interaction. Among them, from the viewpoint of improving electrical conductivity while avoiding damage to the carbon material, a method other than ball mill mixing, for example, a method using electrostatic interaction, is particularly preferred. For example, aluminum particles having an oxide coating that is positively charged in water (the point at which the zeta potential in a 25°C aqueous solution becomes 0 (isoelectric point) is pH>7) may be used in order to electrostatically interact with the carbon material that is negatively charged in water due to having ether groups or carbonyl groups.

The manufacturing method of the present invention preferably further includes a step of applying pressure to the molded body obtained in the pressing step to deform it. By the manufacturing method of the present invention, the aluminum particles in the molded body have a high aspect ratio, and it is possible to obtain a molded body having excellent strength in a specific direction.
In this specification, the step of deforming by applying pressure is not particularly limited, but is preferably a step of applying pressure to a material and stretching it in at least one direction, and examples thereof include a rolling step of stretching a material into a film shape using rollers, an extrusion step using a mold or the like, a step of forming a material into a film shape using a plate press or the like, an injection molding step, a casting step, and the like.

The pressure in the above-mentioned step of applying pressure to cause deformation is not particularly limited, but is preferably 0.1 kN/ ^mm2 or more, more preferably 0.5 kN/ ^mm2 or more, and even more preferably 1 kN/ ^mm2 or more.
Moreover, the pressure is preferably ⁵⁰ kN/mm2 or less, and more preferably 10 kN/ ^mm2 or less.

The temperature in the step of applying pressure to cause deformation is preferably 200° C. or higher, more preferably 300° C. or higher, and even more preferably 400° C. or higher. The temperature is preferably 2000° C. or lower, more preferably 1000° C. or lower, and even more preferably 700° C. or lower.
The step of applying pressure to cause deformation may be carried out in air or in an atmosphere of an inert gas such as nitrogen, helium, or argon.

Among the above-mentioned processes for applying pressure to cause deformation, the extrusion process and rolling process are preferred. The rolling process is particularly preferred in terms of its excellent productivity.

In the following, the extrusion process will be described in detail as an example.
The extrusion load in the extrusion step is preferably 100 kN or more, more preferably 200 kN or more, and even more preferably 300 kN or more.
The pushing load is preferably 1000 kN or less, and more preferably 800 kN or less.

The extrusion speed in the extrusion step is preferably 0.01 m/h or more, more preferably 0.02 m/h or more, and even more preferably 0.04 m/h or more.
The extrusion speed is preferably 1 m/h or less, more preferably 0.5 m/h or less, and even more preferably 0.2 m/h or less.

The extrusion ratio in the extrusion step is preferably 1 or more, more preferably 5 or more, and even more preferably 10 or more.
The extrusion ratio is preferably 200 or less, more preferably 100 or less, and even more preferably 40 or less.
The extrusion ratio is the value obtained by dividing the cross-sectional area of the die on the side pressing the material by the cross-sectional area of the die from which the material comes out in a cross section perpendicular to the extrusion direction of the material.

The extrusion process is preferably, for example, a hot extrusion process. The preferred temperature range in the hot extrusion process is the same as the preferred temperature range in the deformation process.

The preferred form of the aluminum-aluminum carbide composite molded body obtained by the method for producing an aluminum-aluminum carbide composite molded body of the present invention is the same as the preferred form of the aluminum-aluminum carbide composite molded body of the present invention described below.

(Aluminum-aluminum carbide composite molding of the present invention)
The present invention also relates to an aluminum-aluminum carbide composite molded product obtained by using the production method of the present invention.
The aluminum-aluminum carbide composite molded product of the present invention preferably contains 0.01 vol % or more and 5.0 vol % or less of aluminum carbide based on 100 mass % of the molded product.
The aluminum carbide content can be determined by measuring the C content from carbon/sulfur analysis and CHN (Elemental analysis (Carbon, Hydrogen, Nitrogen)) analysis of the molded body of the present invention, and by confirming that no D band peak or G band peak is shown in a spectrum chart obtained by Raman spectroscopy, as described below, after confirming that substantially all of the raw carbon material has been converted to Al ₄ C ₃ , the Al ₄ C ₃ content can be calculated from the C content.
The aluminum-aluminum carbide composite molded product of the present invention contains aluminum carbide, and therefore has sufficient electrical conductivity and also has excellent mechanical strength.
Furthermore, the aluminum carbide content in the aluminum-aluminum carbide composite molded product of the present invention is more preferably 0.1% by volume or more, and even more preferably 0.2% by volume or more, and more preferably 2.0% by volume or less, and even more preferably 1.0% by volume or less.

In the aluminum-aluminum carbide composite molded product of the present invention, the content of the flaky carbon material having a graphene skeleton is preferably 0.1 mass% or less in 100 mass% of the molded product. If the content of the carbon material exceeds 0.1 mass%, the electrical conductivity may not be very good due to the influence of the crystal grain boundaries. Furthermore, the mechanical strength and elongation of the aluminum-aluminum carbide composite molded product of the present invention may not be very good.
It is more preferable that the aluminum-aluminum carbide composite molded product of the present invention does not contain a flaky carbon material having a graphene skeleton.
The absence of a flaky carbon material having a graphene skeleton can be confirmed by confirming that a spectrum chart obtained by Raman spectroscopy does not show a D band peak or a G band peak, as described below.

It is preferred that the aluminum-aluminum carbide composite molded product of the present invention does not exhibit a D band peak in the range of 1250 to 1500 cm ^-1 and does not exhibit a G band peak in the range of 1500 to 1650 cm ^-1 in a spectrum chart obtained by Raman spectroscopy.
The phrase "does not show a D band peak within the range of 1250 to 1500 ^cm " means that a peak attributed to a D band having a peak top within the range of 1250 to 1500 ^cm is not present or does not show a height three or more times the average fluctuation width of the noise height of the spectrum chart.
The same applies to "not exhibiting a G band peak within the range of 1500 to 1650 cm ^-1 ".
The Raman spectroscopy is carried out under the conditions described in the Examples.

The aluminum-aluminum carbide composite molding of the present invention preferably has an aluminum particle content of 95 to 99.99% by volume. The metal particle content is more preferably 97 to 99.9% by volume, even more preferably 98 to 99.8% by volume, and even more preferably 99 to 99.7% by volume. It is particularly preferably 99.2% by volume or more.

The average grain size of the minor axis diameter of the metal derived from the aluminum particles in the aluminum-aluminum carbide composite molded product of the present invention (for example, the grain size when observing a plane perpendicular to the pressure direction in the deformation step in a molded product obtained by applying pressure to the molded product obtained in the pressurizing step) is preferably 1.0 μm or more and 5.0 μm or less, and more preferably 1.5 μm or more and 4.0 μm or less.
When the average grain size is 1.0 μm or more, the effect of excellent strength can be made more remarkable.
The average grain size can be measured by EBSD as described later in the Examples.

The aluminum-aluminum carbide composite molded product of the present invention is a product in which a carbon material is composited with aluminum particles, and examples of such molded products include pellet-like molded products obtained by pressing a carbon material and aluminum particles at a temperature exceeding 600° C., and molded products obtained by applying pressure to such molded products. The pellet-like molded products obtained by pressing have excellent electrical conductivity and mechanical strength, and can be deformed by applying pressure to such molded products.
The shape and size of the aluminum-aluminum carbide composite molded product of the present invention are not particularly limited, and examples thereof include pellets, as well as films, sheets, fibers, columns, cubes, spheres, and the like.

The aluminum-aluminum carbide composite molded body obtained by the preferred embodiment of the method for producing an aluminum-aluminum carbide composite molded body of the present invention described above is also a preferred form of the aluminum-aluminum carbide composite molded body of the present invention. For example, the present invention is also an aluminum-aluminum carbide composite molded body obtained by using a flaky carbon material having a graphene skeleton and aluminum particles, characterized in that the amount of the flaky carbon material having a graphene skeleton used is 1 mass % or less.

The aluminum-aluminum carbide composite molding of the present invention is lightweight yet has excellent electrical conductivity and mechanical strength, making it suitable for use, for example, in electronic components and electrical wires for aircraft and automobiles.

The present invention will be described in more detail below with reference to examples, but the present invention is not limited to these examples. Unless otherwise specified, "parts" means "parts by weight" and "%" means "% by mass."

Various characteristic values were measured by the following methods.
(TEM Observation)
The TEM observation of the molded body (PCPS sintered body) in each example was performed using a field emission transmission electron microscope (HF-2000EDX, manufactured by Hitachi, Ltd.).
(Electrical conductivity measurement)
The measurements were made using a thermoelectric property evaluation device (ZEM-3, manufactured by ULVAC, Inc.).
(Ultimate tensile strength [UTS] and tensile elongation of molded body)
The measurement was carried out using a precision universal testing machine (AUTOGRAPH AG-I 50KN) manufactured by Shimadzu Corporation.

<Example>
<Preparation example of graphene oxide (GO) colloidal solution>
A mixture was prepared by adding 28.75 parts of concentrated sulfuric acid (special grade reagent, manufactured by Wako Pure Chemical Industries, Ltd.) and 1.00 parts of natural graphite (Z-5F, flake graphite, manufactured by Itoh Graphite Industries, Ltd.) to a corrosion-resistant reactor. While stirring the mixture, potassium permanganate (special grade reagent, manufactured by Wako Pure Chemical Industries, Ltd.) was added 20 times into the mixture at 15-minute intervals. The amount of potassium permanganate added each time was 0.15 parts, and the total amount added was 3.00 parts. After the addition of potassium permanganate, the mixture was heated to 35° C. and aged for 2 hours while maintaining the liquid temperature to obtain a reaction slurry (slurry containing graphite oxide). Next, 20.00 parts of the reaction slurry was added under stirring to another container containing 75.58 parts of ion-exchanged water, and 1.08 parts of 30% hydrogen peroxide water (special grade reagent, manufactured by Wako Pure Chemical Industries, Ltd.) was added, and the stirring was continued for 30 minutes before being stopped. After stopping the stirring, the mixture was left to stand overnight to separate into a precipitate layer and a supernatant layer, and the supernatant was then removed. Thereafter, the same volume of ion-exchanged water as the removed supernatant was added to the container to wash the precipitate, and the mixture was stirred for 30 minutes, after which the stirring was stopped and the mixture was left to stand for 5 hours or more, and the supernatant was again removed. The washing with the addition of ion-exchanged water and the removal of the supernatant were repeated until the pH of the supernatant reached 2 or more. An appropriate amount of ion-exchanged water was added to the obtained precipitate layer, and graphite oxide was dispersed by homogenizer treatment to obtain a GO colloidal solution (graphene oxide colloidal solution). The GO concentration in the obtained GO colloidal solution was 1.0%.

(1) Preparation of raw material and GO/Al mixed powder The raw material GO aqueous colloid with a concentration of 1.0 mass% was synthesized by the above method. Al powder with an average particle size of 5.5 μm and a purity of about 99.85% (also referred to as pure Al powder in this specification, meaning that it is not composited with a filler such as GO) produced by gas atomization was provided by Ecka Granules Japan.
12 g of Al powder was put into a beaker containing 200 mL of deionized water. Al colloids were obtained by ultrasonic treatment while stirring in an ice-water bath for 2 hours. At the same time, 2.4 g or 4.8 g of GO colloids were diluted and dispersed in 100 ml or 200 ml of deionized water under ultrasonic treatment to obtain GO suspensions . The GO suspension was then dropped into the Al colloids and stirred for 0.5 hours. After filtration, the mixture was dried at 343 K under vacuum to obtain 0.2% by mass and 0.4% by mass of GO/Al mixed powders.

(2) Preparation of Al ₄ C ₃ /Al matrix composite The GO/Al mixed powder was sintered at a high temperature of 913 K and an applied pressure of 50 MPa for 0.33 h by pulse current pressure sintering (PCPS; Dr. Sinter S511, Sumitomo Metal Mining Co., Ltd.) to synthesize an Al ₄ C ₃ /Al composite. The heating rate of the PCPS was 1200 K/h. The PCPS sintered composite was then processed by a hot extruder (UH-500kN1, Shimadzu Corporation) at 773 K with an extrusion speed of 0.06 m/h and an extrusion ratio of about 20. The average diameter of the cylindrical extruded bar was about 3.4 mm. For comparison, pure Al specimens were also prepared using the same powder mixing and sintering process.

(3) Properties The layer thickness of raw GO was measured by atomic force microscope (AFM; Hitachi SII Nanocute). The crystalline disorder of GO was evaluated by Raman spectroscopy (SOLAR TII Nanobinder, Tokyo Instruments). The thermal stability of GO was confirmed by thermogravimetric analysis (TGA, SDT Q600, Hitachi) in Ar atmosphere using the same conditions of PCPS (e.g., heating rate 1200 Kh ^-1 , maximum temperature 913 K). The apparent density of the bulk sample was determined using Archimedes' principle. The powder or bulk morphology was investigated by field emission scanning electron microscopy (FESEM; JSM-6500F, JEOL Ltd.), electron backscatter diffraction (EBSD; OIM ver.6, TSL Solutions Ltd.) and field emission transmission electron microscopy (TEM; HF-2000EDX, Hitachi Ltd.). Cyclic thermal expansion behavior was tested by thermomechanical analysis (TMA; Q400, TA Instruments Co., Ltd.) in the temperature range of 318 K-573 K with a heating rate of 240 Kh ^-1 under argon gas flow, for 10 thermal cycles. Cylindrical specimens with a diameter of about 3 mm and a length of about 10 mm were prepared from the extruded bars. Vickers hardness tests were performed with a microhardness tester (HM-200, Mitutoyo Corp.). The tensile performance of the composites was evaluated by ^uniaxial tensile testing at a strain rate of 1.67× ¹⁰ s using a precision universal testing machine (AUTOGRAPH AG-I 50KN). Dog-bone shaped tensile specimens with a diameter of about 2.5 mm and a gauge length of about 10 mm were machined from the extruded bars. The electrical conductivity of the hot extruded bars along the extrusion direction was measured at room temperature using a thermoelectric property evaluation device (ZEM-3, ULVAC, Inc.).

<Results and Discussion>
(1) Characteristics of GO and GO/Al mixed powders Figure 1(a) shows a TEM image of raw GO. Figure 1(b) shows the AFM analysis results of raw GO. Figure 1(c) shows an FESEM image of GO placed on silica glass. Figure 1(d) shows the statistical results of the lateral size distribution of GO observed by FESEM. The inset in Figure 1(b) shows the height profile of GO along the white line shown in Figure 1(b).

The morphological features of raw GO are shown in Figure 1. The GO platelets are transparent and can be distinguished by TEM from the edges and wrinkles (Figure 1(a)). A small amount of carbonaceous nanoparticles (see white arrows in Figure 1(a)) resulting from the manufacturing process were located on the flexible GO sheets. AFM analysis showed the characteristics of few-layered (i.e., <10 layers) GO. For example, the GO in Figure 1(b) had a thickness of about 2.8 nm, consisting of four graphene layers, considering the GO interlayer distance of 0.7 nm. FESEM observations showed that the GO platelets had an irregular shape (Figure 1(c)). The lateral size of the GO ranged from 0.5 μm to 2.8 μm, with an average value of about 1.51 μm calculated from 20 FESEM images (Figure 1(d)).

Figure 2(a) shows the Raman spectra of raw GO, GO/Al mixed powder, and PCPS composite. The Raman spectrum was measured under the following conditions: excitation wavelength: 532 nm, objective lens: 20x, CCD capture time: 10 s, number of integrations: 5. Figure 2(b) shows the weight loss of GO powder from 288 K to 913 K measured by TGA. The inset in Figure 2(b) shows a schematic diagram of GO before and after thermal reduction.

The structural disorder of GO was revealed by Raman spectroscopy. This disorder is generally associated with two characteristic peaks for sp2-based carbonaceous materials, such as the defect (D) band at about 1350 cm ^-1 and the graphite (G) band at about 1580 cm ^-1 . As shown in Figure 2(a), the I _D /I _G ratio value of GO was measured to be about 0.95, which is attributed to the presence of abundant intrinsic defects such as edges, vacancies, dangling bonds, and atomic defects in the hexagonal C=C network. These disorders are expected to preferentially react with the Al matrix to form Al ₄ C ₃ during sintering (composite formation). No defect (D) band peak or graphite (G) band peak was observed in the PCPS sintered body prepared at 913 K. This is presumably due to the reaction of GO with Al in the composite to form Al ₄ C ₃ , and the disappearance of GO. The thermal stability of GO was also demonstrated by TGA analysis in Figure 2(b). It is noted that GO was unstable; the thermal decomposition of the surface groups started at low temperatures (~450 K) with the release of gaseous _H2O , CO, or _CO2 , and the residue corresponding to reduced GO (inset in Fig. 2(b)) remained at ~43.8% at the sintering temperature of 913 K.

Figures 3(a) and 3(b) are FESEM images of a 0.2 wt% GO/Al mixed powder, respectively. Figures 3(c) and 3(d) are FESEM images of a 0.4 wt% GO/Al mixed powder, respectively. Figures 3(b) and 3(d) show enlarged portions of the white dashed rectangles in Figures 3(a) and 3(c), respectively.

Figure 3 shows the morphology of GO/Al mixed powders. Due to the opposite surface charges in aqueous solution, GO sheets were adsorbed on the surface of Al powders by electrostatic self-assembly, achieving uniform dispersion (Figure 3(a)). It is noteworthy that the GO-wrapped structure was retained after powder drying as a result of the strong van der Waals forces. The thin, flexible, and high aspect ratio GO easily fit into the spherical Al powders without changing the original properties (e.g., particle size or shape) (Figure 3(b)). When the GO content was increased to 0.4 wt%, each Al particle was almost covered with GO (Figure 3(c)). Occasionally, a GO sheet bridged two adjacent Al powders (Figure 3(d)).

(2) Formation of Al ₄ C ₃ interface during sintering Figures 4(a)-(d) are TEM images of the cross-section of the PCPS-treated 0.4 wt% GO/Al composite. The inset in Figure 4(b) is a schematic showing the formation of cross-linked Al ₄ C _3. The inset in Figure 4(c) shows the diffraction of Al ₄ C ₃ measured from the round spot in Figure 4(c). Figure 4(d) is taken from the dashed square in Figure 4(c).
Figure 4 shows the microstructure of the PCPS-treated 0.4 wt% GO/Al composite. Many rod-like phases (pointed by black arrows in Figure 4(a)) were observed in the composite. The nanorods were single- _{crystalline Al4C3} with rhombohedral structure, as identified by the corresponding selected area electron diffraction (SAED) pattern (inset in Figure 4(c)). Since the raw GO surrounded the surface of the Al powder (Figure 3), the formed _{Al4C3 nanorods} were interposed at the grain boundaries (GB) rather than inside the Al grains during solid-state PCPS (Figure 4(a)).
The HRTEM image in Fig _. 4(d) revealed that the (003) atomic planes of _Al4C3 were continuous, very straight and parallel to each other, with a constant interplanar spacing of about 0.84 nm, suggesting high crystallinity. The vertical formation of _Al4C3 along the (003) planes was faster than the lateral growth, because the surface free energy in the graphite prism planes is higher than that in the basal planes. Correspondingly, _Al4C3 grew in rod-like shapes along the <110> direction (Fig _{. 4(c)). In addition, the Al4C3 crystals contain several} dislocation lines (pointed by arrows in Fig. 4(d)) that are nearly aligned in the <110> direction. This is attributed to the rapid interfacial reaction rate between the defective GO and Al atoms and the insufficient release of internal stress during _{the Al4C3} growth.
As has been well documented in carbon nanotube (CNT)/Al composites, Al atoms preferentially diffuse into CNTs to form Al ₄ C ₃ during high-temperature sintering or annealing without substantial changes to the original shape of CNTs. In contrast, Al ₄ C ₃ has an obviously different shape as raw GO in GO/Al composites. Since GO contains many carbon dangling bonds on the open edges or basal planes, C atoms can easily break the restriction of C=C covalent bonds. Therefore, the nucleation and growth of Al ₄ C ₃ should be related to the interdiffusion of C and Al atoms.

Figure 5(a) is a graph showing the statistical results of the diameter and length of _Al4C3 nanorods measured from 100 TEM images _{, respectively. Figure 5(b) is a graph showing the statistical results of the length and diameter of Al4C3 nanorods measured} from 100 TEM images.
According to the statistical results of 100 TEM images, the average diameter and length of _{the Al4C3} nanorods are about 74.6 nm (Fig. 5(a)) and about 301.4 nm (Fig. 5(b)), respectively, with an aspect ratio of about 4.0. Assuming that GO contains four layers of graphene (Fig. 1), it is calculated that one GO sheet can provide enough carbon source to synthesize about _{5.4 Al4C3} rods.

Figure 4(c) shows that Al ₄ C ₃ is well bonded to the Al matrix and exhibits a clear boundary. Five small faceted planes were clearly detected at the Al ₄ C ₃ /Al interface (Fig. 4(d)). The corresponding SAED pattern (inset of Fig. 4(c)) revealed that the faceted Al ₄ C ₃ planes have a low index and thus retain a relatively stable interface with the matrix. The internal compressive stress is estimated to be a residual stress caused by the mismatch of thermal expansion coefficients between Al ₄ C ₃ (3.6×10 ⁻⁶ K ⁻¹ pcs) and Al (2.5×10 ⁻⁵ K ⁻¹ pcs), which contributes to a clean and closed Al ₄ C ₃ /Al interface.

Through careful TEM observation, we found a remarkable phenomenon; an _{Al4C3 nanorod} (pointed by the thick rightward arrow in Fig. 4(b)) was embedded vertically into two Al particles, forming a bridged structure. This can be understood roughly by the inset in Fig. 4(b). Due to the surface cleaning effect of PCPS, _the amorphous _Al2O3 layer with a thickness of about 5 nm on the surface of the Al powder was partially destroyed (Fig. 4(c)). As a result, _a certain part of the GO sheet was brought into direct contact with the fresh Al surface for the nucleation of _{Al4C3 crystals} . Then, _Al4C3 grew from two opposite directions along the boundary, resulting in _{an Al4C3 bridge} across two adjacent Al grains. In addition to the increased diffusion of Al atoms near the boundary due to the high-energy _PCPS densification, such an _Al4C3 bridge is expected to further improve the bonding conditions between the Al grains, which enhanced the mechanical properties.

No residual GO sheets were measured in the composite by HRTEM (Figure 4). This result is supported by the micro-Raman analysis in Figure 2(a); the two characteristic GO peaks completely disappeared after PCPS. It was concluded that the composite fabricated was an _Al4C3 /Al composite. By considering that GO reacted completely _{to Al4C3} , _{the Al4C3} content in the composite is calculated to be about 0.32 vol% and about 0.64 _vol %, corresponding to 0.2 wt% GO/Al and 0.4 wt% GO/Al of the raw material, respectively (see Table 1).

Table 1 shows the relative density, tensile properties, Vickers hardness, and electrical conductivity of pure Al and Al ₄ C ₃ /Al composites produced by PCPS and hot extrusion. The apparent densities of pure Al, reduced GO, and Al ₄ C ₃ were 2.7 g cm ^-3 , 2.28 g cm ^-3 , and 2.93 g cm ^-3 , respectively. The Vickers hardness values were measured on the cross-section of the specimens.

6(a) is an EBSD-IPF map of pure Al. FIG. 6(b) is an EBSD-IPF map of a 0.32 vol.% Al ₄ C ₃ /Al composite. FIG. 6(c) is an EBSD-IPF map of a 0.64 vol.% Al ₄ C ₃ /Al composite. FIG. 6(d) is a graph showing the grain size distribution of pure Al. FIG. 6(e) is a graph showing the grain size distribution of a 0.32 vol.% Al ₄ C ₃ /Al composite. FIG. 6(f) is a graph showing the grain size distribution of a 0.64 vol.% Al ₄ C ₃ /Al composite.
Figures 7(a)-(c) are TEM images of the longitudinal section of a hot extruded 0.64 vol% Al ₄ C ₃ /Al composite, and Figure 7(c) is an enlarged view of the square area in Figure 7(b).

By combining PCPS and hot extrusion, all the pure Al and Al ₄ C ₃ /Al composites showed a high density of more than 99.7% (Table 1). Detailed crystallographic information was collected by EBSD inverse pole figure (IPF) maps in Fig. 6(a)-(c). Due to the plastic flow of Al during hot extrusion, the matrix rotated to form a structure with <111> oriented grains along the extrusion direction. Also, the average grain size values of the 0.32 vol% Al ₄ C ₃ /Al composite and the 0.64 vol% Al ₄ C ₃ /Al composite measured from the cross section perpendicular to the hot extrusion direction were 2.48 μm (Fig. 6(e)) and 2.02 μm (Fig. 6(f)), respectively, which were smaller than that of pure Al (2.90 μm; Fig. 6(d)).

TEM observation further confirmed the alignment of the hot-extruded Al crystal grains (Figure 7(a)). Individual _Al4C3 nanorods (indicated by arrows in Figure 7(b)) were almost unidirectionally aligned. The average distance _{between adjacent Al4C3 nanorods was approximately} several hundred nanometers. Furthermore, _Al4C3 maintained a close and faceted interface with the matrix, _which could act as an effective reinforcement agent for AMCs (Figure 7(c)).

(3) Mechanical and electrical performance of Al ₄ C ₃ /Al composites Figure 8(a) is a graph showing the nominal tensile stress-tensile strain curves of pure Al, 0.32 vol% Al ₄ C ₃ /Al composite, and 0.64 vol% Al ₄ C ₃ /Al composite, while Figure 8(b) is a graph showing the theoretical calculation results of the composite strength according to each strengthening mechanism of pure Al and Al ₄ C ₃ /Al composite.
Figure 8(a) shows the nominal tensile stress-strain curves of the Al ₄ C ₃ /Al composites, compared with pure Al processed under the same sintering conditions. The tensile strength of Al gradually increased with increasing Al ₄ C ₃ formation, and the elongation at break remained at a high level of >16%. The ultimate tensile strength (UTS) values of the 0.32 vol% Al ₄ C ₃ /Al composite and the 0.64 vol% Al ₄ C ₃ /Al composite were 166.2 MPa and 183.1 MPa, respectively, showing an improvement of about 24.2% and about 36.8% compared with that of pure Al (133.8 MPa). Similarly, the Vickers hardness of Al was significantly improved with the introduction of Al ₄ C ₃ (Table 1). As a result, the synthesized Al ₄ C ₃ nanorods showed a remarkable strengthening effect in AMC.

To understand the strength improvement of _Al4C3 /Al composites, a _{quantitative analysis} of the strengthening mechanisms was performed, mainly including (i) grain refinement, (ii) Orowan loop strengthening of _Al4C3 , and (iii) load transfer from the matrix to _{Al4C3 .} The strengthening effect from grain refinement is predicted by the Hall-Petch equation (1) as follows:
Δσ _GR = k (d _c ^−0.5 −d _m ^−0.5 ) (1)
where k is the constant 0.07 MPa·m ^0.5 for Al, and d _c and d _m are the average grain sizes of the composite and pure Al, respectively.

ここで、Ｍは、テイラー因子（面心立方金属についてはＭ＝３．０６）であり、Ｇは、マトリックスのせん断弾性率（Ｇ＝２５．４ＧＰａ）であり、ｂは、Ａｌのバーガース・ベクトル（ｂ＝０．２８６ｎｍ）であり、Ｄは、Ａｌ_４Ｃ_３の直径（Ｄ＝７４．６ｎｍ）であり、ＳはＡｌ_４Ｃ_３のアスペクト比（Ｓ＝約４．０）であり、ＶはＡｌ_４Ｃ_３の体積率である。 The strength improvement Δσ _OR due to Orowan loop strengthening is calculated by the following formula (2).

where M is the Taylor factor (M = 3.06 for face-centered cubic metals), G is the shear modulus of the matrix (G = 25.4 GPa), b is the Burgers vector of Al (b = 0.286 nm), D is the diameter _{of Al4C3} (D = 74.6 nm), _S is the aspect ratio of _Al4C3 (S = ~4.0), and V is the volume fraction _{of Al4C3} .

In uniaxially oriented short fiber reinforced AMC, the strength increase (Δσ _LT ) caused by the load transfer mechanism can be estimated from the shear lag theory shown in Equation (3) below.
Δσ _LT = 0.5 V _{σ m} (3)
Here, σ _m is the UTS of pure Al, which is 133.8 MPa. According to equations (1)-(3), Δσ _GR =3.3 MPa, Δσ _OR =27.1 MPa, Δσ _LT =0.9 MPa at V=0.32 vol.%; Δσ _GR =8.2 MPa, Δσ _OR =36.3 MPa, Δσ _LT =1.7 MPa at V=0.64 vol.%. It is concluded that the Orowan loop is the dominant strengthening mechanism in the Al ₄ C ₃ /Al composite synthesized in this example. Figure 8(b) summarizes the contribution of each strengthening mechanism in the Al ₄ C ₃ /Al composite. The predicted strengthening effect of Al ₄ C ₃ is in good agreement with the experimental value.

Considering electrical applications, the electrical conductivity of pure Al and Al ₄ C ₃ /Al composites was evaluated at room temperature, as shown in Table 1. Due to the presence of clean Al ₄ C ₃ -Al interface (FIG. 7(c)), the electrical conductivity of Al was slightly decreased to about 2.3% or about 7.6% by the introduction of 0.32 vol.% or 0.64 vol.% Al ₄ C ₃ , respectively. In general, the tensile strength and electrical conductivity of AMCs show a trade-off trend, i.e., the increase in strength comes at the expense of electrical conductivity and vice versa. Compared with widely used fillers (e.g., CNTs, ceramics, GPLs [graphene platelets], or heat-resistant nanoparticles), Al ₄ C ₃ has been proven to be more effective for fabricating high-performance AMCs with excellent electrical conductivity-strength combination. For example, the 0.64 vol.% Al ₄ C ₃ /Al composite had a high UTS of 183.1 MPa and electrical conductivity of 57.4% IACS (Table 1).

FIG. 9 is a graph showing the cyclic thermal expansion behavior of a 0.64 vol. % Al ₄ C ₃ /Al composite.
Recently, Zhou et al. combined the Hummer process with powder metallurgy techniques to fabricate a 0.4 vol.% few-layer graphene GPL/Al composite with a UTS of about 173 MPa and electrical conductivity of about 59.1% IACS (see Non-Patent Document 4). However, in Non-Patent Document 4, the GPL remains, and the interface bond between the GPL and the Al matrix is mechanically bonded, which may lead to instability of mechanical properties at high temperatures when applied to electric wires. To further confirm the effectiveness of the Al ₄ C ₃ /Al composite of this example, the thermal expansion behavior under cyclic thermal loading as shown in FIG. 9 was evaluated, and the hysteresis phenomenon of thermal expansion was clearly observed between heating and cooling in the GPL/Al composite, indicating that the interface was unstable. However, the cyclic thermal expansion behavior of the Al ₄ C ₃ /Al composite was linear and reversible in the temperature range of 318 K to 573 K (FIG. 9). This result suggests that interfacial slip does not actually occur at the Al ₄ C ₃ -Al interface. Therefore, it is believed that the Al ₄ C ₃ /Al composite of this example exhibits superior mechanical performance at high temperatures to the GPL/Al composite described in Non-Patent Document 4.

We proposed _a strategy to utilize few-layer GO to synthesize _Al4C3 /Al composites by high-energy sintering. GO sheets were uniformly coated on the surface of Al powder under electrostatic attraction through a _{heterocoagulation} process. During high-temperature PCPS, GO reacted with Al atoms and completely transformed into single-crystalline _Al4C3 nanorods with an average diameter of about 74.6 nm and a length of about 301.4 nm. Subsequent hot extrusion uniaxially rearranged the _{Al4C3 nanorods} individually, exhibiting an intimate and faceted interface with the matrix.
The tensile strength of the Al matrix was significantly enhanced by the introduction of Al ₄ C ₃ , mainly due to the Orowan loop strengthening of Al ₄ C ₃ . Furthermore, the Al ₄ C ₃ /Al composite retained high electrical conductivity of >57.4% IACS (about 92.4% of pure Al) and showed stable interfacial bonding at high temperatures. As a result, the Al ₄ C ₃ /Al composite has great application prospects, for example in electrical conductor applications.

Claims (5)

1. A method for producing an aluminum-aluminum carbide composite compact, comprising the steps of:
The production method includes a step of pressing a flaky carbon material having a graphene skeleton and aluminum particles at a temperature exceeding 600° C.,
The aluminum-aluminum carbide composite molded product is characterized in that it does not exhibit a D band peak in the range of 1250 to 1500 cm ^-1 and does not exhibit a G band peak in the range of 1500 to 1650 cm ^-1 in a spectral chart obtained by Raman spectroscopy. The method for producing an aluminum-aluminum carbide composite molding according to claim 1, characterized in that the flake-like carbon material having a graphene skeleton is graphene oxide. The method for producing an aluminum-aluminum carbide composite compact according to claim 1 or 2, characterized in that the aluminum particles are composed of aluminum alone and/or an alloy. The method for producing an aluminum-aluminum carbide composite compact according to any one of claims 1 to 3, characterized in that the production method further includes a step of applying pressure to the compact obtained in the pressing step to deform it. An aluminum-aluminum carbide composite molded body obtained by using a flaky carbon material having a graphene skeleton and aluminum particles,
the flaky carbon material having a graphene skeleton is graphene oxide, and the amount of the graphene oxide used is 1 mass % or less;
An aluminum-aluminum carbide composite molded product, characterized in that a spectrum chart obtained by Raman spectroscopy does not show a D band peak in the range of 1250 to 1500 cm ⁻¹ and does not show a G band peak in the range of 1500 to 1650 cm ⁻¹ .

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