JP2024518384A - Manufacturing method of sintered carbonized body - Google Patents

Abstract

The present invention relates to a method for the manufacture of sintered carbide bodies, comprising mixing tungsten carbide powder (WC powder) with a metallic binder material comprising cobalt powder (Co), nickel (Ni) and aluminium (Al) in a mixing and/or grinding process, preferably a wet grinding process, forming a powder mixture, producing a green compact from at least a part of the powder mixture by pressing, sintering the green compact under the action of temperature and pressure in a sintering step, and forming a sintered carbide body after a cooling step following the sintering step.To create an easily controllable and reliable method for the manufacture of sintered carbide bodies, characterized by an increased wear resistance and at the same time a high fracture strength, according to the present invention, it is provided to add nickel aluminide as intermetallic phase material to the mixing and/or grinding process, preferably nickel aluminide powder, in particular Ni3Al powder.

Description

The present invention relates to a method for producing a sintered carbide body, comprising mixing tungsten carbide powder (WC powder) with a metal binder material comprising cobalt powder (Co), nickel (Ni) and aluminium (Al) in a mixing and/or grinding process, preferably a wet grinding process, to form a powder mixture, producing a green compact from at least a part of the powder mixture by pressing, sintering the green compact under the action of temperature and pressure in a sintering step, and forming a sintered carbide body after a cooling step following the sintering step.

EP 2 691 198 describes a sintered carbide material, i.e. a hard alloy body (tungsten carbide material) and a method for its production. According to this known method, a powder containing coarse-grained tungsten carbide and a larger than stoichiometric proportion of carbon and cobalt powder is mixed. Powdered tungsten is then added to this powder. The tungsten powder and the cobalt powder had an average grain size of about 1 μm. The coarse-grained tungsten carbide had an average grain size of 40.8 μm.

The powder was then ground in a ball mill, to which hexane and paraffin wax were added. A green compact was produced from the mixture by pressing, and the green compact was then sintered. After the sintering process, the resulting sintered carbide material was subjected to a heat treatment, in which it was heated to 600°C and held at this temperature for 10 hours.

After a subsequent cooling step, the cemented carbide material was subjected to analysis, where it was found to have nanoparticles in the binder phase, the nanoparticles having a size of less than 10 nm. The nanoparticles were formed by eta phase ( _Co3W3C ) or ( _Co6W6C ) or _theta phase ( _Co2W4C ). The grain size of the nanoparticles was less than _{10 nm} .

Nanoparticles have been found to be associated with a strengthening of the binder phase, which can increase the hardness of cemented carbide materials. The drawback of these materials is the lack of thermal stability of nanoparticles, which makes them only suitable to a limited extent for high temperature applications or applications where high temperature inputs are present.

During rock processing and cutting of asphalt and concrete, friction generates very high temperatures on the tool surface. The hard material tungsten carbide is not significantly affected by these temperatures, as it has a high hot hardness even at these temperatures. However, the strength of the metal binder decreases dramatically at these temperatures. When the strength of the metal binder decreases, the stresses of use result in increased wear and tear and dissolution of the binding phase. As a result, the tungsten carbide grains can no longer be held in the hard alloy.

The objective of the present invention is to create an easily controllable and reliable method for producing sintered carbide bodies that are characterized by high fracture strength while at the same time having improved wear resistance.

This problem is solved by adding nickel aluminide, preferably nickel aluminide powder, in particular Ni ₃ Al powder, as intermetallic phase material to the mixing and/or milling process.

For the purposes of this invention, nickel aluminide can be an intermetallic phase material that contains at least Ni and Al, where the Ni and Al are bonded to each other in the crystal structure.

Within the scope of the present invention, nickel aluminides are also understood to mean alloyed nickel aluminides (e.g. nickel aluminides alloyed with tin).

In the sense of the present invention, nickel aluminide is also understood to mean an amorphous material, especially in powder form.

Nickel aluminide can be added to the mixing or grinding process without the need for special precautions in terms of work safety or worker protection. In particular, no deoxidation is required, as would be necessary if Al were added in elemental form. Furthermore, this allows the nickel aluminide to be added to the mixing and/or grinding process in precisely metered amounts, which in turn facilitates reliable and reproducible production.

The method according to the invention makes it possible to produce cemented carbide materials, in particular hard alloys (e.g. known as "hard metals"), which already have a reinforcing binder phase and/or are prepared to form a reinforcing binder phase. In both cases, the binder phase is strengthened by the intermetallic phase material resulting from the Ni and Al content of the added nickel aluminide.

According to one preferred variant of the invention, it can be provided that the cooling step and/or the heat treatment of the sintered carbide body subsequent to the sintering step is controlled so that in the sintered carbide body, intermetallic phase materials are formed in the binder phase of the sintered carbide body, at least part of which is preferably formed according to the structural formula (M,Y) ₃ (Al,X), where M=Ni, Y=Co and/or other elements and X=tungsten and/or other elements.

Here, it is advantageously possible to provide that during the cooling step, the sintered body is held at a temperature ranging between 400°C and the solvus temperature of the sintered body for 0.25 to 24 hours, in order to ensure that intermetallic phase material is formed in sufficient quantity and size to achieve an effective strengthening of the binder phase.

If intermetallic phase material is already formed during the cooling process and is present in the bonding phase, this immediately strengthens the bonding phase. This makes the process particularly easy to operate.

When the cemented carbide material has dissolved elements Ni and Al in the binder phase after the sintering process, the cemented carbide material can be heat treated to form an intermetallic phase material that strengthens the binder phase and provides the desired improved wear resistance.

During this heat treatment, for example, at least a portion of the binder phase elements Ni, Co, W and Al combine to form an intermetallic phase material, at least a portion of which may be formed according to the formula (M,Y) ₃ (Al,X), where M=Ni, Y=Co and/or other elements, and X=tungsten and/or other elements.

Here, heat treatment of the cemented carbide material can be carried out in a variety of ways suitable for forming the intermetallic phase material as desired.

Here, in particular, there may be a heat treatment, in particular the introduction of heat from outside into the sintered carbide material. The heat treatment may, for example, be caused by an active source generating heat or cold, which introduces heat into the material or removes heat from the material. Thus, for example, the formation of the intermetallic phase material may take place in a furnace into which the sintered carbide material is introduced. It is also conceivable to expose at least a part of the surface of the sintered carbide material to a heating device, for example a burner.

It is also contemplated that there is an excitation source that introduces energy into the cemented carbide material to generate heat in the cemented carbide material. This may be, for example, an induction coil or a laser device.

It is also conceivable to generate heat by passive heating, i.e. by using the sintered carbide material, preferably during the intended operating conditions or during the processing, in particular the method steps, in which the sintered carbide material is used.

Here, heat is generated in the sintered carbide material, for example by friction, as occurs, for example, during the intended use of the sintered carbide material, in particular during the intended use, in particular during the use of a tool. In the case of a tool, the tool is moved relative to the workpiece (for example, in the case of a road cutting chisel, the road cutting chisel is moved relative to the road surface), which generates frictional energy and thus heat in the sintered carbide material. The heat generated in this case can be used to achieve a self-reinforcing effect of the binder phase by at least partial formation of intermetallic phase material in the sintered carbide material.

The intended operating conditions can also be understood as the use of a cemented carbide material suitable for operation at an operating temperature suitable for the formation of intermetallic phase material.

It is also conceivable that the sintered carbide material is passively heated during the process step in which it is applied to a support, for example a tool base or carrier. The heat generated in this case can be used to form an intermetallic phase material. Possible joining methods here are welding methods, for example friction welding, electron beam welding, brazing methods, for example hard brazing, furnace brazing, induction brazing, diffusion brazing, cladding methods, for example explosive bonding.

The intermetallic phase material forms crystalline intercalations in the metal binder.

This intermetallic phase material has significantly higher strength, especially at high temperatures, than the metal binder material in which it is intercalated. At the surface of the cemented carbide material subjected to wear attack, the intermetallic phase material reduces erosion and dissolution of the metal binder material, for example when used in earth-excavating tools.

The movement of the soil excavation tool and the loosened and remaining soil materials cause abrasive and mechanical loads on the cemented carbide material. The tungsten carbide grains in the cemented carbide material are then sufficiently resistant to this abrasive attack. The problem in the prior art is that the strength of the binder material is significantly lower than that of tungsten carbide. Now, the intermetallic phase material is incorporated or formed in the binder phase, which prevents rapid erosion and dissolution of the metallic binder material.

Furthermore, it has surprisingly been found that the internal structure of the cemented carbide material can also be strengthened by the intermetallic phase material. In the event of strong shock-like stresses, the crystals of the intermetallic phase material reduce or prevent the sliding of the tungsten carbide grains in the region of the binder phase connecting them and thus excessive plastic deformation of the binder phase. In particular, in this case, the individual crystals of the intermetallic phase material support one another. This is of great advantage, especially at high tool operating temperatures. For example, if Co is used in the binder phase, at such temperatures the strength of Co in the binder phase decreases, but the intermetallic phase material still ensures that it provides sufficient support for the binder material.

Overall, it has been found that the method according to the invention allows a significant increase in the wear resistance of the finished sintered carbide material to be achieved. For example, tests have shown that when sintered carbide material is used in the form of cutting tips in round shank chisels for road milling machines, the wear resistance can be increased by up to 50%. It has been found that such a significant increase in wear resistance can be achieved when milling asphalt and concrete road surfaces.

Sintered carbide materials can in particular be used to construct the working areas of tools for the treatment, loosening, transport and processing of plant, mineral or building materials, in particular in agriculture, forestry or in the fields of road construction, mining or tunnel construction.

If intermetallic phase materials are present in the sintered carbide material, then according to one variant of the invention it can be provided that the proportion of metal binder material in the sintered carbide material is 1 to 28% by weight, preferably 1 to 19% by weight. In this case, the entire or almost entire proportion of this metal binder material, with the exception of unavoidable impurities, can be formed by Co. Such a material selection results in a particularly tough binder phase, which is effectively reinforced by existing or forming intermetallic phase materials.

According to one variant of the invention, it can be provided that the sum of the elements Ni and Al in the cemented carbide material is 1-28% by weight, advantageously 1.5-19% by weight, in at least one region of the cemented carbide body. This range of data takes into account both the Ni and Al of the intermetallic phase material that may be present, and also the Ni or Al dissolved in the binder phase. Such a composition makes it possible to manufacture earth-digging tools made of hard alloys, for which particularly high demands are imposed.

In order to be able to produce such a sintered carbide material, according to one possible variant of the invention, it can be provided that the green compact contains, as intermetallic phase materials, 70-95% by weight, preferably 80-95% by weight, of tungsten carbide (WC), 1-28% by weight, preferably 1-19% by weight, of cobalt (Co) and 1-28% by weight, preferably 1.5-19% by weight, of nickel aluminide.

If no or only little intermetallic phase material is present in the sintered carbide material, then according to one variant of the invention it can be provided that the proportion of binder phase in the sintered carbide material is 5 to 30% by weight, preferably 5 to 20% by weight. In this case, half or the majority of the binder phase may be formed by the metallic binder material Co. Furthermore, Al and Ni may be dissolved in the binder phase as metallic binder materials. Finally, further elements and unavoidable impurities may also be present in the binder phase.

According to the present invention, it is also possible to provide that the binder phase contains other elements besides Co, Ni and Al, in particular dissolved W, C and/or Fe, excluding inevitable impurities.

In accordance with the present invention, the intermetallic phase material optionally present in the binder phase may be formed according to the structural formula (M,Y) ₃ (Al,X), or the intermetallic phase material may be formed according to this structural formula based on the elemental composition in the finished cemented carbide material, where M=Ni, Y=Co and/or other components, and X=tungsten and/or other components.

When intermetallic phase material is present in the binder phase, it is advantageous that Y=Co and X=W, at least in the majority of the crystals of the intermetallic phase material. The composition of the soluble components in the binder phase in the cemented carbide material is therefore selected so that the intermetallic phase material can thus form by heat treatment or the action of heat (see above).

Furthermore, for some or all of the crystallites of the intermetallic phase material (Al,X), X can be present not only in the form of W, but also in the form of Mo and/or Nb and/or Ti and/or Ta and/or Cr and/or V. The composition of the dissolved components in the binder phase in the finished cemented carbide material can therefore be selected such that the intermetallic phase material can thus be formed by heat treatment or by the action of heat (see above).

According to the invention, it is possible to provide that the binder phase comprises two or more intermetallic phase materials or only a single intermetallic phase material, and/or that the cemented carbide material is prepared such that two or more intermetallic phase materials or only a single intermetallic phase material is formed by heat treatment or the action of heat.

According to a preferred embodiment of the invention, the binder phase has the following chemical element composition:
Ni>25 wt.%, Al>4 wt.%, balance Co and dissolved binder components such as W and/or C;
Advantageously, Ni>35% by weight, Al>5% by weight, balance Co and dissolved binder components such as W and/or C;
Particularly preferred is a composition containing Ni>40 wt.%, Al>6.5 wt.%, the remainder Co and dissolved binder components such as W and/or C.
It is possible to provide for the device to have the following:

Here, these data relate to the total content of the respective substances in the binder phase. These data therefore take into account the respective elements in dissolved form and/or also present in a bonded state in the intermetallic phase material.

The specified data values were found to provide an effective strengthening effect of the bonding phase.

In order to be able to produce an optimal strengthening of the binder phase by the intermetallic phase material, it is possible within the scope of the present invention to provide that the ratio of the weight percentage of Al to the weight percentage of Ni is >0.10, advantageously >0.12.

Here, it is advantageously possible to provide that the ratio of the weight proportion of Al to the weight proportion of Ni is ≦0.46, preferably ≦0.18, particularly preferably ≦0.16.

The above data on the ratio of the weight percentage of Al to the weight percentage of Ni take into account the total weight percentages, i.e. both the dissolved elements Al and Ni and the Al and Ni in the intermetallic phase material (if present). The detection of the weight percentages of Al and Ni can be carried out by conventional ICP measurements.

One possible variant of the invention is that the sintered carbide material has at least two volume regions, the first volume region having a higher relative proportion of intermetallic phase material per unit volume than the second volume region. The configuration of at least two volume regions allows targeted influence of the properties of the sintered carbide material, in particular the tool properties. For example, a zone exposed to high levels of wear and abrasion can have a volume region with a higher relative proportion of intermetallic phase material. In contrast, a zone that has to meet special toughness requirements can have a volume region with a lower relative proportion of intermetallic phase material. For completeness, it is pointed out here that the second volume region, which has a lower relative proportion of intermetallic phase material compared to the first volume region, may not contain any intermetallic phase material.

In one possible alternative form of the invention, the first volume region having a high relative proportion of intermetallic phase material is arranged to be bounded by at least one region of the surface of the cemented carbide material. In this way, high wear resistance is created at the surface of the cemented carbide material. Advantageously, the second volume region is not located adjacent to the surface of the cemented carbide material, but is present in the interior of the cemented carbide material. This ensures high fracture stability.

It is also possible to provide that the second volume region, which has a lower relative proportion of intermetallic phase material compared to the first volume region or is free of intermetallic phase material, borders at least one region of the surface of the cemented carbide material, and preferably the first volume region is not adjacent to the surface of the cemented carbide material. This makes it possible to produce a tool whose cutting contour is optimally adapted to the cutting task during tool use in order to achieve a good service life. In particular, a so-called re-sharpening effect can be realized.

A possible variant of the invention is that the binder phase, in particular the metallic binder material and/or the intermetallic phase material, comprises Nb and/or Ti and/or Ta and/or Mo and/or V and/or Cr, advantageously one or more of these materials being present in the binder phase in dissolved form and/or in the form of carbides. This allows for an increase in the solvus temperature and strength of the existing intermetallic phase material and/or the intermetallic phase formed by the action of heat or heat treatment to be achieved. As a result, less intermetallic phase material is required while the strength of the sintered carbide material remains the same. Alternatively, the addition increases the binder strength and thus the hot strength.

However, it is also contemplated that one or more of the aforementioned components are incorporated into at least a portion of the crystal lattice of the intermetallic phase material, or can be incorporated into the crystal lattice of the intermetallic phase material by thermal treatment (heat treatment or the action of heat). For example, titanium atoms (or another member of the aforementioned group) will occupy primarily Al or W lattice sites in the crystal lattice of the intermetallic phase material, and similarly to W, will increase the precipitation temperature of the intermetallic phase material.

Thereby, if intermetallic phase materials are present in the sintered carbide material, on the one hand, they can be precipitated more effectively during the sintering process, since precipitation starts already at high temperatures, since in this case the diffusion rate is significantly higher.

On the other hand, as already mentioned, this measure achieves higher hot strength, since the solvus temperature of the cemented carbide material is increased; in other words, the temperature required to be able to remelt the intermetallic phase material in the cemented carbide material is increased.

According to the invention, it is also possible to provide that the proportion of Mo and/or Nb and/or Ti and/or Ta and/or Cr and/or V in the binder phase is ≦15 at%. The above-mentioned elements form carbides as a rule. Within the scope of the invention, it is now possible to provide that the material composition is selected such that these elements are dissolved in small amounts in the binder phase depending on their solubility product and their affinity for carbon, so that they can be incorporated into the crystal lattice of the intermetallic phase material and/or that they can dissolve in the metallic binder phase. If a sintered carbide material with a high toughness of the binder phase is desired, it is desirable to limit the proportion of carbide forms to small amounts. In that case, it is desirable for the material to be present in a total proportion of ≦15 at%.

Furthermore, it is advantageous that the powder mixtures for the production of sintered carbide materials (in particular the powder mixtures of green compacts) can be adjusted stoichiometrically with respect to the carbon proportion, since titanium (and/or Mo and/or Nb and/or Ti and/or Ta and/or Cr and/or V) takes over the role of tungsten.

According to one variant of the invention, it is possible to provide that the carbon content in the powder mixture (in particular the powder mixture of the green compact) is adjusted to be stoichiometric or less than stoichiometric. This measure prevents or minimizes graphite precipitation in the sintered material, which would be caused by a carbon content greater than stoichiometric. The inventors have recognized that such intercalation has a detrimental effect on the fracture strength of the sintered carbide material.

According to the invention, the carbon percentage in the cemented carbide material is
C _{stoichiometric amount} (wt%)-0.003 x binder content (wt%) to C _{stoichiometric amount} (wt%)-0.012 x binder content (wt%), advantageously
A range from C _{stoichiometric amount} (wt%) - 0.005 x binder content (wt%) to C _{stoichiometric amount} (wt%) - 0.01 x binder content (wt%) is particularly provideable.

In the context of the present invention, the above-mentioned advantageous effects are particularly pronounced in the case of coarse-grained hard alloys. In one preferred variant of the present invention, it is therefore provided that the dispersed tungsten carbide is present in the sintered carbide material in the form of grains whose average grain size, measured according to DIN ISO 4499 Part 2, is in the range of 1 to 15 μm, preferably in the range of 1.3 to 10 μm, particularly preferably in the range of 2.5 to 6 μm.

Advantageously, it is provided that the maximum proportion of Fe in the binding phase or in the green compact is 5% by weight and/or that other unavoidable impurities are present in the binder material.

If it is provided that the intermetallic phase (M,Y) ₃ (Al,X), existing or formed by heat treatment or the action of heat, has the crystal structure _L12 (space group 221) according to ICSD (Inorganic Crystal Structure Database), when the sintered carbide body is subjected to high loads, a microstructure of the binder phase arises in which the crystals of the intermetallic phase can effectively support each other in the metallic binder material.

Advantageously, and preferably for use in earth excavation tools, it is provided that the intermetallic phase material, existing or formed by heat treatment or the action of heat, has a maximum size of 1500 nm, advantageously a maximum size of 1000 nm.

According to one preferred configuration variant of the invention, it can be provided that the cemented carbide material is free or as free as possible from eta phase and/or Al ₂ O _3. The inventors have realized that it is desirable for the maximum proportion of eta phase or the maximum proportion of Al ₂ O ₃ to be at most 0.6% by volume relative to the total cemented carbide material. If both substances are present in the cemented carbide material, it is advantageous for the sum of eta phase material and Al ₂ O ₃ to be at most 0.6% by volume.

The grain size of the _Al2O3 and _/ or eta phase material is preferably at most 5 times the average grain size of the WC, whereby the average grain size of the WC and the grain size of _{the Al2O3} and/or eta phase material can be determined by the intercept method (in accordance with DIN ISO 4499 Part 2).

The toughness of the cemented carbide material can be adversely affected by the eta phase or _{by Al2O3} . If the proportion of eta phase is high, the cemented carbide material is conditionally suitable for use in earth-digging tools where high demands are placed on it. _The same is true for _Al2O3 .

The above-mentioned object of the present invention is also achieved by a method for manufacturing a tool, in particular a grinding tool, a soil excavation tool, preferably for a road mill, a recycler, a stabilizer, an agricultural or forestry soil excavator, which comprises a base body having a working area, in which at least one working element made of a sintered carbide material manufactured according to any one of claims 1 to 22 is held, preferably by a material connection, in particular by a brazing connection, in particular by a hard brazing connection.

Advantageously, the sintered carbide material in the working area forms a cutting body having a cutting tip, blade or cutting or working edge. It is also conceivable that the sintered carbide material is a protective covering.

Thus, as mentioned above, it is possible according to the invention to provide that the working element is advantageously formed in the form of a cutting element having at least one cutting edge and/or at least one cutting tip, or in the form of a wear protection element, in particular a protective plate, a protective bar, a protective pin, a protective lug or a protective bolt.

Particularly preferred applications of the present invention arise when it is provided that the tool is a cutting tool, a cutting chisel, in particular a road cutting chisel or a mining chisel, a ploughshare, a cultivator tip, a drill tool, in particular an earth drill, a crushing tool, for example a crushing chisel or a crushing bar, a mulching tool, a wood chipping tool or a shredder tool, a sorting tool, for example a screen.

A further particularly preferred use of the present invention is configured such that the cutting chisel has a chisel head and a chisel shaft connected directly or indirectly thereto, and a working element is held on the chisel head.

For example, it is possible to provide that the working element is formed by a sintered carbide material according to the invention, which forms a support for an ultra-hard cutting tip, for example made of PCD material.

As noted above, the cemented carbide material can be a hard alloy having a strengthening binder phase. This strengthening can be achieved by precipitation of intermetallic phase material during cooling in the sintering process and/or can be achieved when the intermetallic phase material is formed in a thermal process subsequent to the sintering process, where the cemented carbide material is brought to a temperature that allows for the precipitation of intermetallic phase material in the cemented carbide material.

For the production of hard alloys according to the invention, the nominal composition of the raw materials at the initial weighing may be selected as follows: 70-95 wt. % WC, 1-28 wt. % metallic binder and 1-28 wt. % nickel aluminide (e.g. intermetallic phase). The metallic binder may contain the element Co, and optionally Fe and/or other components. The intermetallic phase is advantageously _Ni3Al at the initial weighing.

The subject of the invention is also a method for producing a cemented carbide material, in which in a first method step a cemented carbide material precursor, in particular a hard alloy, is produced which comprises 70-95% by weight, preferably 80-95% by weight, of tungsten carbide in dispersed form and a binder phase, the binder phase comprising a metallic binder material which comprises Co, the binder phase comprising the dissolved elements Ni and Al, the binder phase having the following chemical element composition:
Ni>25 wt.%, Al>4 wt.%, balance Co and dissolved binder components such as W and/or C;
Advantageously, Ni>35% by weight, Al>5% by weight, balance Co and dissolved binder components such as W and/or C;
Particularly preferred is a composition containing Ni>40 wt.%, Al>6.5 wt.%, the remainder Co and dissolved binder components such as W and/or C.
and in a further method step subjecting the cemented carbide material precursor to a heat treatment to form a cemented carbide material comprising an intermetallic phase material in a binder phase, wherein at least a portion of the intermetallic phase material is formed according to the structural formula (M,Y) ₃ (Al,X), where M=Ni, Y=Co and/or other constituents, and X=tungsten and/or other constituents.

Here, the heat treatment can include at least one heating step or at least one cooling step.

According to the invention it is advantageously possible to provide that after the heat treatment at least some of the crystals of the intermetallic phase material (M,Y) ₃ (Al,X) have the crystal structure L12 (space group 221) according to ICSD (Inorganic Crystal Structure Database).

In order to be able to achieve a particularly effective strengthening of the binder phase, in particular in coarse-grained hard alloys as cemented carbide materials, it is advantageously possible to provide that at least a portion of the intermetallic phase material has a maximum size of 1500 nm (measured by the section method on microscopic specimens), advantageously a maximum size of 1000 nm.

Advantageously, it can be provided that the coercivity HcM of at least one region of the cemented carbide material produced by the method according to the invention, advantageously in the region where the intermetallic phase material is present, is:
HcM [kA/m]>(1.5+0.04×B)+(12.5−0.5×B)/D+4 [kA/m],
Advantageously, HcM [kA / m] > (1.5 + 0.04 × B) + (12.5 - 0.5 × B) / D + 6 [kA / m],
Particularly preferably, HcM [kA/m]>(1.5+0.04×B)+(12.5−0.5×B)/D+10 [kA/m]
where B is the proportion of binder phase in the cemented carbide material (wt.%) and D is the grain size of the dispersed WC determined by the intercept method according to DIN ISO 4499 Part 2.

For typical hard alloys containing Co in the binder phase and no intermetallic phase material, the coercivity is usually used to indirectly determine the average grain size of WC at a given binder content. According to the present invention, the intermetallic phase material significantly increases the coercivity. Thus, the coercivity can be indirectly evaluated as an indicator of the strengthening of the binder phase by the intercalated intermetallic phase material. The higher the coercivity, the greater the sum of the interfaces between the metal binder material, the intermetallic phase material and the WC. The greater the number of precipitated intermetallic phase materials, the better the support of the individual crystals of the intermetallic phase material to each other in the binder phase, especially at high temperatures (especially at high tool use temperatures).

At least one region of the sintered carbide material has a coercivity HcM [kA/m] > (1.5 + 0.04 x B) + (12.5 - 0.5 x B)/D + 4 [kA/m], which can be used primarily for the above-mentioned wear protection applications, e.g. wear protection.

The coercivity of at least one region of the sintered carbide material, preferably HcM [kA/m] > (1.5 + 0.04 x B) + (12.5 - 0.5 x B)/D + 6 [kA/m], can be used primarily in the above-mentioned earth-digging tools where high demands are imposed.

The coercivity of at least one region of the sintered carbide material, preferably HcM [kA/m] > (1.5 + 0.04 x B) + (12.5 - 0.5 x B)/D + 10 [kA/m], can be used primarily for the high-performance tools mentioned above.

According to one variant of the invention, the coercivity of at least one region of the cemented carbide material is 20% higher than the coercivity of a hard alloy body having the same composition and WC grain size as the cemented carbide material, and it is also possible to provide that the binder phase is formed only from a metallic Co binder, which does not contain any intermetallic phase material.

Thus, a hard alloy body having the same composition is a hard alloy body comprising 70-95% by weight of dispersed tungsten carbide and a binder phase, the binder phase comprising a metallic binder material but no intermetallic phase material, the proportion of metallic binder material in the sintered carbide material being 5-30% by weight, the binder material otherwise having the same or nearly the same composition as the binder material of the sintered carbide material according to the invention.

As mentioned above, coercivity indirectly indicates the degree of strengthening of the binder phase.

In the scope of the present invention, the sintered carbide material produced by the method according to the present invention may be configured such that the sintered carbide material has a hot compressive strength of ≥ 1650 [MPa] at a temperature of 800 ° C and a strain rate of 0.001 [1 / s] and / or such that the sintered carbide material has a hot compressive strength of ≥ 1600 [MPa] at a temperature of 800 ° C and a strain rate of 0.01 [1 / s] (measured on cylindrical test pieces with a diameter of 8 mm and a height of 12 mm). With a sintered carbide material in which the proportion of the binder phase is 5-7 wt. % and the proportion of WC is in the range of 93-95 wt. %, advantageously in this case WC is present in the form of coarse grains with an average grain size in the range of 2-5 μm, cutting tips, in particular for road cutting chisels, can be produced.

A particularly preferred variant of the invention provides for producing nickel aluminide, preferably _Ni3Al powder, in a molten metallurgical process and/or adding nickel aluminide, preferably _Ni3Al powder, as a material produced in a molten metallurgical process to the mixing and/or grinding process.

Particularly good processability of the powder mixture in the mixing and/or grinding process is obtained if provision is made to add nickel aluminide, preferably _Ni3Al powder, with an average particle size FSSS<70 μm, preferably with an average particle size FSSS<45 μm, to the mixing and/or grinding process.

According to one variant of the invention, it is possible to provide that in the preparation step, advantageously in the first grinding step, nickel aluminide is mixed with a grinding liquid and with coarse-grained tungsten carbide, preferably with a mean grain size FSSS of >20 μm, particularly preferably with a mean grain size FSSS in the range of 30-60 μm, for example in the form of macrocrystalline and/or monocrystalline tungsten carbide, in order to form crushed nickel aluminide, advantageously crushed nickel aluminide powder, in particular crushed Ni ₃ Al powder, from the nickel aluminide, whereby an effective comminution of the nickel aluminide is achieved.

Here, the method operations may, for example, consist of adding a pressing aid, at least one alloying component and/or cobalt powder and mixing this with the nickel aluminide and/or crushed nickel aluminide during the preparation step and/or the subsequent grinding step.

To achieve particularly good grinding performance, it is advantageous to provide that the proportion of nickel aluminide in the grinding mixture of the preparation process is in the range of 8 to 50% by weight, preferably in the range of 9 to 25% by weight.

It is particularly preferred to add WC powder to the pre-ground product obtained in the preparation step in a subsequent grinding step so that the proportion of WC powder in the resulting ground mixture is in the range of 70 to 95% by weight, and to grind the crushed nickel aluminide in the subsequent grinding step to obtain finely ground nickel aluminide.

Advantageously, the finely crushed nickel aluminide makes the distribution of the intermetallic phase in the sintered body more uniform. Furthermore, the WC structure, especially the WC coarse grain structure, is better preserved in the WC powder, since the milling process of the added WC is shorter.

Advantageously, according to one variant of the invention, the green compact is sintered in a furnace by a liquid phase sintering process at a sintering temperature in the range of 1350°C to 1550°C.

A particularly effective strengthening of the binder phase can be achieved if it is provided that at least a portion of the intermetallic phase material in the binder phase has a maximum size of 1500 nm, advantageously a maximum size of 1000 nm (measured by sectioning on microscopic specimens) and/or if it is provided that at least a portion of the crystals of the intermetallic phase material (M,Y) ₃ (Al,X) in the binder phase have the crystal structure L12 (space group 221) according to ICSD (Inorganic Crystal Structure Database).

One possible method variant is to process atomized nickel aluminide, preferably atomized _Ni3Al powder, as the intermetallic phase material in the milling and/or mixing process. This material is easy to handle. It requires less milling effort to obtain the desired fine grain structure.

Advantageously, the mixing and/or grinding process is a multi-stage process having at least two mixing and/or grinding steps, and advantageously it is possible to provide that the nickel aluminide is added before the last grinding and/or mixing step.

Manufacturing (including explanation of measurement methods)
Production:
The following describes how cemented carbide materials containing intermetallic phase material in the binder phase can be produced by a powder metallurgy process routine, which is divided into the process steps of producing a pressable powder mixture, compacting, and finally sintering to obtain a compact dense cemented carbide body.

As starting material for the production of the powder mixture, WC powders of various grain sizes can be used, in particular coarse-grained WC with a grain size FSSS>25 μm. The starting powders of the binder phase are ultrafine-grained cobalt powder (FSSS 1.3 μm) and nickel aluminide, preferably nickel aluminide powder, in particular Ni ₃ Al powder.

For example, nickel aluminide powder (Ni-Al powder), e.g. Ni-13Al powder with an aluminum content of about 13.3 wt. %, can be used. The particle size of the Ni-Al powder is FSSS<70 μm, preferably FSSS<45 μm. To adjust and balance the targeted carbon content, W metal powder (FSSS<2 μm) and carbon black are used. To alloy the binder phase with alloying elements such as Ti, Ta, Mo, Nb, V, Cr, their carbide powders or W-containing mixed carbides with particle size<3 μm are used.

The powder mixture is produced according to the prior art by wet grinding, preferably in a ball mill equipped with hard alloy balls. As grinding media, ethanol and hexane are used. Other possible grinding media would be acetone, or an aqueous medium with suitable inhibitors.

For the production of powder mixtures of cemented carbide materials with binder contents >15%, a single grinding step is sufficient due to the high binder content and enhanced recrystallization. On the other hand, for binder contents below 15%, a multi-stage wet grinding process is advantageous to effectively grind the Ni-Al powder and minimize oxide formation during the grinding process.

In the first method step, the Ni-Al powder is intensively mixed with the grinding liquid and coarse-grained tungsten carbide with an average grain size FSSS>20 μm, preferably 30-60 μm. If necessary, pressing aids, small amounts of alloying components and cobalt powder can already be added here.

The grinding parameters (time, ratio of grinding balls to grinding material, grinding media) and the ratio of WC to Ni-Al powder vary depending on the WC grain size set in the cemented carbide material.

In the second process step, 50-80% by weight of WC raw material of a given particle size is added to the pre-grind and mixed, with the main focus on reducing agglomerates and obtaining as homogeneous a mixture as possible.

If alloy adjustment and addition of pressing aid have not yet been carried out in the first grinding step (pre-grinding VM), this can also be carried out in the second method step.

The slurry obtained by wet milling is dried by conventional techniques and converted into a powder suitable for pressing. Preferably, this is done by spray drying.

The forming is preferably carried out directly by axial pressing in a mechanical, hydraulic or electromechanical press.

Sintering is carried out in a vacuum at 1350-1550 °C, preferably in an industrial sintering HIP furnace, in which an overpressure is generated after liquid phase sintering by introducing an inert gas, in which any residual porosity present can be removed.

FIG. 4 shows an exemplary phase diagram of WC-Co-Ni _{3 Al for 3 wt. % Co and 3 wt. % Ni 3 Al} , which illustrates the formation of these precipitates.

After solidification of the melt, initially only WC and a solid solution of Co, Ni, Al, W and C are present. Only below the solvus temperature do intermetallic phase materials precipitate out of the solid solution, with the intermetallic phase materials being formed according to the structural formula (M,Y) ₃ (Al,X), where M=Ni, Y=Co and/or other elements, and X=tungsten and/or other elements. These intermetallic phase materials can be visible with a scanning electron microscope.

In Fig. 2 and Fig. 3, two different cemented carbide materials according to the invention in the form of hard alloys are shown by such scanning electron microscope images. The binder phase of such hard alloys can be clearly seen, and the intermetallic phase material (light phase) 10 and the metallic binder material 30 (dark phase) can be distinguished. The WC grains 20 are bonded to each other via the binder phase.

It has been found that the intermetallic phase material is uniformly distributed in the binder phase, with the crystals of the intermetallic phase material having a cubic shape and preferably smaller than 1500 nm. The crystals of the intermetallic phase material (M,Y) ₃ (Al,X) have the crystal structure L1 ₂ (space group 221) according to ICSD (Inorganic Crystal Structure Database).

Such sintered carbide materials can be bonded to a steel substrate, whereby they form the working elements of tools, for example grinding tools, soil excavation tools, preferably for road mills, recyclers, stabilizers, agricultural or forestry soil excavators. In this case, the working elements are arranged in the working area of the tool. The joining to the substrate is carried out by brazing, in particular by hard brazing. In this case, heat is applied to the tool in order to form the brazed joint. The tool is then quenched, for example in a water and oil emulsion.

Due to the at least partial re-melting of the intermetallic phase material during the brazing process, the components of the intermetallic phase material are present as dissolved components in the cemented carbide material after quench hardening. Thus, in this way, a cemented carbide material precursor is formed.

This sintered carbide material precursor is then subjected to a heat treatment, as already mentioned above on several occasions. Here, heat can be introduced into the sintered carbide material by the heat treatment, the temperature being preferably below the solvus temperature and preferably above 400° C. The treatment time, i.e. the time during which the heat treatment is carried out, is 0.25 to 24 hours. Intermetallic phase material is again formed in at least some regions of the sintered carbide material during the heat treatment, resulting in a strengthening of the binder phase.

Here, the heat treatment can be an active process in which heat is introduced into the sintered carbide material in a targeted manner using a heat source. Advantageously, the heat treatment is carried out passively, for example when the sintered carbide material precursor comes into contact with a workpiece, for example a road pavement, during use of the tool. During this contact, heat is introduced into the sintered carbide material precursor, bringing it to a temperature at which an intermetallic phase material is formed. In this way, the tool is automatically strengthened according to the invention, with the sintered carbide material according to the invention being formed in the areas subject to wear.

It is also conceivable that the cemented carbide material configured as described above is produced by a sintering process in which the intermetallic phase is formed. The product can then be advantageously brought to a temperature above the solvus temperature at which the intermetallic phase at least partially re-dissolves. The material is then quenched to form a cemented carbide material precursor. The cemented carbide material precursor is then subjected to a heat treatment to form the cemented carbide material according to the invention.

In order to facilitate precipitation of the intermetallic phase material in the binder phase, it can be advantageously provided that the (M,Y) ₃ (Al,X) content in the binder phase is ≧40%, with the carbon proportion being adjusted to be stoichiometric or less than stoichiometric.

It was found that the higher the solubility of tungsten in the binder phase, the more stable the precipitation of the intermetallic phase was, because the crystal structure of the intermetallic phase material was filled with " _Co3W ", shifting the precipitation range to higher temperatures.

Here the elements Mo, Nb, Cr, V and especially Ti, Ta show a similar effect and can be added to the alloy in small amounts (<15 at% in the binder).

The amount of alloy that can be used depends on the solubility product of each of the metal carbides. Even if the magnitude is negligible, a surprisingly clear effect has been shown that is not due to a grain refinement effect.

Based on the improved stability and precipitation behavior, the addition of other elements to the alloy allows the proportion of intermetallic phase material in the binder to be reduced, even below 40%. Furthermore, in the presence of, for example, Ti and Ta, these elements take over the role of tungsten as a stabilizer, making it unnecessary to adjust the carbon proportion below the stoichiometric amount.

The effect of the precipitation of intermetallic phase materials on hot strength can be impressively demonstrated by hot compression tests. Figure 1 shows the hot compression strength of a hard alloy containing 6% binder at different test temperatures and strain rates. In particular, at a test temperature of 800°C, the intermetallic phase materials increase the strength by about 40-50%.

Physical property values are measured on samples of the sintered carbide material according to the invention, which serve to characterize the material and its properties.

For hard alloys, the measurement of the coercive force H _cM and the specific saturation magnetization 4 ps is an established non-destructive testing method.

Both measurement parameters are also determined for the characterization of the sintered carbide material according to the invention using Förster Koerzimat® 1.097.

Another parameter for the characterization of materials is density, which is determined by weighing according to Archimedes' principle.

The hardness of the material is determined for hard alloys on metallographically prepared polished samples in accordance with valid standards. Advantageously, the Vickers HV10 hardness test with a test load of 10 kp is used (ISO 3878).

Similarly, the porosity of the sintered material (DIN ISO 4499-4) and the aluminium oxide particles are checked and evaluated by optical microscopy on the polished samples. Comparative images of the A or B porosity can be used to estimate the volume percentage of aluminium oxide in the microstructure, where A08 or B08 corresponds to a volume percentage of approximately 0.6% by volume. For optical microscopy examination of the eta phase, this is etched with Murakami's reagent according to the standard (DIN ISO 4499-4). The average grain size of the WC is determined according to DIN ISO 4499-2. In this case, the SEM (scanning electron microscope) images are evaluated by the section method.

The proportion of intermetallic phases in the binder and the maximum size of the precipitated particles are also determined by SEM imaging, but using an in-lens BSE detector. For this purpose, images are taken at several points on the sample and an evaluation is carried out on representative cross sections by image processing and determination of the area proportion by gray-scale division.

Example:
The following table shows examples of sintered carbonized bodies according to the invention, which in principle can be produced by the same methods as described above:

Thus, in accordance with the above explanation, the present invention relates to a cemented carbide material, in particular a hard alloy, comprising 70-95% by weight, advantageously 80-95% by weight, of tungsten carbide in dispersed form and a binder phase, the binder phase comprising a metallic binder material which comprises Co, dissolved Ni and dissolved Al, the binder phase optionally comprising an intermetallic phase material, if present, being formed according to the structural formula (M,Y) ₃ (Al,X), where M=Ni, Y=Co and/or other components, X=tungsten and/or other components, the binder phase having the following chemical element composition: Ni>25% by weight, Al>4% by weight, balance Co and dissolved binder components, e.g. W and/or C.

The present invention will now be described in more detail with reference to the embodiment examples shown in Figures 5 to 12.

FIG. 1 is a diagram showing the hot compressive strength of a hard alloy containing 6% binder when the test temperature and strain rate are changed. FIG. 1 shows a cemented carbide material according to the invention in the form of a hard alloy. FIG. 1 shows a cemented carbide material according to the invention in the form of a hard alloy. 1 is a phase diagram of WC-Co-Ni ₃ Al for 3 wt % Co and 3 wt % Ni ₃ Al. FIG. 2 is a vertical cross-sectional view of the chisel tip. 6 shows the chisel tip according to FIG. 5 along the section line marked VI-VI in FIG. 5. 7 is a vertical cross-sectional view of a chisel tip 50 according to FIGS. 5-6, but with a modified texture. 7 is a vertical cross-sectional view of a chisel tip 50 according to FIGS. 5-6, but with a modified texture. 7 is a vertical cross-sectional view of a chisel tip 50 according to FIGS. 5-6, but with a modified texture. 7 is a vertical cross-sectional view of a chisel tip 50 according to FIGS. 5-6, but with a modified texture. 7 is a vertical cross-sectional view of a chisel tip 50 according to FIGS. 5-6, but with a modified texture. 7 is a vertical cross-sectional view of a chisel tip 50 according to FIGS. 5-6, but with a modified texture.

Figures 5 to 12 show chisel tips 50 in the form of a sintered carbide material 40. These chisel tips 50 are preferably used in cutting tools, in particular cutting chisels, round shank chisels, road cutting chisels, mining chisels, etc.

The chisel tip 50 is formed and prepared to be joined, preferably brazed, to the steel body. For this purpose, the steel body usually has a head with an integrally molded shank, preferably a round shank. Towards the direction away from the shank, the head has a receiving portion for the chisel tip 50. The chisel tip 50 can be fixed in or on this receiving portion.

The chisel tip 50 is constructed in one piece and has a base part 51. The base part 51 allows the chisel tip 50 to be joined to the steel body. Advantageously, the base part 51 has a joining surface 51.1. In order to fix the chisel tip 50 to the steel body, a brazing material for a hard brazing joint can be arranged between the joining surface 51.1 and the steel body.

In order to keep the thickness of the brazing gap between the chisel tip 50 and the steel body as constant as possible, it is possible to provide that a spacer 51.2 is molded on the chisel tip 50 in the region of the joining surface 51.1, which spacer 51.2 is formed so as to protrude beyond the joining surface 51.1 and rest on the steel body in such a way that the joining surface 51.1 is held at a distance from the mating surface of the steel body for the brazing process.

It can furthermore be provided that one or more recesses 52 are present in the region of the joining surface 51.1. Here, the recesses can advantageously be configured to transition via a convex radius of the joining surface 51.1 into a concave portion, which is preferably formed as a concave valley. The recesses 52 make it possible to reduce the material required for the chisel tip 50. Furthermore, the recesses 52 form a reservoir for excess brazing material in the region of the joining surface 51.1.

The base portion 51 preferably has a circumferential edge 51.3, which may be formed as a convex shape at least in some areas, and the edge 51.3 may be formed as a transition between the base portion 51 and the transition portion 53.

The transition 53 has a first region formed as a concave region 53.1. A truncated cone-shaped geometry or a combination of the concave region 53.1 and an at least partially truncated cone-shaped geometry may also be provided. The first region causes the chisel tip 50 to taper in a direction from the base 51 towards the tip 54 of the chisel tip 50.

The transition 53 may further have a cylindrical region 53.2 that continues to the first region on the opposite side of the base 51. As shown in FIG. 5, at least in the region of the chisel tip 50, the transition between the first region and the cylindrical region is formed so as to be continuously differentiable, preferably continuously in the direction of the longitudinal central axis of the chisel tip, so that a break in continuity is avoided.

The chisel tip 50 can advantageously have recesses 53.3 in the region of the transition 53. These serve to reduce material and optimize the evacuation of the soil material that is undermined during use of the tool.

The chisel tip 50 has a tip 54 which is preferably continuous with a cylindrical region 53.2 at the transition 53. The connection 54.1 may be formed as a convex bend. The connection 54.1 is followed by a tapered section 54.2 which transitions into an end section 54.3. The end section 54.3 is preferably formed in the form of a convex bend, particularly preferably in the form of a spherical cap.

Figure 6 shows a cross-sectional view and a top view of the chisel tip 50. The indentations 53.3, which are evenly distributed around the circumference of the chisel tip 50, are clearly visible.

The chisel tip 50 has at least one first volume region 70 and at least one second volume region 60.

In the second volume region 60, the cemented carbide material 40 comprises tungsten carbide grains (WC grains) 20 which are bonded to one another via the metal binder material 30 of the bonding phase. It is possible to provide that in the second volume region 60, no intermetallic phase material 10 is present or that the intermetallic phase material 10 is present in a concentration of preferably less than 30% by weight/unit volume, preferably less than 25% by weight/unit volume, particularly preferably less than 15% by weight/unit volume.

In the first volume region 70, the cemented carbide material 40 comprises tungsten carbide grains (WC grains) 20 bonded together via the metal binder material 30 of the bonding phase. In the first volume region 70, the intermetallic phase material 10 is preferably present in a concentration of ≧30% by weight/unit volume, preferably in the range of 30-70% by weight/unit volume, particularly preferably in the range of 35-60% by weight/unit volume, and even more preferably in the range of 40-50% by weight/unit volume.

The intermetallic phase material is formed according to the structural formula (M,Y) ₃ (Al,X), where M=Ni, Y=Co and/or other elements, and X=Tungsten and/or other elements. The binder phase has the following chemical element composition by weight: Ni>25%, Al>4%, balance Co and dissolved binder elements such as W and/or C.

The relative proportion Y of intermetallic phase material per unit volume in the first volume region 70 is greater than that in the second volume region 60.

As shown in the drawings, the first volume region 70 and the second volume region 60 can border at least one region of the surface of the cemented carbide material 40.

The first volumetric region 70, which has a high relative proportion of intermetallic phase material 10, forms part of the surface of the chisel tip 50, particularly the tip 54 of the tapered portion 54.2 and/or the terminal portion 54.3.

Starting from the tip 54, the first volume region 70 may further extend to the base 51. Advantageously, the region of the base 51 may also be formed by the second volume region 60. The second volume region 60, which has a relatively low content of intermetallic phase material 10 or is free of intermetallic phase material 10, is advantageously arranged in the region of the transition 53. As shown in FIG. 6, both the first volume region 70 and the second volume region 60 may be adjacent to the surface of the chisel tip 50 in the region of the transition 53.

The arrangement of the first and second volumetric regions 70, 60 according to the example embodiment shown in Figures 5 and 6 has the following technical advantages:
Due to the tip 54 being formed by the first volume region 70, in this case a high wear resistance is provided, particularly in the region at risk of wear.

In areas where there is no risk of wear, toughness and therefore fracture stability is improved.

Furthermore, it is contemplated that the volume regions 60, 70 allow for targeted adjustment of a particular wear pattern, for example to assist in resharpening the chisel tip.

Figures 7 to 12 show further configuration variants of the chisel tip 50. Here, the chisel tip 50 is configured essentially identically to the chisel tip 50 shown in Figures 5 and 6. In this respect, reference is made to the above description in order to avoid repetition, and only the differences are described. The chisel tip 50 according to Figures 7 to 12 differs in particular in the arrangement and configuration of the first and second volume regions 70, 60.

As shown in FIG. 7, the first volume area 70 is arranged in the area of the tip 54, preferably partially in the cylindrical area 53.2 of the transition 53. However, it is also conceivable that the first volume area 70 is arranged only in the area of the tip 54. In this case, the first volume area 70 effectively protects against wear and tear in the area of the tip 54.

8 shows that the first volume region 70 may be arranged completely or partially within the chisel tip 50. The first volume region 70 may in this case be configured to extend in the direction of the longitudinal central axis of the chisel tip 50, preferably over the entire area of the transition 53. As already mentioned above, the first volume region 70 is characterized by a particularly high shear strength due to the presence of the intermetallic phase material 10. In this way, the transition region 53, which is at risk of fracture, is effectively reinforced by the first volume region 70.

Conversely, as shown in FIG. 9, it is also possible to arrange the second volume area 60 completely or partially within the chisel tip 50. In this case, the first volume area 70 advantageously completely surrounds the second volume area 60. To prevent wear and tear particularly effectively and to prevent breakage of the chisel tip 50 in the region of the transition 53, the first volume area completely or almost completely forms the surface of the chisel tip 50. Due to its large extent transverse to the longitudinal central axis of the chisel tip 50, the first volume area 70 also provides a high axial bending resistance moment.

Figure 10 shows that in a development of the variant according to Figure 7, the first volume region 70 with a high relative proportion of intermetallic phase material 10 can extend beyond the region of the tip 54 and the transition 53, whereby the first volume region 70 may advantageously completely form the surfaces of the tip 54 and the transition. The first volume region 70 can also be led up to or into the base 51, as shown in Figure 10.

Figure 11 shows that the first volume region 70 may extend within the chisel tip 50 from the distal end 54.3 of the tip 54 to the base 51 to form a continuous volume region.

Figure 12 shows that it is also possible to provide an embodiment in which the first and second volume regions 70, 60 are interchanged, contrary to the embodiment shown in Figures 5 and 6.

The above figures should be understood as schematic illustrations. This is especially the case when the first or second volume regions 70, 60 are not precisely sharply defined as shown in the figures, but rather a transition is formed between the two volume regions 70, 60.

In order to produce or form the cemented carbide material 40 shown in Figures 5 to 12, according to the invention, a method is first used in a first method step to produce a cemented carbide material precursor, in particular a hard alloy, which comprises 70-95 wt. %, preferably 80-95 wt. %, of tungsten carbide in dispersed form and a binder phase. The binder phase comprises at least Co as metallic binder material and also the dissolved elements Ni and Al. The binder phase has the following chemical element composition:
Ni>25wt%, Al>4wt%, balance Co and dissolved binder components such as W and/or C
has.

In a further method step, the cemented carbide material precursor is subjected to a heat treatment as described above to form a cemented carbide material 40 including intermetallic phase material 10 in a binder phase in at least a first volumetric region 70. The intermetallic phase material 10 is formed according to the structural formula (M,Y) ₃ (Al,X), where M=Ni, Y=Co and/or other elements, and X=tungsten and/or other elements.

For this purpose, in a further method step, the sintered carbide material precursor can be held during heat treatment in a temperature range between 400° C. and the solvus temperature for a period of 0.25 to 24 hours. To form the individual volume regions 60, 70 in a targeted manner, the targeted heating can be carried out, for example, by means of a laser or an induction coil.

Claims (22)

A method for producing a sintered carbide body, comprising mixing tungsten carbide powder (WC powder) with a metal binder material comprising cobalt powder (Co), nickel (Ni) and aluminium (Al) in a mixing and/or grinding process, preferably a wet grinding process, to form a powder mixture, producing a green compact from at least a part of said powder mixture by pressing, sintering said green compact under the action of temperature and pressure in a sintering step, and forming a sintered carbide body after a cooling step following said sintering step,
A method, characterized in that nickel aluminide, preferably nickel aluminide powder, in particular Ni ₃ Al powder, is added as intermetallic phase material to said mixing and/or grinding process. 2. The method of claim 1, wherein a cooling step and/or a heat treatment subsequent to the sintering step of the sintered carbide body is controlled to form an intermetallic phase material in the binder phase of the sintered carbide body, at least a portion of which is preferably formed according to the structural formula (M,Y) ₃ (Al,X), where M=Ni, Y=Co and/or other elements, and X=Tungsten and/or other elements. The method according to claim 1 or 2, wherein during the cooling step, the sintered body is held at a temperature range between 400°C and the solvus temperature of the sintered body for 0.25 to 24 hours. The method according to any one of claims 1 to 3, wherein the green compact comprises 70-95 wt%, preferably 80-95 wt%, tungsten carbide (WC), 1-28 wt%, preferably 1-19 wt%, cobalt (Co), and 1-28 wt%, preferably 1.5-19 wt%, nickel aluminide as intermetallic phase materials. The method according to any one of claims 1 to 4, wherein the nickel aluminide, preferably the _Ni3Al powder, is produced by a molten metallurgical process and/or the nickel aluminide, preferably the _Ni3Al powder, as a material produced by a molten metallurgical process is added to the mixing or grinding process. The method according to any one of the preceding claims, wherein the nickel aluminide, preferably _Ni3Al powder, having an average particle size FSSS<70 μm, preferably having a particle size FSSS<45 μm, is added to the grinding and/or mixing process. 7. The method according to claim 1, wherein in a preparation step, preferably a first grinding step, the nickel aluminide is mixed with a grinding liquid and coarse-grained tungsten carbide, preferably having an average particle size FSSS of >20 μm, particularly preferably having an average particle size FSSS in the range of 30 to 60 μm, for example in the form of macrocrystalline and/or monocrystalline tungsten carbide, to form crushed nickel aluminide, preferably crushed nickel aluminide powder, in particular crushed Ni ₃ Al powder, from the nickel aluminide. The method of claim 7, wherein a pressing aid, at least one alloying component and/or cobalt powder is added and mixed with the nickel aluminide and/or the crushed nickel aluminide during the preparation step and/or the subsequent grinding step. The method according to claim 7 or 8, wherein the proportion of nickel aluminide in the ground mixture of the preparation step is in the range of 8 to 50% by weight, preferably in the range of 9 to 25% by weight. The method according to any one of claims 7 to 9, wherein WC powder is added to the pre-ground product obtained in the preparation step in a subsequent grinding step so that the proportion of WC powder in the resulting ground mixture is in the range of 70 to 95% by weight, and the crushed nickel aluminide is crushed in the subsequent grinding step to obtain finely crushed nickel aluminide. The method according to any one of claims 1 to 10, wherein the green compact is sintered in a furnace by a liquid phase sintering process at a sintering temperature in the range of 1350°C to 1550°C. 12. The method of any one of claims 1 to 11, wherein during the liquid phase sintering process at the sintering temperature, the cobalt and the intermetallic phase are at least partially dissolved together in the melt, and upon cooling and/or heat treatment, the intermetallic phase material is formed in the binder phase, the intermetallic phase material being formed according to the structural formula (M,Y) ₃ (Al,X), where M=Ni, Y=Co and/or other elements, and X=tungsten and/or other elements. The method according to any one of claims 1 to 12, wherein at least a portion of the intermetallic phase material in the binder phase has a maximum size (measured by sectioning on microscopic specimens) of 1500 nm, preferably 1000 nm. 14. The method of claim 1, wherein at least a portion of the crystals of the intermetallic phase material (M,Y) ₃ (Al,X) in the binder phase have the crystal structure L12 (space group 221) according to ICSD (Inorganic Crystal Structure Database). 15. The method according to claim 1, wherein in the grinding and/or mixing process, atomized nickel aluminide, preferably atomized _Ni3Al powder, is processed as intermetallic phase material. The method according to any one of claims 1 to 15, wherein the mixing and/or grinding process is a multi-stage process having at least two mixing and/or grinding steps, and preferably the nickel aluminide is added before the last grinding and/or mixing step. The method according to any one of claims 1 to 16, wherein the mixing and/or grinding process involves processing Nb and/or Ti and/or Ta and/or Mo and/or V and/or Cr, and the binder phase of the sintered cemented carbide material, in particular the metallic binder material (30) and/or the intermetallic phase material (10), comprises Nb and/or Ti and/or Ta and/or Mo and/or V and/or Cr, advantageously one or more of said materials being present in dissolved form in the binder phase. The method according to any one of claims 1 to 17, wherein the carbon proportion in the cemented carbide material (40) is adjusted to be stoichiometric or less than stoichiometric. The carbon percentage in the cemented carbide material (40) is
C _{stoichiometric amount} (wt%)-0.003 x binder content (wt%) to C _{stoichiometric amount} (wt%)-0.012 x binder content (wt%), advantageously
20. The method of claim 18, wherein the binder content ranges from C _{stoichiometry} (wt%) - 0.005 x binder content (wt%) to C _{stoichiometry} (wt%) - 0.01 x binder content (wt%). The method according to any one of claims 1 to 19, wherein the proportion of Mo and/or Nb and/or Ti and/or Ta and/or Cr and/or V in the binder phase is ≦15 at%. The method according to any one of claims 1 to 20, wherein the dispersed tungsten carbide (WC grains 20) is present in the cemented carbide material (40) in the form of grains having an average grain size, measured according to DIN ISO 4499 Part 2, in the range of 1 to 15 μm, preferably in the range of 1.3 to 10 μm, particularly preferably in the range of 2.5 to 6 μm. The method according to any one of claims 1 to 21, wherein the maximum proportion of Fe in the binder phase of the cemented carbide material is <5 wt. % and/or other unavoidable impurities are present in the binder phase.

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