KR102320583B1 - Cermet having improved toughness and method for manufacturing the same - Google Patents

Abstract

An object of the present invention is to provide a cermet having superior toughness by overcoming the low toughness of the existing simulated cermet and the low hardness of the cermet without the simulated structure. According to one aspect of the present invention, particles having a core structure consisting of a core region and a rim region in a complete solid solution phase of carbides of two or more metals selected from Group IVa, Va, and VIa metals in the periodic table, including Ti; and a binder made of a metal; is provided, including a cermet.

Description

Cermet having improved toughness and method for manufacturing the same}

The present invention relates to a cermet having excellent toughness used for cutting or mining, and a method for manufacturing the same.

WC-based cemented carbide, TiC or Ti(CN)-based various cermet alloys, other ceramics, CBN or high-speed tool steel are used as the main cutting tools or wear-resistant tools used for cutting and mining.

Among them, cermet is generally made of hard TiC or Ti(CN) and metals such as Ni, Co and Fe as binders, and carbides, nitrides and carbonitrides of Group IVa, Va and VIa metals on the periodic table. It refers to a ceramic matrix-like composite sintered body containing as an additive. _{The production of the cermet is performed by mixing hard additive carbides such as WC, NbC, TaC, Mo 2} C, etc., in addition to TiC or Ti(CN), and metal powders such as Co, Ni and Fe, which are matrix phases for bonding them, and these It is manufactured by sintering in a vacuum or gas atmosphere. The first mass production of such a cermet is a TiC-Mo ₂ C-Ni cermet manufactured by Ford Motor in the United States in 1956. Used for semi-finishing and finishing. In the 1960s and 1970s, attempts were made to add various kinds of elements in order to improve toughness, which is a major weakness of the TiC-Ni cermet system, but no clear results were achieved.

Meanwhile, Ti(C,N), a more thermothermally stable phase, was synthesized by adding TiN to TiC in the 1970s, and as a result, there was some improvement in toughness. Since Ti(C,N) has a finer structure than TiC, toughness could be improved, and chemical stability and mechanical impact resistance could also be improved. On the other hand, _{many added carbides such as WC, MO 2} C, TaC, and NbC have been used to improve toughness, and various products have been commercialized until now. In this way, when carbide is added to improve toughness, the general microstructure of the TiC-based or Ti(CN)-based cermet sintered body is observed as a core structure. surrounded by a binder.

In the core structure, the core region is TiC or Ti(CN) that is not dissolved in a liquefied metal binder (Ni, Co, Fe, etc.) during sintering, and has a high hardness. On the other hand, the rim region surrounding these cores contains TiC or Ti(CN), which is a component of the core partially dissolved in a liquefied metal binder (Ni, Co, Fe, etc.), and It is a solid-solution between addition carbides (represented in the form of (Ti,Me1,Me2,...)(C,N), etc.), which is a hard phase with higher toughness than the core. Since the rim area has good wettability with metal binders (Ni, Co, Fe, etc.), it exhibits effects such as sintering well and improving toughness. In this way, the cermet was able to increase the toughness to some extent by solving the problem of low wettability, which is a fatal weakness of simple cermets such as TiC-Ni or Ti(CN)-Ni, through the formation of the rim structure.

However, in the case of the cermet having the above-mentioned core structure, the ^{toughness is still quite low at the level of 6-8 MPa·m 1/2} compared to cemented carbide such as WC-Co, and it still cannot completely replace WC-Co. The reason has been found to be due to the fact that a very high strain energy is accumulated at the core/rim interface of the SIM card structure, and thus it is easily broken in an environment to which a mechanical impact is applied.

Accordingly, an attempt to develop a cermet with improved toughness through the formation of a complete solid solution phase without a simile structure has been steadily made by tool companies such as Sumitomo and various researchers. In the case of a fully solid solution cermet without a core structure, toughness equivalent to cemented carbide such as WC-Co can be secured at the level of ^{11 to 14 MPa·m 1/2.} couldn't get

The present invention is to solve such a problem in the prior art, and overcomes the low toughness of the existing simulated cermet and the low hardness of the cermet without the simulated structure, and provides a cermet with much higher toughness than these. The purpose. However, these problems are exemplary, and the scope of the present invention is not limited thereto.

According to one aspect of the present invention, a particle having a core structure consisting of a core region and a rim region, which is a complete solid solution phase of carbides of two or more metals selected from Group IVa, Group Va, and Group VIa metals in the periodic table, including Ti, and A cermet including a binder made of metal is provided.

According to an embodiment of the present invention, the composition of the group VIa metal in the core region may have a lower value than the composition in the rim region.

According to an embodiment of the present invention, the lattice constant in the rim region may have a larger value than the lattice constant in the core region.

According to an embodiment of the present invention, the rim region may have a compressive stress state compared to the core region.

According to an embodiment of the present invention, the composition of the group VIa element at the interface between the core region and the rim region may gradually increase from the core region to the rim region. According to an embodiment of the present invention, the core region and The difference in composition of the group VIa in the rim region may range from 1 at% to 10 at%.

According to an embodiment of the present invention, the group VIa element may include any one or more of W and Mo.

According to an embodiment of the present invention, the complete solid solution phase of the carbide of the metal may have a NaCl-type face-centered cubic structure as a solid solution of TiC and WC.

According to an embodiment of the present invention, the binder may include any one or more of Ni, Co, and Fe.

According to another aspect of the present invention, a complete solid solution phase carbide powder in which two or more metal carbides selected including Ti from group IVa and VIa metals in the periodic table are completely solid-dissolved; an additive comprising at least one metal carbide powder selected from Va and Group VIa metals; and one or more metal powders; There is provided a method for producing a cermet by sintering a mixed powder comprising a.

According to an embodiment of the present invention, the step of sintering may include: completely dissolving the first complete solid solution phase carbide powder in the binder; and precipitating a second perfect solid solution phase carbide having a different composition from the first perfect solid solution phase carbide from the binder.

According to an embodiment of the present invention, the step of precipitating the second complete solid-solution carbide may include precipitating a core region; and depositing a rim region having a relatively higher composition of the Group VIa metal than the core region around the core region after the core region is deposited.

According to an embodiment of the present invention, the perfect solid solution powder may have an average size of 60 nm or less.

According to an embodiment of the present invention, it may have a NaCl-type face-centered cubic structure as a solid solution of TiC and WC.

According to an embodiment of the present invention, the additive may _{include any one or more of WC, Mo 2} C, TaC, NbC and VC.

According to an embodiment of the present invention, the metal powder may include any one of Ni powder, Co powder, and Fe powder.

The cermet according to an embodiment of the present invention made as described above has a microstructure differentiated from that of the prior art, thereby achieving the effect of remarkably improving the mechanical properties of the product as it has very excellent toughness. Of course, the scope of the present invention is not limited by these effects.

1 is a result of observation with a scanning electron microscope (SEM) of _{(Ti 0.88} W _0.12 ) C powder used for manufacturing a cermet according to an experimental example of the present invention.
2 is an X-ray diffraction analysis result of the (Ti _0.88 W _{0.12 )C powder.}
3 is an X-ray diffraction analysis result of a cermet according to an experimental example of the present invention.
4 is a result of observing the microstructure of a cermet according to an experimental example of the present invention with a scanning electron microscope (SEM).
5 is a result of observing the microstructure of a cermet according to an experimental example of the present invention with a transmission electron microscope (TEM).
6 is a result of analyzing the W content in the core region and the rim region of the cermet according to the experimental example of the present invention.
7 is a line scanning result for analyzing the composition distribution of cermet carbide particles according to an experimental example of the present invention.
8 is a result of observation with a transmission electron microscope of cermet carbide particles according to an experimental example of the present invention.
9 is a result of analyzing Vickers hardness and fracture toughness of a cermet according to an experimental example of the present invention.
10 is a result of observing the crack shape of the cermet according to the experimental example of the present invention.
11 and 12 are results of observation of cermet carbide particles according to an experimental example of the present invention with a transmission electron microscope (TEM) and a scanning electron microscope (SEM), respectively.
13 shows the Vickers hardness (Hv30) and toughness results of the cermet according to the experimental example of the present invention, along with the results of commercial WC-Co cemented carbide materials, commercial cermets, and cermets reported in the literature.

Hereinafter, embodiments of the present invention will be described in detail with reference to the accompanying drawings. However, the present invention is not limited to the embodiments disclosed below, but can be implemented in various different forms. It is provided to fully inform In addition, in the drawings for convenience of description, the size of the components may be exaggerated or reduced.

A cermet according to an embodiment of the present invention includes carbide particles forming a complete solid solution phase of a hard phase and a metal binder for bonding them.

The metal carbide constituting the perfect solid solution phase may form a complete solid solution phase with at least one of a carbide of a Group IVa metal essentially containing Ti, and a carbide of a Group Va and Group VIa metal. For example, the carbide particles may form a complete solid solution phase with TiC and at least one of group Va and group IVa metal carbides. The Group Va metal carbide may include VC, TaC, and NbC, and the Group VIa metal carbide may include WC, Mo ₂ C, and the like.

For example, the perfect solid solution carbide particles may have a structure in which a part of Ti atoms are substituted with other metal atoms in a TiC lattice having a NaCl-type face-centered cubic structure. In this case, the substituted metal atom may be any one or more of group IVa, Va, and VIa metal atoms. The dissolved metal atoms are added up to the solid solution limit of TiC. When the solid solution limit is exceeded, a second carbide is precipitated in addition to the complete solid solution phase carbide. For example, in the complete solid solution phase (Ti, W)C, when the content of W exceeds the solid solution limit, WC may be precipitated as a second carbide.

According to an embodiment of the present invention, the perfect solid solution carbide particles have a core structure having a core region and a rim region. The core region and the rim region are made of the same element and have the same crystal structure, but there is a composition difference between the elements of the same kind, and due to the composition difference, the core region and the rim region may be divided. For example, when the carbide particles are (Ti, W)C with a NaCl-type face-centered cubic structure in which TiC and WC are completely solid phases, the W content in the core region is smaller than the W content in the rim region. Therefore, the Ti content in the core region may have a relatively larger value than the Ti content in the rim region. However, the carbide particles have a simulating structure, but a change in composition occurs gradually and continuously at the boundary between the core region and the rim region. Therefore, the boundary between the core area and the rim area may not be clear and may appear unclear. The difference in composition of the group VIa, for example, W, in the core region and the rim region may have a value of 1 at% to 10 at%.

Although the perfect solid solution carbide particles have the same lattice structure, local differences in stress state exist due to the characteristics of local composition. In the metal carbide particles according to the embodiment of the present invention, the lattice constant in the rim region exhibits a larger value than the lattice constant in the core region. For example, when the carbide particle is (Ti, W)C composed of a perfectly solid solution of TiC-WC, the lattice constant of the TiC lattice is will change When the particle to be substituted has a larger atomic radius than Ti, the lattice constant will increase, and as the content of the substituted metal atom increases, the increase in the lattice constant will be larger. For example, when the carbide particles are (Ti, W)C in which TiC and WC are completely solid phases, the atomic radius of W has a larger value than that of Ti, and therefore the rim region with a high W content. It shows a larger lattice constant compared to the core region, and also shows a compressive stress.

The cermet of the present invention exhibits superior toughness properties compared to the prior art due to the carbide particles having the core/rim structure of the core/rim structure in which the composition change occurs gradually and the carbide particles having the rim region under compressive stress.

In the conventional cermet structure, the composition changes rapidly at the interface between the core region and the rim region. It is known to break easily.

In contrast, in the cermet according to the embodiment of the present invention, since the composition change at the interface of the core region and the rim region is continuous and gradual, it is possible to solve the problem of deterioration of toughness due to stress caused by a sudden change in composition.

On the other hand, it is known conventionally that Ti(CN) powder is used as a complete solid solution cermet of a core/rim structure. Compared with the conventional perfect solid solution phase cermet, the perfect solid solution phase cermet of the present invention is different in that it is a perfect solid solution phase between different carbides. That is, the present invention does not contain N, unlike the conventionally known Ti(CN) perfect solid solution phase, and due to this, the content of a metal component that does not have a high affinity with N, for example, W, is higher in the case of the present invention. . In addition, since the perfect solid solution carbide of the perfect solid solution phase cermet of the present invention is a substitution type solid solution phase in which one metal atom is substituted for another metal atom, the lattice distortion in the solid solution phase is higher than that of conventional Ti(CN). show the difference

The cermet having such a microstructure is manufactured by sintering a mixed powder obtained by mixing a metal carbide powder and a metal carbide additive powder forming a complete solid phase with a metal powder constituting a binder. As a method of manufacturing the sintered body, high-temperature sintering, pressure sintering, energizing pressure sintering, etc. may be used, and a series of processes ranging from powder mixing, molding, and sintering for manufacturing cemet are well-known techniques in the field, so detailed description will be omitted. .

In the present invention, it is necessary to adjust the size of the perfect solid solution carbide particles used as a raw material to be smaller than that of the prior art. Hereinafter, these reasons will be described.

The amount in which the carbide is dissolved may be expressed by Equation 1 below.

:용해된 탄화물 i의 평균 농도( mol/cm³,)

: mean concentration of dissolved carbide i ( mol/cm ³ ,)

: 탄화물 i의 중량당 평균 표면적(cm²/g),

: average surface area per weight of carbide i (cm ² /g),

: 탄화물 i 입자의 농도 (g/cm³)

: Concentration of carbide i particles (g/cm ³ )

: 탄화물 i의 평균 용해 속도 (mol·(cm²·s)^-1)

: Average dissolution rate of carbide i (mol·(cm ² ·s) ^-1 )

: 용해 시간(초)

: Dissolution time (sec)

값)을 향상시키기 위해서는 사용하는 분말의 표면적을 가능한 작게 하여야 한다. 즉, 수학식 1에서,

는 소결 조건 및 조성이 결정되면 고정되는 값이다. 따라서,

값을 조정함으로써, 동일한 시간 동안

값을 크게 할 수 있다. The amount of dissolved carbide constituting the sintered body during grain is basically governed by the thermodynamic stability of the carbide. In general, in the case of carbides having a solid solution phase, for example, (Ti, W)C, which is a solid solution phase of TiC and WC, it has thermodynamically higher stability than a single-phase carbide other than a solid solution phase, for example TiC and WC, and, thus, dissolves It's slow. Therefore, when referring to Equation 1, the dissolved amount (

value), the surface area of the powder used should be as small as possible. That is, in Equation 1,

is a value that is fixed when the sintering conditions and composition are determined. thus,

By adjusting the value, for the same time

value can be increased.

For this reason, when the raw material carbide used for manufacturing the sintered body is a solid solution phase carbide, the surface area is increased by reducing the size of the raw material carbide compared to when a normal carbide is used instead of a solid solution phase to produce a sintered body faster and more uniformly. be able to According to an embodiment of the present invention, the average particle size (particle size) of the carbide having a solid solution phase may be 60 nm or less (greater than 0), preferably 5 nm to 50 nm, more preferably 10 nm to 40 nm.

The metal carbide additive comprises at least one metal carbide powder selected from Va and Group VIa metals, including W. For example, it may include WC, Mo ₂ C, TaC, NbC, VC, and the like. The metal carbide additive is a supply agent for supplying the same metal as at least one of the metal components included in the perfect solid solution phase carbide into the inside of the perfect solid solution phase carbide powder particles in the process of being mixed with the perfect solid solution carbide powder and sintered can perform the role of Due to the addition of such metal carbide, mechanical properties of the cermet, for example, hardness or toughness may be improved. The metal carbide additive may have a range of 1 to 5 μm, preferably 2 to 4 μm.

The binder is made of one or more metals that are melted during sintering and serve to bind unmelted carbide particles. The binder may be selected from, for example, metals such as Ni, Fe, and Co.

Hereinafter, in order to help the understanding of the present invention, an experimental example to which the above-described technical idea is applied will be described. However, the following experimental examples are only for helping understanding of the present invention, and the present invention is not limited by the following experimental examples.

[Experimental example]

Table 1 shows the type and composition of the powder used as a raw material for manufacturing the cermet. The composition is expressed in wt% (wt%).

Experimental example (Ti _0.88 W _0.12 )C WC Mo ₂ C NbC TaC Ni Co One 70 10 0 0 0 20 0 2 60 20 0 0 0 20 0 3 50 30 0 0 0 20 0 4 40 40 0 0 0 20 0 5 40 40 0 0 0 10 10 6 55 30 0 0 0 15 0 7 40 30 5 3 2 20 0

_{The (Ti 0.88} W _0.12 )C powder, which is the complete solid solution powder used in the preparation of the Experimental Example _{, was quantified according to the mole fraction of TiO 2} , WO ₃ , and C (graphite) and was subjected to high-energy milling, followed by carbonization at 1150° C. for 4 hours. was prepared. 1 is a result of observing the prepared (Ti _0.88 W _0.12 )C powder with a scanning electron microscope (SEM), and FIG. 2 is a result of X-ray diffraction analysis. Table 2 shows the BET results showing the particle size and specific surface area of the powder.

powder Particle size (㎛) BET surface area (m ² /g) (Ti _0.88 W _0.12 )C 0.03 ^a /0.029 ^b 35.685 WC 2.73 - Ni 2.40 - ^a Particle size by BET measurement
^b Crystallite size calculated by the Rietveld refinement method

1 and Table 2, the (Ti _0.88 W _0.12 )C powder was approximately spherical and had a powder size corresponding to about 30 nm, and the BET value was about 35.6 m ² /g. Meanwhile, in Table 2, the particle sizes of the WC and Ni powders used in the preparation of the Experimental Example are also shown.

Referring to the X-ray diffraction analysis result of FIG. 2 _{, it can be seen that the (Ti 1-x} W _x )C powder consists of a single phase showing a TiC crystal structure, which is a NaCl-type face-centered cubic structure. That is, (Ti _0.88 W _0.12 )C is interpreted as having a lattice of a TiC crystal structure, but having a solid solution crystal structure in which Ti atoms in the TiC lattice are substituted with W. From this, it can be confirmed that the (Ti _0.88 W _0.12 )C powder has a single phase in which TiC and WC are completely dissolved.

The powders shown in Table 1 were mixed using a ball-mill and then uniaxially pressed at 150 MPa in a cylindrical die to make a compact, and the compact was put into a vacuum furnace and sintered at 1510° C. for 1 hour to prepare a cermet.

3 shows the results of X-ray diffraction analysis on the sintered body corresponding to Experimental Example 4 after sintering is completed.

Referring to the X-ray diffraction analysis result of FIG. 3 , it can be confirmed that only the (Ti,W)C diffraction peak, which is the fully dissolved phase of TiC and WC, and the diffraction peak corresponding to the nickel binder are found. From this, it can be confirmed that the carbide of the simulating structure constituting the sintered body is a single-phase complete solid solution phase. In addition, the lattice constant obtained from Bragg's law using the (220) peak among the X-ray diffraction peaks showed a value of 4.3087 Å, which was very similar to the value derived from electron diffraction of a transmission electron microscope, which will be described later.

4 shows the results of observation of the microstructure of the cermets of Experimental Examples 1 to 4 by FESEM-BSE. Referring to (a) to (d) of Figure 4, the cermet according to the present experimental examples is composed of (Ti, W)C particles and a binder (binder) binding them in the vicinity. Referring to FIG. 4(a), in the case of Experimental Example 1 in which the amount of WC added was 10wt%, the dark rim area and the bright core area were clearly observed, and similar results were obtained in Experimental Example 2 in which the addition amount of WC was 20wt%. appear. When observed with FESEM-BSE, it provides brightness contrast by backscattered electron signals sensitive to atomic number. The bright area of the picture observed by FESEM-BSE corresponds to the area where the heavy element with large atomic number is concentrated. From this, in the case of Experimental Examples 1 and 2, it can be inferred that the relatively bright core region (white arrow) has a larger amount of tungsten (W) added than the dark rim region (black arrow). In addition, as shown in (a) and (b) of Fig. 4, the particle shape of the cermet has a relatively facet, and abnormal particles with a particle size significantly larger than the average were often observed.

On the other hand, in the case of Experimental Examples 3 and 4 in which the amount of WC added was 30wt% and 40wt%, a clear difference was shown in the microstructure compared to Experimental Examples 1 and 2 in which the amount of WC added was 20wt% or less. Referring to (c) and (d) of Figure 4, it can be confirmed that the particles having an abnormally large particle size are no longer observed and the particle shape has a relatively rounded shape. This means that the surface anisotropy of (Ti, W)C solid solution particles decreases with increasing WC content.

In addition, in the case of Experimental Example 3 in which the WC content was 30 wt%, the difference in brightness between the core region and the rim region was not noticeable, so that the boundary between the core region and the rim region was unclear. Furthermore, in the case of Experimental Example 4 in which the content of WC is 40wt%, the inversion of the brightness in which the core area is darker than the rim area appears. This means that in the case of Experimental Example 3, the W content in the core region and the rim region is similar, and in the case of Experimental Example 4, the W content in the rim region is rather higher than in the core region.

In the cermet according to the experimental example of the present invention, the sim structure is interpreted as that the core/rim structure is formed through complete solid solution and uniform precipitation of _{(Ti 0.88} W _{0.12 )C nanoparticles, and such a sim structure creation mechanism} Accordingly, it is determined that the structure exhibits a different structure from the conventional core/rim structure.

The core region of the conventional Ti(CN)-based cermet corresponds to partially dissolved Ti(CN) particles. Therefore, in the case of the conventional Ti(CN)-based cermet, the content of constituent elements was drastically different based on the core/rim boundary, and thus a large lattice strain was formed at the core/rim boundary. Due to this large lattice deformation, high stress was exhibited at the core/rim boundary.

However, in the case of the experimental example of the present invention, during sintering, (Ti _0.88 W _0.12 )C nanoparticles are completely dissolved in the metal binder and then uniformly precipitated to form a core/rim structure, and thus, unlike the prior art, at the core/rim interface does not show a sharp difference in the elemental content of This means that the stress at the core/rim boundary is considerably relaxed compared to the conventional core/rim structure.

5 shows the results of analyzing the local compositions of the phases present in the cermets corresponding to Experimental Examples 1 to 4 using a scanning transmission electron microscope - Energy dispersive spectroscopy (STEM-EDS), analyzed through STEM-EDS. High-angle annular dark-field (HAADF) results showing local spots. The HAADF analysis results provide a Z-contrast signal that is particularly sensitive to elements with high atomic numbers, so the bright regions correspond to regions where heavy metal elements are concentrated. Referring to FIGS. 5A to 5D , it can be seen that the brightness results in the core area and the rim area show the same trend as the results shown in FIG. 4 .

A portion indicated by a yellow circle in FIG. 5 is a region where the local composition is analyzed, and the analysis result is shown in FIG. 6 . Figure 6 (a) shows the content (at%) of W in the core region and the rim region, (b) shows the content (at%) of W and Ti in the Ni binder.

Referring to FIG. 6A , it can be seen that the W concentration in the core region is approximately constant regardless of the added WC content, while the W concentration in the rim region is greatly affected by the added WC content. In addition, in the cemet of Experimental Examples 1 to 4, the W content of the core region is about 40 at%, which is a completely different value compared to the _{(Ti 0.88} W _{0.12 )C powder inputted as a raw material in the cermet manufacturing step, which is a significantly higher value} indicates In order to distinguish this, the (Ti, W)C powder used as a raw material for manufacturing the cermet is referred to as a first perfect solid solution carbide, and (Ti, W)C perfect solid solution particles precipitated in the sintering step for manufacturing the cermet are used. It may be referred to as a second perfect solid solution carbide.

As described above, the distribution of W composition in the core region and the rim region shows that the (Ti _0.88 W _0.12 )C powder is completely dissolved in the binder during the sintering process and then uniformly precipitates again to form the core/rim structure of the cermet. because it forms The core region is formed by precipitation at a specific temperature during sintering, and thus the core region exhibits a composition close to thermodynamic equilibrium during formation, for example, 30 to 50 at%. Meanwhile, since the rim region is formed around the core region after the formation of the core region, when W is sufficiently present in the Ni binder, the W content in the rim region can be accommodated up to the limit of solubility.

Referring to (b) of FIG. 6 , as the content of added WC was higher, the content of W in the Ni binder was higher, and thus the content of Ti was relatively decreased. Similar to the W content in the Ni binder, the W content in the rim region also showed similar results according to the WC content. Specifically, referring to FIG. 6 (a), when the added WC content was 10 and 20 wt%, the concentration of W in the rim region was much lower than that in the core region, and the WC content was 30 wt%. In the case of , a value almost similar to the concentration in the core region was shown. When the WC was 40 wt%, the W content in the core region and the rim region was reversed, indicating a higher W content in the rim region. The results of this compositional analysis are consistent with the FESEM-BSE and HAADF analysis results shown in FIGS. 4 and 5 .

Referring to FIG. 5D , it can be seen that a relatively dark core region is formed in the center of the carbide particles, and a rim region, which is a relatively bright region, is formed in the periphery of the core region. As a result of analyzing the W component in the core region and the rim region, it was 41.4 at% in the core region and 47.2% in the rim region, indicating that the W content in the rim region is higher than in the core region. Accordingly, the Ti content in the core region is relatively higher than that in the rim region.

Line scanning was performed from the rim region of the carbide particles through the core region to the rim region again for the sintered body corresponding to Experimental Example 4 in order to confirm the composition distribution according to regions within the carbide particles. is shown in FIG. 7 .

Referring to FIG. 7 , as the rim region progresses to the core region, the content of W continuously decreases, and then maintains a lower value in the core region than in the rim region. Next, it can be confirmed that the content of W is continuously increased at the boundary between the core region and the rim region as it enters the rim region again. Accordingly, Ti shows a continuous increase/decrease pattern in the opposite direction to the change of W.

Referring to FIG. 7 , in the sintered body of Experimental Example 1, it can be seen that the W content in the core region is lower than that in the rim region. In particular, it can be seen that at the interface between the core region and the rim region, the composition change occurs continuously and gradually, rather than the rapid change in the W content in a step shape. Due to this interface structure, it is judged that the boundary between the core region and the rim region appears unclear when observing the carbide particles. Such an interface structure contributes to the improvement of the toughness of the cermet, which will be described later.

8 shows the results of observing the carbide in the sintered body corresponding to Experimental Example 4 with a transmission electron microscope, and performing electron diffraction analysis on the core region and the rim region. Fig. 8 (a) is an electron diffraction analysis result of the rim region, and (c) is an electron diffraction analysis result of the core region.

Referring to (a) and (c) of FIG. 8 , it can be seen that both the core region and the rim region show the same crystal structure of the NaCl-type face-centered cubic structure, which is consistent with the results of the X-ray diffraction analysis of FIG. 3 . am. However, the lattice constant in the core region was 4.2952 Å, whereas the lattice constant in the rim region was 4.40022 Å, indicating a larger lattice constant in the rim region. That is, the average lattice constant of the carbide is about 4.34 Å, which is almost identical to the result from the XRD peak analysis (FIG. 3), but shows a difference in the lattice constant locally in the carbide.

The difference in the lattice constant in the carbide particles is due to the difference in composition. As described above, the carbide particles are the perfect solid solution phase of TiC and WC, and have a crystal structure in which W replaces Ti atom sites in the TiC crystal structure. Since the atomic radius of W has a larger value than that of Ti, the lattice constant increases as the content of substituted W increases. As described above, the W content in the rim region has a higher value than that of the core region, and thus the lattice constant in the rim region has a larger value due to the difference in W content.

Due to the difference in lattice constant between the core region and the rim region in these carbide particles, the rim region is in a compressive stress state. That is, in a single-phase carbide having different lattice constants, compressive stress is received in a region having a large lattice constant, and a tensile stress is relatively formed in a region having a small lattice constant. In this experimental example, the lattice constant is larger in the rim region with a high W content, which leads to the formation of compressive stress in the large rim region. In the carbide structure having such a core structure, the toughness of the sintered body is improved as it has a compressive stress in the rim region.

In the cermet according to this experimental example, carbide particles, which are hard phases, form a single-phase completely solid phase, but have a core/rim core structure due to a difference in composition. However, at the interface between the core region and the rim region, the change in composition changes gradually and continuously, and if the rim region has a structure with compressive stress, the structure of these carbide particles acts as an important factor in improving toughness.

9 shows the results of measuring the Vickers hardness (Hardness) and fracture toughness (Toughness) of the cermet according to the WC content. The Vickers hardness was obtained by taking the average value for 5 press-insertions under 30 kg for 15 seconds of loading time, and the fracture toughness was derived using the length of the crack after measuring the Vickers hardness.

Referring to FIG. 9 , in the case of Vickers hardness, an almost constant value was maintained regardless of the change in the amount of WC added. On the other hand, in the case of fracture toughness, when WC was added at 30 wt % or more, the result was significantly improved. In particular, when the added amount is 40 wt%, the fracture toughness is about 22 MPam ^{1/2, which} is very high compared to the conventional one.

When the amount of WC added is 30 wt%, the W content of the core region and the rim region is similar as shown in FIG. is considered to appear. On the other hand, when the amount of WC added was 40 wt%, the fracture toughness was better than that when the amount of WC was 30 wt%.

10 shows the shape of a crack due to press-fitting of the cermet corresponding to Experimental Examples 1 to 4;

Experimental example Intragranular cracking rate (%) Intergranular crack rate (%) Experimental Example 1

Experimental Example 2

Experimental Example 3

Experimental Example 4

Referring to FIG. 10 , the prepared cermet may have inter-granular fracture (white arrow in FIG. 10) occurring along the boundary between the carbide and the binder depending on the content of the added WC or intragranular passing through the inside of the carbide particle. The shape of the crack (trans-granular fracture, black arrow in FIG. 10) was shown. Table 3 shows the ratio of crack types according to each experimental example. As shown in FIGS. 10(a) and 10(b), when the amount of WC added was 20 wt% or less, intragranular cracking mainly occurred, and as shown in FIGS. 10(c) and (d), the amount of WC added was In the case of 30 wt% or more, intergranular cracks were mainly observed. In particular, in the case of Experimental Example 4 in which the amount of WC added was 40 wt%, the ratio of intergranular cracks reached almost 90%.

The intragranular cracks shown in FIGS. 10 (a) and (b) showed a smooth shape that proceeded linearly along the cleavage plane of the brittle material, whereas the intergranular cracks shown in FIGS. 10 (c) and (d) exhibited a very rough shape typically observed in ductile materials.

The rate of intergranular cracking increases as the amount of WC added increases, which may be related to the increase in fracture energy along the cut surface of the carbide particles. This is considered to be because the lattice strain energy due to the difference in W concentration at the core/rim interface decreased as the amount of WC added increased.

On the other hand, the difference in W content at the core / rim interface showed a larger value in Experimental Example 4, in which the amount of WC was added at 40wt%, compared to Experimental Example 3, in which the amount of WC was added at 30wt%, but the rate of intergranular cracking was A larger value was shown in Experimental Example 4. This means that increasing the W content in the rim region can increase the intragranular fracture energy along the cleavage plane of the carbide particles.

The size and shape of the carbide particles can also influence the type of crack initiation and growth. The larger the size of the carbide particles, the more difficult it is to refract the propagating crack. As the carbide particle has a round shape, the refractive function of the propagating crack becomes higher. Therefore, as shown in (c) and (d) of Fig. 10, in the case of Experimental Examples 3 and 4, in which the size of the carbide particles are relatively small and have a round shape, the probability of occurrence of intergranular cracks increases.

In Experimental Examples 5 to 7 of Table 1, specific manufacturing compositions of cermets prepared by adding Co in addition to Ni as a binder or adding Mo _{2 C, NbC, and TaC in addition to WC as carbides to be added are shown.}

11 (a) to (c) are the results of observation with a scanning electron microscope of the microstructure of the cermet prepared from the powders corresponding to Experimental Examples 5 to 7, respectively. -> Modification has been completed.) Referring to FIG. 11 , in the cermets according to Experimental Examples 5 to 7, a relatively dark core region is formed at the center of the carbide particles, and a relatively bright region is formed around the core region. It can be seen that the rim region has a sim structure formed therein.

12 is a result of observing the carbide particles of the cermet of Experimental Example 7 with a scanning electron microscope (SEM) and analyzing the components of a specific part. Referring to FIG. 12 , in the cermet of Experimental Example 7 ((Ti,W)C-30WC-5Mo ₂ C-3NbC-2TaC-20Ni), a rim region, which is also a relatively bright region, is formed at the periphery of the core region. can confirm. Also, as in Experimental Example 4, the content of W in the core region showed a lower value than the content in the rim region, and the content of Mo, which is a Group VIa metal such as W, also exhibited a higher value in the rim region.

13 shows the Vickers hardness (Hv30) and toughness results of the specimens corresponding to Experimental Examples 4 to 7, along with the results of commercial WC-Co cemented carbide materials, commercial cermets, and cermets reported in the literature. The values indicated by 1), 2), 3) and 4) in the upper left of FIG. 9 are values corresponding to Experimental Examples 4, 5, 6 and 7, respectively. Referring to FIG. 9 , the cermet according to the experimental example of the present invention exhibits superior toughness compared to the conventional commercial cermet or commercial WC-Co cemented carbide. It can be confirmed that the hardness value is similar or slightly smaller than that of the cermet reported in the literature, but the toughness shows a superior value. Therefore, it is expected that the application of the cermet can be greatly expanded, unlike several conventional cermets that have been used in limited applications due to their low toughness. As described above, the cermet using the manufacturing method of the present invention can greatly improve the problem of low toughness of the existing cermet and cemented carbide materials. ^{(Hardness: 10~15GPa, toughness: 15~40MPa·m 1/2} level can be secured depending on the adjustment of hardness and toughness by adjusting the binder and added carbide content)

Since the present invention has significantly improved toughness unlike conventional cermets, it is likely to be applicable to various extreme environment fields requiring high toughness (defense, aviation, space, marine, nuclear fusion power generation) and bearings. Specific examples include high-temperature structural materials, armor for tanks, etc., bullets from kinetic energy bombs, and materials for extreme environments inside nuclear fusion reactors. In addition, when the insufficient hardness is compensated for by applying the _{ceramic coating (Al 2} O ₃ , TiN and TiAlN, etc.) that is generally performed on the tool material to the cermet base material manufactured in the present invention, a longer life is secured than the existing tool material. It looks like it could be used as a tool.

Although the present invention has been described with reference to the embodiment shown in the drawings, which is merely exemplary, those skilled in the art will understand that various modifications and equivalent other embodiments are possible therefrom. Therefore, the true technical protection scope of the present invention should be determined by the technical spirit of the appended claims.

Claims (7)

It is formed from the first perfect solid solution carbide, which is a perfect solid solution phase of carbides of two or more metals selected from Group IVa, Group Va, and Group VIa metals, including Ti, in the periodic table, and has a nano-sized core structure composed of a core region and a rim region. particle; and
A binder made of metal; including,
The first perfect solid solution carbide is _{a perfect solid solution of TiC and WC prepared from TiO 2} , WO ₃ , and C (graphite), in which a part of the Ti atoms are substituted with W in the lattice of TiC having a NaCl-type face-centered cubic structure. have a crystal structure,
The nano-sized particles are
As a second full solid solution carbide having a single phase completely dissolved in solid solution, as the composition of W in the core region is lower than that in the rim region, the lattice constant in the rim region is the lattice constant in the core region has a larger value than, the rim region has a compressive stress state compared to the core region,
The boundary between the core region and the rim region is unclear,
In the interface, as the composition of W gradually increases from the core region toward the rim region, the compressive stress also gradually changes.
Cermet with improved toughness. The method of claim 1,
wherein the difference in composition of the group VIa in the core region and the rim region is between 1 at% and 10 at%,
Cermet with improved toughness. The method of claim 1,
The binder includes any one or more of Ni, Co, and Fe.
Cermet with improved toughness. preparing a first perfect solid solution carbide in a perfect solid solution phase from at least two metal carbides selected including Ti from group IVa and VIa metals in the periodic table; and
Sintering the mixed powder comprising the first complete solid solution carbide powder, an additive including one or more metal carbide powders selected from Va and group VIa metals, and a binder including one or more metal powders,
The step of preparing the first perfect solid solution carbide,
Including; TiO ₂ , WO ₃ , performing high-energy milling on C (graphite);
The sintering step is
The first completely solid solution carbide powder is completely dissolved in the binder; and
and precipitating a second perfect solid carbide having a different composition from the first perfect solid carbide from the binder,
The step of precipitating the second perfect solid solution carbide,
precipitating the core region; and
After the core region is precipitated, a rim region having a relatively higher composition of Group VIa metal than the core region is precipitated around the core region,
The first perfect solid solution carbide is
TiO ₂ , WO ₃ , a complete solid solution of TiC and WC prepared from C (graphite), and has a crystal structure in which a part of the Ti atoms are substituted with W in the lattice of TiC having a NaCl-type face-centered cubic structure,
The second perfect solid solution carbide is
As a solid solution having a single phase completely dissolved, the boundary between the core region and the rim region is unclear,
In the interface, as the composition of W gradually increases from the core region toward the rim region, the compressive stress also gradually changes,
A method for manufacturing a cermet with improved toughness. 5. The method of claim 4,
The average size of the perfect solid solution carbide powder has a range of 60 nm or less (greater than 0),
A method for manufacturing a cermet with improved toughness. 5. The method of claim 4,
The metal powder includes any one of Ni powder, Co powder and Fe powder,
A method for manufacturing a cermet with improved toughness. 5. The method of claim 4,
The additive comprises _{any one or more of WC, Mo 2} C, TaC, NbC and VC,
A method for manufacturing a cermet with improved toughness.

Images