JP2019157178A - Oxide ceramic base composite material, and manufacturing method therefor - Google Patents

Abstract

To provide an AlO-FeAl composite material excellent in high temperature oxidation resistance at 900°C or higher, and dense and excellent in mechanical properties such as hardness or toughness, by integrating and sintering a binding phase and a hard phase, especially an oxide ceramic group composite material having mechanical properties equal to cermet and excellent high temperature oxidation resistance compared to the cermet, and a manufacturing method therefor.SOLUTION: An oxide ceramic group composite material has a hard phase with alumina (AlO) hard particle as a base phase, and a binding phase containing iron aluminum (FeAl), average particle diameter of the alumina (AlO) hard particle is 0.2 μm to 0.6 μm, content of the hard phase containing alumina (AlO) is in a range of 50 vol.% and 75 vol.%, iron aluminum (FeAl) alloy contains aluminum (Al) at a percentage of more than 23 atom% and 50 atom% or less, and the binding phase and the hard phase are integrated and sintered.SELECTED DRAWING: None

Description

  The present invention relates to an oxide ceramic matrix composite material and a method for producing the same.

  WC-Co cemented carbide is used as a mold and tool material in the high precision machining field. WC-Co is a material in which tungsten carbide (WC) particles, which are hard phases, are baked and hardened using metallic cobalt (Co) as a binder phase, and is known as a composite material that achieves both high hardness and toughness. ing.

  Many technologies related to WC-Co cemented carbide have been proposed to date to improve the mechanical properties of cemented carbide. For example, Patent Document 1 proposes a tough cemented carbide having improved toughness, strength, and hardness by increasing the lengths of both the a-axis and the c-axis in the hexagonal tungsten carbide crystal axis. .

  However, WC-Co cemented carbide has a problem in high-temperature oxidation resistance, and when exposed to air at 900 ° C or higher, the oxidation of WC and Co becomes severe, making it unusable for use as a mold or tool material.

To improve high-temperature oxidation resistance, ceramics with excellent oxidation resistance such as titanium carbonitride (TiCN) and alumina (Al ₂ O ₃ ) are used on the surfaces of WC-Co cemented carbide molds and tools. It has been devised such as coating. However, when these coatings are lost due to wear, there is a problem that wear due to oxidation rapidly develops. For this reason, it is important that the material itself has high-temperature oxidation resistance.

  As a composite material excellent in high-temperature oxidation resistance, there is a cermet in which hard particles having titanium carbonitride (TiCN) as a matrix are combined with nickel. This composite material exhibits relatively excellent oxidation resistance even in the atmosphere at 900 ° C. while maintaining mechanical properties close to those of cemented carbide. However, since the oxidation gradually proceeds when kept in the atmosphere at a high temperature for a long time, the cermet may be insufficient in terms of durability.

In order to produce materials with excellent durability even in the atmosphere of 900 ° C or higher, it is necessary to use oxide ceramics such as Al ₂ O ₃ as the base material and to harden them with metals and alloys with excellent high-temperature oxidation resistance. It is valid. A candidate for such a composite material is a composite material in which an alloy containing aluminum is used as a binder phase and Al ₂ O ₃ is baked and hardened. Patent Document 2 proposes a composite material in which Al ₂ O ₃ is baked and solidified using a ductile Al compound with the aluminum (Al) content suppressed to 20 atomic% or less as a binder phase. Patent Document 3 proposes a composite material in which Al ₂ O ₃ is a hard phase and iron aluminum (FeAl) containing Al in an amount of 15.9 atomic% to 22.4 atomic% is a binder phase.

  However, when an Al compound with a low Al content, especially a compound with a metal with poor oxidation resistance such as iron (Fe), is used in the binder phase, the oxidation resistance at a high temperature of 900 ° C. or higher is greatly improved. It is not possible.

In order to improve the durability of the Al compound against high-temperature oxidation, it is necessary to form an oxide film on the surface, and for that purpose, it is important that the Al ratio in the Al compound is at least 23 atomic% or more. Non-Patent Document 1 reports a method of using a combustion synthesis of Fe and Al, forming FeAl during firing, and producing an Al ₂ O ₃ —FeAl composite material.

However, an Al ₂ O ₃ —FeAl composite material having sufficient hardness and strength has not been obtained by the method using combustion synthesis as described above. The mechanical properties such as hardness and strength of Al ₂ O ₃ -FeAl composites using combustion synthesis are low due to the poor wettability of the hard phase Al ₂ O ₃ and the binder phase FeAl, and the state of the liquid phase This is thought to be because the FeAl repels Al ₂ O ₃ and the binder phase and the hard phase are not sintered together.

On the other hand, Non-Patent Document 2 reports that an Fe ₃ Al—Al ₂ O ₃ composite material in which Al ₂ O ₃ is dispersed in an Fe ₃ Al alloy was obtained by electric pressure sintering. A dense body having a relative density of 96% or more is obtained by using electric current pressure sintering. However, the dispersion ratio of Al ₂ O ₃ is limited, and composite materials containing 20% by weight or more (about 30% or more by volume) of Al ₂ O ₃ are not targeted. That is, it is not possible to densify an Al ₂ O ₃ —Fe ₃ Al composite material having hard Al ₂ O ₃ particles as a parent phase by simply using electric pressure sintering. This is because the wettability of Al ₂ O ₃ and Fe ₃ Al is poor as well as the poor wettability of Al ₂ O ₃ and FeAl, so a small amount of Al ₂ is dispersed in the Fe ₃ Al matrix. This is because O ₃ does not have a great influence on densification, but in a system using Al ₂ O ₃ as a parent phase, it is a significant hindrance to densification.

From the above, using a FeAl alloy containing Al of 23 atomic% or more as the binder phase and sintering the hard phase Al ₂ O ₃ hard particles, the hard phase and binder phase are integrated and sintered. A dense composite material having excellent mechanical properties such as hardness and toughness and excellent durability against high-temperature oxidation has not been obtained.

JP 7-258786 A International Publication No. 1997/043228 Chinese Patent Application Publication No. 20121210097

S. SCHICKER et al., Reaction Synthesized A12O3-based Intermetallic Composites, Acta Materialia, 1996 46-7 p.2485-2492 Y. Bai et al., Study on preparation and mechanical properties of Fe3Al-20wt.% Al2O3 composites, Materials and Design, 2012, 39, p.211-219

The present invention has been made in view of the above circumstances, is excellent in high-temperature oxidation resistance of 900 ° C. or higher, and is dense and sintered with a binder phase and a hard phase integrated, such as hardness and toughness. The objective is to provide an Al ₂ O ₃ —FeAl composite material with excellent mechanical properties.

In order to solve the above problems, the oxide ceramic matrix composite material of the present invention comprises a hard phase having alumina (Al ₂ O ₃ ) hard particles as a parent phase and a binder phase containing an iron aluminum (FeAl) alloy. Have
The average particle size of alumina (Al ₂ O ₃ ) hard particles is 0.2 μm or more and 0.6 μm or less,
The content of the hard phase containing alumina (Al ₂ O ₃ ) is in the range of 50% to 75% by volume,
The iron aluminum (FeAl) alloy contains aluminum (Al) in a proportion of more than 23 atomic% and 50 atomic% or less,
The binder phase is characterized by being sintered integrally with the hard phase.

The method for producing an oxide ceramic matrix composite material of the present invention is a method for producing the oxide ceramic matrix composite material,
A process of preparing powder by mixing alumina (Al ₂ O ₃ ) powder, metal or alloy powder for forming an iron aluminum (FeAl) alloy, and additives;
And a step of applying current while pressing the powder at 10 MPa or higher in a vacuum of 100 Pa or lower and controlling the maximum temperature to be 1150 ° C. to 1350 ° C. and sintering.

According to the oxide ceramic matrix composite material of the present invention, it has excellent high-temperature oxidation resistance of 900 ° C. or higher, and has a dense structure in which a binder phase and a hard phase are integrally sintered, and has a machine such as hardness and toughness. Excellent mechanical characteristics.
According to the production method of the present invention, high temperature oxidation resistance of 900 ° C. or more is excellent, and the oxidation is excellent in the mechanical properties such as the hardness and the toughness, which are sintered by combining the binder phase and the hard phase. A ceramic composite material is obtained.

2 is a structure photograph (secondary electron image) and element map of an electric current pressure sintered body of Example 2. FIG. 6 is an external view photograph of vacuum sintered bodies of Comparative Examples 4, 5, and 6.

  The present invention is described in detail below.

The oxide ceramic matrix composite of the present invention has a hard phase having alumina (Al ₂ O ₃ ) hard particles as a parent phase and a binder phase containing an iron aluminum (FeAl) alloy.
In the oxide ceramics-based composite material of the present invention, the average particle diameter of Al ₂ O ₃ hard particles as a parent phase is 0.2 μm or more and 0.6 μm or less. In order for Al ₂ O ₃ hard particles to have a particle size of less than 0.2 μm, it is necessary to use a fine raw material powder of less than 0.2 μm. Grain growth occurs at a sintering temperature (between 1150 ° C. and 1350 ° C.) in a production method using sintering, and it becomes difficult to suppress the average particle size to less than 0.2 μm. On the other hand, when a relatively coarse raw material powder larger than 0.6 μm is used as the raw material powder, almost no grain growth of the Al ₂ O ₃ hard particles after sintering occurs. Since the densification is not promoted at the sintering temperature, the densification of the composite becomes insufficient. When the sintering temperature is 1350 ° C. or higher, the FeAl alloy bonded phase melts and precipitates on the surface, making it difficult to produce a sintered body in which the hard phase and the alloy bonded phase are integrated.

The average particle diameter of the Al ₂ O ₃ hard particles in the oxide ceramic matrix composite of the present invention is obtained by observing with an electron microscope or the like, and converting the hard particles present in the obtained image into an equivalent circular diameter. . Specifically, each measured particle is compared with a circle having a diameter of 0.1 μm from 0.1 μm to 2 μm and assigned to a circle having the closest diameter. Ignore fine particles less than 0.1 μm. In this way, the number N _i of particles corresponding to a circle having each diameter D _i is obtained. The average particle diameter D _ave of the hard particles in the composite material is obtained from the following equation (1) as an equivalent area diameter.
Here, A _i is the total area of particles having a diameter D _i .

In the oxide ceramic matrix composite of the present invention, the hard phase has Al ₂ O ₃ hard particles as a parent phase. However, ceramic carbide hard particles may be contained in a hard phase having Al ₂ O ₃ hard particles as a parent phase. The content of the hard phase is in the range of 50% to 75% by volume in the oxide ceramic matrix composite material. When the volume ratio is within this range, the bending strength, Vickers hardness, and fracture toughness of the oxide ceramic matrix composite material can be increased. When the volume ratio is less than 50%, it does not have hardness that contributes to a mold or a tool material. On the other hand, when the volume ratio exceeds 75%, the strength and fracture toughness are low, and the properties required for the mold and the tool material are insufficient.

Examples of the ceramic carbide hard particles constituting the hard phase include tungsten carbide (WC) and titanium carbide (TiC). These may be used alone or in combination of two or more. The oxide ceramic matrix composite of the present invention preferably contains ceramic carbide hard particles in a hard phase in a volume ratio smaller than that of alumina (Al ₂ O ₃ ) hard particles as a parent phase. Ceramic carbide hard particles such as WC and TiC are added in consideration of improvements in sinterability and hardness, but the volume fraction of carbide hard particles in the hard phase is higher than that of Al ₂ O ₃ hard particles. If the amount is too large, it will be difficult to carry out current-pressure sintering using the driving force for densification of Al ₂ O ₃ hard particles described later.

  In the oxide ceramic matrix composite material of the present invention, the FeAl alloy of the binder phase contains Al in a proportion of greater than 23 atomic% and not greater than 50 atomic%, preferably not less than 26 atomic% and not greater than 39 atomic%. . When the Al content is 23 atomic% or less, the high temperature oxidation resistance of the composite material produced decreases. On the other hand, when the proportion of Al exceeds 50 atomic%, the binder phase becomes brittle and the strength and toughness of the produced composite material are lowered.

  The oxide ceramic matrix composite of the present invention may contain boron (B) in the binder phase for the purpose of improving high temperature characteristics and hardness. When boron is used, the binder phase preferably contains 0.01% by weight or more and 0.4% by weight or less based on the total weight of the hard phase and the binder phase. When the boron content in the binder phase is too large, it becomes an obstacle to densification, and the mechanical properties and oxidation resistance of the composite material produced may be reduced.

The oxide ceramics-based composite material of the present invention is a process of preparing a powder by mixing Al ₂ O ₃ powder, metal or alloy powder for forming an FeAl alloy, and additives (hereinafter also referred to as a powder preparation process). ), Applying a current while pressing the powder at a pressure of 10 MPa or higher in a vacuum of 100 Pa or lower, and controlling the maximum temperature to be 1150 ° C. to 1350 ° C. (hereinafter also referred to as a sintering step). It can manufacture by the method containing.

By using energized pressure sintering in which the powder is energized at a pressure of 10 MPa or more in a vacuum of 100 Pa or less, the generation of pores is suppressed, and the maximum temperature of sintering is 1150 ° C to 1350 ° C. It is possible to obtain a dense oxide ceramic matrix composite by utilizing the driving force of both the densification of the hard matrix Al ₂ O ₃ and the FeAl alloy that is the binder phase. .

In the powder preparation process, the metal or alloy powder for forming the binder phase of the FeAl alloy includes, for example, metal powder such as Fe powder and Al powder, and alloy powder containing a metal compound such as FeAl ₂ and FeAl. Can be mentioned. Whichever of these powders is used, it reacts in the sintering process to produce FeAl.

  In the powder preparation step, examples of the additive include ceramic carbide hard particles such as WC and TiC described above, and boron (B).

  In the powder preparation step, mixing may be either dry or wet, but in the case of wet, it is preferable to use an organic solvent instead of water. When water is used, metal powders and alloy powders containing Al react with water, causing rapid oxidation and danger of explosion, etc., and using powders prepared using water, the strength and toughness of the sintered body This leads to a decrease in mechanical properties.

  In the sintering step, current-pressure compression sintering is used in which a current is applied while pressing the powder at 10 MPa or higher in a vacuum of 100 Pa or lower. Thereby, formation of pores can be significantly suppressed. If current-pressure sintering is not used, a composite material in which the FeAl alloy phase is precipitated on the surface is obtained due to poor wettability, the relative density is low, and the strength is low. Generally, a graphite mold is used as the mold for the electric current pressure sintering, but other materials may be used. On the other hand, if the degree of vacuum exceeds 100 Pa, the oxidation of the alloy powder is likely to proceed, which causes the mechanical properties of the sintered body to deteriorate. Regarding the pressurization at the time of sintering, the upper limit of the pressure is not particularly limited, but a range in which the mold for current-pressure compression sintering does not break is preferable.

In producing the oxide ceramic composite material of the present invention, it is basically important to suppress the liquid phase of the binder phase during sintering. When a large amount of liquid phase is formed by pressure sintering, the liquid phase overflows from the sintering mold and cannot be controlled to the target composition. For this reason, it is necessary to sinter at a temperature lower than the temperature at which the FeAl alloy is judged to have formed a liquid phase. However, a temperature at which FeAl itself becomes dense by itself is necessary. On the other hand, Al ₂ O ₃ also needs to have a sufficiently dense driving force at the sintering temperature. Considering these points, it is preferable to sinter after designing the raw materials so that the densification temperature of Al ₂ O ₃ is lower than the liquid phase formation temperature of FeAl and the difference is within 30 ° C. The liquid phase formation temperature of FeAl in electric pressure sintering is affected by various conditions such as the composition of FeAl, the size of the graphite mold, and the size of the pressure, but it exists in the range of 1150 ° C to 1350 ° C. In order to make the Al ₂ O ₃ densification temperature lower than the liquid phase formation temperature of FeAl and to keep the difference within 30 ° C, it is necessary to use Al ₂ O ₃ particles that are fine enough to be densified at low temperature There is. Specifically, it is necessary to prepare the mixed powder so that the average particle diameter of Al ₂ O ₃ after sintering is 0.2 μm or more and 0.6 μm or less. By using the sintering process as described above, a dense oxide ceramics-based composite material can be produced by making use of the driving forces of both Al ₂ O ₃ densification and FeAl densification.

The oxide ceramic matrix composite obtained by the above production method has a structure in which a binder phase containing a FeAl alloy of 5 μm or less is finely dispersed with respect to a hard phase having Al ₂ O ₃ as a parent phase, that is, a hard phase and It has a structure in which the binder phase is integrated and sintered. Since the finely dispersed binder phase has sufficient toughness and high-temperature oxidation resistance, the oxide ceramics-based composite material of the present invention has a high level of hardness and toughness and is excellent in high-temperature oxidation resistance. Become.

  In the oxide ceramic matrix composite material of the present invention, the Vickers hardness is preferably in the range of 10 GPa to 20 GPa, more preferably 12 GPa to 18 GPa. In the present invention, the Vickers hardness is determined as Hv by a measurement method defined in JIS Z2244 at room temperature (15 to 30 ° C.).

In the oxide ceramic matrix composite of the present invention, the bending strength is preferably in the range of 0.9 GPa to 2.0 GPa, and more preferably 1.2 GPa to 2.0 GPa. The bending strength is obtained by conducting a three-point bending test on a rectangular parallelepiped specimen. The room temperature bending strength σ (GPa) is calculated from the following formula (2) from the obtained fracture load. Here, F is the breaking load (kN), L is the span length (mm), B is the specimen width (mm), and H is the specimen thickness (mm).

Terms similar to the above, the oxide ceramic matrix composite of the present invention, fracture toughness is preferably in the range of 4.6MPam ^0.5 or more 13MPam ^0.5 or less, more preferably 7.0 MPAM ^0.5 or more 11MPam ^0.5 or less. In the present invention, fracture toughness (K _IC ) is determined as K _IC by an indentation press-in method (IF) method estimated from the length of a crack extending from a Vickers indentation specified in JIS R1760 at room temperature (15 to 30 ° C.). .

  The oxide ceramic matrix composite of the present invention is characterized by having both the bending strength, Vickers hardness and fracture toughness in the above ranges. This is because the binder phase and the hard phase are sintered together.

The oxide ceramic matrix composite material of the present invention preferably has a weight increment of 5.0 × 10 ⁻⁴ g / cm ² or less after being kept in the atmosphere at 900 ° C. for 24 hours, and 1.0 × 10 ⁻⁴ g / cm 2. Particularly preferred is ² or less. When the weight increment is 1.0 × 10 ⁻⁴ g / cm ² or less, it becomes a significant difference compared to the high temperature oxidation resistance of commercially available cermets, and is highly useful.

  EXAMPLES Hereinafter, the present invention will be described in more detail with reference to examples, but the present invention is not limited to these examples.

Commercially available α-Al ₂ O ₃ powder with an average particle size of 0.3 μm (manufactured by Kojundo Chemical Laboratory Co., Ltd.), α-Al ₂ O ₃ powder with an average particle size of 1.4 μm (manufactured by Showa Denko KK), average WC powder with a particle size of 0.7μm (manufactured by Nippon Shin Metal Co., Ltd.), FeAl ₂ powder with a particle size of 300μm or less (Fe: 33mol%, Al: 67mol%), Fe powder with a particle size of 3-5μm Each raw material powder was mixed so as to have the composition shown in Table 1, using B powder (manufactured by Kojundo Chemical Laboratory Co., Ltd.) and B powder (manufactured by Kojundo Chemical Laboratory Co., Ltd.) having an average particle size of 45 μm as raw materials. Specifically, all the raw material powder was put into a 420 ml stainless steel pot, 1200 g of 5 mm diameter carbide balls and 100 ml of acetone were added, and wet ball mill mixing was performed, but only Comparative Example 1 had an average particle size of 1.4 μm. using α-Al ₂ O ₃ powder, other for using α-Al ₂ O ₃ powder having an average particle diameter of 0.3 [mu] m. speed in uniform 100 rpm, milling time was 48 hours or 96 hours. the resulting The obtained mixed powder was dried using an evaporator and a vacuum dryer.

  The mixed powder after drying was sintered by the following two methods. The first is a method in which the mixed powder is formed by uniaxial pressing and then sintered under a vacuum of 10 Pa or less (hereinafter referred to as vacuum sintering). The second is a method in which the mixed powder is placed in a graphite mold and uniaxially pressed. This is a method of performing energization pulse sintering (hereinafter referred to as energization and pressure sintering). In vacuum sintering, the molding pressure was 100 MPa and the sintering temperature was uniformly 1450 ° C. On the other hand, in electric pressure sintering, the mixed powder is put into a graphite mold with an inner diameter of 10 mm or an inner diameter of 23 mm, and an electric pulse sintering apparatus is used. Sintering was performed in the range of ˜1230 ° C. As a preliminary step to determine the sintering temperature, in order to estimate the liquidus temperature of the FeAl alloy contained in the mixed powder, the heating rate was fixed at 10 ° C / min in the temperature range above 1100 ° C, The amount of contraction was confirmed with a displacement meter. The shrinkage rate suddenly increases from a certain temperature of 900 ° C. or more and then gradually decreases, but then the shrinkage rate apparently increases again. The increase in the shrinkage rate again occurs because the liquid phase of the produced FeAl alloy oozes out of the mold. For this reason, the temperature at which the shrinkage rate increases again at 1100 ° C. or higher was regarded as the liquid phase formation temperature. The oxide ceramic matrix composite of the present invention is sintered at a temperature lower than that with reference to the estimated liquid phase formation temperature. In addition, the temperature described here has shown the temperature measured by matching the hole drilled in the graphite type | mold with the radiation thermometer.

After confirming the presence or absence of surface precipitation of the alloy phase, the sintered sample was cut and surface ground, processed into a test piece having a width of 4 mm and a thickness of 2 mm, and density measurement was performed by Archimedes method. Table 2 shows each sample sintering method, the estimated liquid phase formation temperature, the sintering temperature theoretical density, the apparent density obtained from the Archimedes method, and the relative density obtained by dividing the apparent density by the theoretical density. Here, the theoretical density represents a value calculated from the ratio of each compound and the individual density when it is assumed that the hard phase such as Al ₂ O ₃ , WC, TiC and the FeAl alloy bonded phase do not react.

  Evaluation of sintered sample includes measurement of average particle size of hard particles, analysis of Al and Fe ratio in FeAl alloy binder phase, evaluation of integration of hard phase and binder phase, bending strength, Vickers hardness, Young The rate and fracture toughness were measured. A summary of these results is shown in Table 3. For comparison, some characteristics of high toughness TiCN-Ni cermets are also shown. The method for obtaining each characteristic value shown in Table 3 will be described below.

  The measurement of the average particle diameter of the hard particles and the analysis of the ratio of Al and Fe in the FeAl alloy binder phase were performed as follows. First, using the polished sample, we take microstructure photographs by SEM at various magnifications and focus on the part corresponding to the binder phase by elemental analysis by energy dispersive X-ray (EDX). The element ratio was determined. Specifically, the elements present in the binder phase were first estimated by qualitative analysis, and then the ratio of each element was determined using simple quantification for the corresponding elements. Moreover, the average particle diameter of the hard particles was obtained by calculating the area equivalent diameter constituting each sample based on the formula (1) from the obtained tissue image. In determining the average particle size of the hard particles, the polished sample was thermally etched at 1200 ° C. and then subjected to SEM observation in order to clearly observe the grain boundaries of the hard particles. FIG. 1 shows the secondary electron image of the structure of the oxide ceramic matrix composite material of the present invention and the distribution of the corresponding elements (O, Fe, Al) (Example 2, no etching).

  The evaluation of the integration of the hard phase and the binder phase was performed by visually evaluating the sample. Specifically, as shown in FIG. 2, when it can be seen that the FeAl alloy is spherically deposited on the surface of the sintered sample, it is considered that the hard phase and the binder phase have not been successfully integrated. .

  The bending stress was obtained by processing a sample into a rectangular parallelepiped test piece having a width of 4 mm and a thickness of 2 mm, and performing a three-point bending test on the test piece. The displacement speed was 0.5 mm and the span was 10 mm. The room temperature bending strength σ (GPa) was determined from the above-described equation (2) from the obtained fracture load.

The Vickers hardness was determined based on JIS Z2244 from the size of the indentation after polishing a part of the surface of the broken piece after the bending test and driving in a diamond indenter. Fracture toughness K _IC was determined based on JIS R1760 from cracks generated from indentations. Since the Young's modulus of the sample is required for calculation of fracture toughness, the longitudinal wave velocity and the transverse wave velocity of each sample are measured with an ultrasonic elastic modulus measuring instrument, and the Young's modulus E (GPa) is calculated from the following equation (3). Was calculated.
Here, ρ is the sample density (g / cm ³ ), and V _L and V _S are the longitudinal wave velocity and the transverse wave velocity (m / s), respectively.

Comparing the corresponding element distribution with the structure photograph (secondary electron image) of Example 2 in FIG. 1, it can be seen that the white part corresponds to the FeAl alloy phase and the black part which is the parent phase corresponds to the Al ₂ O ₃ hard phase. . In addition, the binder phase is not a structure that spreads between the hard particles, but a structure in which the binder phase having a particle shape of 5 μm or less is uniformly dispersed in the hard phase particles and the binder phase is integrally sintered. I know that there is. All the complexes of the present invention have such characteristics. The binder phase does not spread between the hard particles because the wettability of the FeAl alloy and Al ₂ O ₃ is poor, but the binder phase with a particle shape of 5 μm or less is finely dispersed in the hard phase, that is, hard It turns out that the composite material which covered the bad wettability and improved the intensity | strength and toughness was produced because the phase and the binder phase were sintered integrally.

On the other hand, in Comparative Examples 4, 5, and 6 obtained by vacuum sintering, as shown in FIG. 2, FeAl alloy is seen to be deposited in a spherical shape on the surface of the sintered sample, and bonded to the hard phase. It can be seen that the phases are not integrated and sintered. Further, the relative density is lower than 92% except for Comparative Example 6 (see Table 2), which is an incomplete dense body. The size of the precipitate tends to increase as the addition ratio of the raw material powder for constituting the FeAl alloy binder phase increases. From this, it is considered that in the sintering method using vacuum sintering, the liquid phase FeAl has poor wettability with the Al ₂ O ₃ hard phase, so that densification is not promoted and is deposited on the surface.

Comparing the results of Table 1 and Table 3 regarding the composition of the FeAl binder phase, it can be seen that the proportion of Fe in the binder phase of the actual sample is higher than the theoretical FeAl composition. This is because the specific surface area increases and surface oxidation occurs because the particle diameter of the coarse FeAl ₂ powder, which is a coarse material, becomes sharply fine during pulverization. Surface oxidation causes Al ₂ O ₃ to form from the FeAl alloy bonded phase during sintering, and the Al component is deprived, so the proportion of Fe in FeAl becomes relatively large. In the present invention, the proportion of Al in the FeAl alloy bonded phase is specified to be more than 23 atomic% and not more than 50 atomic%, but this can be controlled by the ratio of raw material powder and preparation conditions.

  The average particle size of the hard particles in the sintered body largely depends on the average particle size of the hard particles as a raw material from the results shown in Table 3, and there is almost no influence of the wet mill time or the composition of the FeAl binder phase at this time. I understood it. This is because there was no significant difference in the pulverization effect of the hard particles due to changes in the wet mill time and the ratio of the alloy binder phase, and the hard particles refined during sintering caused grain growth by being incorporated into other large particles. Therefore, it is considered that the difference in particle size is reduced. Regarding Comparative Example 1 in which the raw material powder was larger than the other samples, the average particle size of the hard particles after sintering was clearly larger than that of the other samples.

  Mechanical properties such as Young's modulus, bending strength, hardness, fracture toughness, etc. in Examples 1 to 9 and Comparative Example 2 produced by electric pressure sintering are generally not as good as those of high toughness cermets, but have characteristics close to them. . These mechanical properties depend on the composition of the binder phase of the FeAl alloy and the mill time of the raw material powder. In fact, the higher the proportion of the binder phase of the FeAl alloy, the higher the hardness and Young's modulus and the lower the fracture toughness. Moreover, the longer the mill time of the raw material powder, the higher the hardness and Young's modulus and the lower the fracture toughness.

On the other hand, Comparative Examples 1, 3, and 4 have a relative density of less than 90% and incomplete densification, so that both bending strength and hardness are lower than in all Examples. In Comparative Examples 5 and 6, the relative density exceeded 91%, but the strength was lower than in all Examples. As shown in FIG. 2, since Comparative Examples 4, 5, and 6 use vacuum sintering, the fact that the hard phase and the binder phase are not integrated and sintered is also a factor of low mechanical properties. Conceivable. In Comparative Examples 1, 4, and 5, the progress of cracks could not be clearly confirmed, and the fracture toughness could not be measured. In Comparative Example 3, the progress of cracks was confirmed, but the value was as low as 2.9 MPa ^0.5 .

A high-temperature oxidation resistance test was conducted at 900 ° C. in the atmosphere using each sample processed into a bending test piece. The sample was placed in an electric furnace heated to 900 ° C. and heated for 24 hours. After cooling to room temperature, the sample was taken out and the weight increment was measured. Table 4 shows the amount of increase in oxidation relative to the surface area (unit: g / cm ² ).

  It can be seen that Examples 1 to 9 have a smaller oxidation increase of 10 to 100 times and superior oxidation resistance compared to the high toughness TiCN-Ni cermet. On the other hand, in Comparative Examples 1, 3, and 4, densification was not achieved (90% or less), and therefore, the area exposed to oxygen, that is, the specific surface area was larger than that of the Example, so that the increase in oxidation was 10 About 100 times larger. In Comparative Example 2, the density is sufficient (95.4%), but the ratio of Al contained in the binder phase is as low as 7 atomic%, and an oxide film of Al is not formed. About 100 times larger, almost the same as cermet. For this reason, the composite material of the present invention has excellent high-temperature oxidation resistance of 900 ° C. or higher as long as it is sufficiently densified and the proportion of Al in the binder phase satisfies a specific range (greater than 23 atomic%). It turns out that it becomes a material.

  The oxide ceramics-based composite material of the present invention is superior in durability to the atmosphere at 900 ° C. or higher than cermet known as a material excellent in high-temperature oxidation resistance, and has mechanical characteristics comparable to cermet 900. It can be suitably used as a material for a cutting tool or a peripheral component for casting that is exposed to a high temperature environment of ℃ or higher.

Claims (7)

A hard phase having alumina (Al ₂ O ₃ ) hard particles as a matrix and a binder phase containing an iron aluminum (FeAl) alloy;
The average particle size of alumina (Al ₂ O ₃ ) hard particles is 0.2 μm or more and 0.6 μm or less,
The content of the hard phase containing alumina (Al ₂ O ₃ ) is in the range of 50% to 75% by volume,
The iron aluminum (FeAl) alloy contains aluminum (Al) in a proportion of more than 23 atomic% and 50 atomic% or less,
An oxide-ceramic matrix composite material in which the binder phase is sintered together with the hard phase.   The oxide ceramic matrix composite material according to claim 1, wherein the binder phase is not precipitated on the surface. _3. The oxide ceramic matrix composite material according to claim 1, wherein the hard phase contains ceramic carbide hard particles in a volume ratio smaller than that of alumina (Al ₂ O ₃ ) hard particles as a parent phase.   The oxide ceramic matrix composite material according to any one of claims 1 to 3, wherein the binder phase contains boron in a range of 0.01 wt% to 0.4 wt% with respect to the total weight of the hard phase and the binder phase. . The oxide ceramic base according to any one of claims 1 to 4, wherein the bending strength is 0.9 GPa or more and 2.0 GPa or less, the Vickers hardness is 10 GPa or more and 20 GPa or less, and the fracture toughness is 4.6 MPaam ^0.5 or more and 13 MPaam ^0.5 or less. Composite material. The oxide ceramic matrix composite material according to any one of claims 1 to 5, wherein a weight increment after being kept in an atmosphere of 900 ° C for 24 hours is 1.0 × 10 ^-4 g / cm ² or less. A method for producing an oxide ceramic matrix composite material according to any one of claims 1 to 6,
A process of preparing powder by mixing alumina (Al ₂ O ₃ ) powder, metal or alloy powder for forming an iron aluminum (FeAl) alloy, and additives;
A method for producing an oxide ceramic matrix composite material, comprising: energizing the powder while pressing it at a pressure of 10 MPa or higher in a vacuum of 100 Pa or lower, and controlling and sintering so that the maximum temperature is 1150 ° C. to 1350 ° C. .