DE2856513A1 - HARD ALLOY CONTAINS MOLYBDAEN AND TUNGSTEN - Google Patents

Description

Hard alloy containing molybdenum and tungsten

The invention relates to a hard alloy containing molybdenum and tungsten. The invention also shows a method for producing this hard alloy.

So far, cemented carbide was used as the starting material Tungsten carbide powder (WC) used as the predominant component, together with a suitable binder metal, which is typically an iron group metal, depending on the requirements of the to be produced alloy carbides or carbonitrides of metals with high melting point such as titanium (Ti), tantalum (Ta), niobium (Nb), molybdenum (Mo), Hafnium (Hf), vanadium (V) and chromium (Cr) can be added. It is true, however, that tungsten is a relatively expensive metal and that it is only in occurs in a few parts of the world. Accordingly, it is considered a so-called "strategic" material and its availability may be subject to political considerations. An increase in the demand for cemented carbide containing tungsten carbide thus becomes inevitable Wise the problem of natural supplies ^ if there is it is possible to replace the tungsten carbide with the carbide of another refractory metal ^ so such an exchange occurs of great importance in the industry.

Molybdenum monocarbide (MoC) is considered a useful substitute material because only this carbide has the same crystal structure of simple hexagonal type as tungsten carbide and it also has similar mechanical ones Has properties like tungsten carbide. However, the existence of hexagonal molybdenum monocarbide as a simple substance continues to this day doubted and attempts to stabilize the molybdenum monocarbide, have so far only been through the formation of a solid solution with Tungsten carbide. This procedure was first

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described by W. Dawihl in 1950, but this solid solution not examined in detail and their economic value was not recognized at the time.

Recently, however, as a result of the rise in the price of tungsten Investigations into the use of the solid solution (Mo W) C, where-

x y

rin χ + y = 1. It's an interesting question as to why Studies of this solid solution and attempts to use it have not previously been actively carried out.

Molybdenum carbide is stabilized as a monocarbide with a crystal structure of the simple, hexagonal type when a solid solution is formed with tungsten carbide. If this stable carbide (Mo, W) C can be easily produced, replacing tungsten with molybdenum appears to be possible. For this purpose, a method for stably producing (Mo, W) C has been proposed (Japanese Patent Application OPI No. 146303/1976; U.S. Patent 4,049,380). If the (Mo, W) C powder obtained by this method is used as the starting material for the production of a (Mo, W) C-Co alloy as a substitute for WC, it turns out, however, that the MoC in the alloy is not stable, and that Mo "C tends to be precipitated frequently. This method has also not been used in practice because it requires longer heat treatments.

Furthermore, the production of a molybdenum-tungsten carbonitride with the crystalline structure of tungsten carbide has been proposed, for which purpose molybdenum and tungsten in a combined form and carbon in a ratio sufficient to form the monocarbide are heated in a nitrogen-containing atmosphere ( Japanese Patent Application (OPI) No. 104617/1978). The aim of this process, however, is to stabilize the alloy by incorporating nitrogen, so that no (Mo, W) ₂ C is precipitated.

⁹¹⁰³ §09827 / 0991

The object of the invention is accordingly to provide a hard alloy containing molybdenum to provide, which corresponds to a cemented carbide that mainly consists of tungsten carbide (WC), part of which has been replaced by molybdenum carbide (MoC).

Another object of the invention is to provide a hard phase alloy composed of a hexagonal monocarbide of (Mo, W) C in which (Mo, W) ₂ C is dispersed.

Another object of the invention is to provide a hard alloy with a Provide hard phase, which consists of a compound of tungsten and molybdenum with carbon, nitrogen and oxygen, the one Has a crystal structure of a simple hexagonal type.

The object is achieved according to the invention by a hard alloy, consisting of composed of a hard phase composed of a compound of (Mo, W) C of simple hexagonal type and a binding phase composed of at least one Element that is selected from the group consisting of iron, cobalt, Nickel and chromium, in which a compound represented by (Mo, W) «C and has a crystal structure of a hexagonal type, is uniformly dispersed as a hard phase.

According to the invention, a hard alloy is provided, consisting of a hard phase of at least one compound of the simple, hexagonal type MC, where M is metal and C is carbon, selected from the group consisting of mixed arbides, carbonitrides and carboxynitrides of molybdenum and tungsten as the predominant component, and a binder phase of at least one element selected from the group consisting of iron, cobalt, nickel and Chromium, in which a hard phase consists of a compound of the type M «C is uniformly dispersed with a hexagonal crystal structure. the

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Amount of the compound of the type M "C, which is eliminated, is preferably 30 vol -% or less, particularly 5 to 25 vol -.%, Because at more than 30 vol -.% Forms the M ₀ C large particles, so that the effect of an alloy strengthened by dispersion is not achieved. As an essential condition for uniform dispersion of the hard phase of the M ₀ C type, it is necessary that the atomic ratio of the carbon content of the hard phase to the theoretical carbon content of the hard phase of the MC type is 0.98 to 0.80.

Briefly summarized and according to a preferred embodiment of the invention, a hard alloy is provided, consisting of one or more carbide phases consisting of 80 wt .-% or more of a carbide of the type MC in solid solution, containing molybdenum and tungsten and having a crystal structure of simple , hexagonal type, and 20 wt .-% or less of a mixed carbide of the type M _{0 C, which contains Mo 0} C as the main component and is of a granular or spherical shape with a particle size of 10 μπα or less, the carbide of the type M ₀ C is dispersed in the alloy, as well as from 3-50 wt. % Of a binder or binder phase, which consists of a metal of the iron group.

The invention is explained in more detail with the aid of the figures. Show it;

1 shows a microscopic image enlarged 200 times

a molybdenum-containing cemented carbide alloy according to the prior art, in which the occurrence of an acicular precipitated carbide of the type (Mo, W) ₂ C is visible;

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2 shows a micrograph, enlarged 200 times, of a (Mo, W) C- (Mo, W) «C -C o alloy according to the invention, in which a carbide of the type (Mo, W) ₂ C is uniformly dispersed;

Fig. 3 shows an X-ray diffraction pattern of one according to the invention

Alloy;

4 shows a graph comparing the high temperature strength of a WC-Co alloy according to the prior art with a (Mo, W, Cr) C-Co alloy according to the invention, in which WC-IO% Co is replaced by A, (Mo, W, Cr) - (9% Co + 5% Ni) is represented by B, WC-15% Co by C and (Mo, W, Cr) -IS % Co by D;

5 shows a graph of the compressive strength and tensile strength of a WC-Co alloy as a function of the load, in which WC-5% Co through A, WC-IO% Co through B, WC-25% Co through C, WC-30 % Co is represented by D, WC-7% Co by E, WC-12% Co by F, and WC-15 % Co by G, the Co percentages being by weight and

6 shows a graphical illustration of the comparison of the pressure deformation as a function of the load in the case of a (Mo, W) C-Co alloy according to the invention and a WC-Co alloy according to the prior art, in which WC-II% Co by H, WC-16% Co represented by I, (Mo _Q? W _{0 3} ) C-Il% Co represented by J, (Mo _{0 5} W _{Q &} ) C-19% Co represented by K and WC-24% Co represented by L, where the amount of the binder metals Co and Ni are given as percentages by volume.

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In various investigations of mainly the relationship between the carbon content and the strength of a silicon carbide alloy which have been made from (Mo W) C as a starting material

χ y

the following facts emerged:

If the carbon content of the alloy is less than the theoretical carbon content, a molybdenum- _{tungsten mixed carbide of the type M 0} C is precipitated in the form of needle-shaped crystals, as is the case in the prior art (FIG. 1). If, on the other hand, the alloy contains a micro-quantity of an impurity element, the mixed material does not consist of needle-shaped crystals, but is precipitated in fine-grained form (FIG. 2). As a result of measurements of the strengths of alloys containing acicular precipitated molybdenum-tungsten composite carbide of the type M ₀ C and of alloys containing granular composite carbide of the type M ₀ C, it was found that the latter compared to the former of superior strength. It is known that the strength of such an alloy depends in large part on differences in the precipitated forms just described. In the first case, stress is concentrated on the needle-shaped crystals of molybdenum-tungsten mixed carbides (M ₀ C, M ₃ C ₀ ), so that fractures are formed and the strength of the alloy is reduced, while in the second case an alloy of a so-called dispersed type present, in the granular molybdenum tungsten carbides of the type M ₀ C are evenly or widely distributed, so that a concentration of the load acting on the mixed carbides is avoided and an external force acting on the alloy is absorbed, which is an increase the strength of the alloy means.

The reason why the granular crystals of a molybdenum-tungsten mixed carbide of the type M ₀ C are precipitated is not known in detail, but can be considered as follows:

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With a simple (Mo, W) C alloy, needle-shaped crystals of a molybdenum-tungsten mixed carbide of the type M ₀ C are precipitated because when the liquid phase cools, in which tungsten, molybdenum and carbon are dissolved in the binder phase, the precipitation temperatures of tungsten and molybdenum are different. That is, WC is precipitated at a relatively high temperature and only Mo and C remain in the binder phase until the end. At a temperature below 118O ⁰ C, MoC is _{decomposed to Mo 0} C and C and consequently Mo ₀ C remains as an agglomerate. Free carbon and Mo ₀ C, after precipitation, can be well dispersed by a rapid cooling treatment, so that agglomeration is avoided. Such a procedure can be used with a sufficient cooling effect on an alloy sample of small size and with small heat capacity. However, it is difficult to treat a large sample of a cemented carbide alloy of high heat capacity by this procedure.

As a result of an X-ray diffraction analysis of the binder phase of an alloy in which acicular Mo ₀ C is precipitated, it was found that the lattice constant of the binder phase is not different from that of the pure metal and that the binder phase is neither alloyed nor brittle. Consequently, it is assumed that an alloy of sufficient strength can be produced _{by dispersing acicularly precipitated Mo 0 C in granular or spherical form.} When microscopic amounts of impurity elements are in the binder phase, Mo, W and C on seeds of these elements combine to form or precipitate a number of seeds of molybdenum-tungsten composite carbide of the MgC type before the liquid phase disappears or solidifies and Mo and C do not occur in the binding phase. Consequently, there is no precipitation of needle-shaped crystals of the type M ₀ C, not even at a temperature of 118O ⁰ C or less at which the

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liquid phase disappears. In general, a large number of MqC germs are excreted in a spherical or rod-like form. In order to disperse the molybdenum-tungsten mixed carbide of the type M ₀ C in the alloy and to precipitate it in a finer form, it is effective _{to inhibit the precipitation and the growth of the molybdenum-tungsten mixed carbide of the type M 0} C in that rapid cooling from the sintering temperature to the solidification temperature of the liquid * phase is carried out.

In order to disperse the M "C phase evenly in the alloy, a method is provided in which a previously prepared carbide (Mo, W) pC is added to the starting powders to be mixed and the sintering conditions are controlled in such a way that a (Mo, W) _o C-phase is excreted evenly. Further, a method is provided in which the generation of not only a complete solid solution of (Mo, W) C is performed during the step for producing a carbide, but also a carbide in the surface layer of fine (Mo, W) ₀ C is dispersed, a metal of the iron group such as Co, Ni or Fe is added to the carbide and the mixture is sintered _{with precipitation of (Mo, W) 0 C.} Finally, a method is provided in which Mo and W are added to a carbide of the type (Mo, W) C and thus the Mo and W dissolved in the binding phase is precipitated as (Mo, W) ₀ C during the sintering step.

Investigations of the conditions for the dispersion of (Mo, W) ₀ C in the alloy have shown that the addition of a micro amount of one or more impurity elements to the alloy leads to the M ₀ C phase around the impurity nuclei during sintering and cooling

is excreted around, whereby the M ₀ C phase is uniformly dispersed in a spherical shape. Especially during the formation of the carbide, an impurity element can be added and uniformly

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be dispersed. Impurities, such as iron, act to promote the carburization reaction and the resulting Fe "C" forms the nuclei for the dispersion of (Mo, W) -C. If the impurity is not added, there is a tendency for a needle-shaped M ₀ C phase to precipitate as the primary crystal. In order to prevent such precipitation, it is necessary to control the amount of Mo and W dissolved during sintering. For this purpose, the amount of W dissolved in the binding metal is increased beyond the amount of Mo, whereby (Mo, W) ₀ C is evenly precipitated.

Examples of the impurity elements to be added to the binding metal are beryllium, magnesium, calcium, boron, silicon, phosphorus, manganese, iron and rhenium, whereby one or more of these elements can be added % By weight added because "" if more than 3% by weight is added, the molybdenum-tungsten mixed carbide of the M ₀ C phase becomes brittle and the strength is not increased. The addition of titanium, zirconium, hafnium, tantalum and niobium as elements for inhibiting the precipitation and growth of the molybdenum-tungsten composite carbide of the M ₀ C type is also effective in promoting the dispersion precipitation of the composite carbide.

The carbides or mixed carbides of the type M ₀ C mentioned in this description include not only (Mo, W) ₀ C and (Mo, W) "C _O , but

Ct Ci O Ci

also other lower carbides containing other metals.

The particle size of the granular precipitated molybdenum-tungsten mixed carbide of the type M ₀ C is preferably 0.1-10 μm, and more effectively 1.0-2 μm, because too coarse precipitated particles reduce the strength and hardness of the alloy and if the particles are too small, the mixed carbide at the boundary of (Mo, W) C

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or the binder phase thereof is deposited so that the boundary strength is decreased and the strength of the alloy is decreased. When a carbide of the MC type is dispersed in the alloy, the amount of carbon is preferably 80-98 % of the theoretical amount that results if all carbides are considered to be carbides of the MC type. This corresponds to a presence of 2-30 % by volume.

A granular or spherical molybdenum-tungsten mixed carbide of the MgC phase has a great influence on the properties of the alloy, depending on the amount of mixed carbide. An X-ray diffraction analysis carried out with CuKa at 4C JW ₉ 80 mA, a sampling frequency of 4,000 Hz and a time constant of 0.2 seconds shows that the properties of the alloy are at least similar to those of WC-Co alloys when the ratio of the maximum in the X-ray diffraction diagram for M ₁₅ C to the maximum for MC, which appear at X-ray diffraction angles (2θ) close to 39.4 ° and 48.4 °, in the range from 0.01 to 0.5, in particular 0.05 to 0 , 20 lying. For purposes of explanation, M "C" molybdenum-tungsten mixed carbide is sometimes represented by M "C or Mo" C in this specification. Although W, Co, Ni, N and / or O are dissolved in M “C or Mo ₃ C and the ratio of the metallic components and non-metallic components varies at about 2: 1, the beneficial effects given by the invention are not lost.

As the binder metal is preferably a metal of the iron group at a ratio of 3-50 wt -.%, Based on the alloy composition used because below 3 wt -.%, The alloy is brittle and above 50 wt -.%, The Hoehtemperatereigenschaften poor will. The metal of the iron group used as the binding phase can naturally dissolve metals of groups IVa, Va and VIa,

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and it is possible to add other elements that are soluble in it such as aluminum, silicon, calcium, silver, etc., while maintaining the beneficial effects of the invention.

The basic concept of the invention remains, even if part of the molybdenum and tungsten carbide is replaced by a mixed carbide of the type B1, the titanium, zircon, hafnium, vanadium, niobium, tantalum, chromium, molybdenum and / or tungsten in a ratio of 30 wt -.% or less, preferably 0. 5 - 25 percent - is replaced.%. Further, a similar relationship also exists in the case of an alloy in which part of the C in the carbide is replaced with nitrogen and / or oxygen. Examples of the preferred embodiment in this case are as follows.

The first embodiment consists of the incorporation of N into (W, Mo) C, to give (W, Mo) (C, N), thus without prolonged heat treatment obtain a stable starting material of hexagonal WC type leaves.

The second embodiment consists of the incorporation of O into (W, Mo) (C, N) to give (W, Mo) (C, N, O) which is more stable.

The third embodiment consists of incorporating Cr in (W, Mo) (C, N) or (W, Mo) (C, N, O) to create (W, Mo, Cr) (C, N) or (W , Mo, Cr) (C, N, O), thereby obtaining a starting material of low weight and low price.

The fourth embodiment is that, in preparing these raw material powders, a mixture of oxides, metals, carbides and / or carbon is used in a part of the carburizing step of an atmosphere having a nitrogen partial pressure of 300 Torr or more

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at a temperature of 700 C or more is exposed to a to form stable starting powder.

The fifth embodiment is that when combining the the above-described starting powder containing an iron group metal, the presence of two or more kinds of hard phases of simple hexagonal toilet type, of different composition, is brought about in the finished alloy, whereby the alloy is given a high strength.

In these five embodiments, part of the phase of the type MC can also be replaced by a solid solution of the type B1, which contains one or more of the metals of groups IVa, Va and VIa and non-metallic elements, or the simple additives to cemented carbides, such as silver, silicon, bismuth, copper, aluminum, etc., also added to the binding metal from the iron group the advantages and effects of the invention will be apparent.

The embodiments described above are now intended in detail to be discribed:

In the important system of the invention, in which a simple, hexagonal Phase containing molybdenum and tungsten and an M “C phase occurs, was for the sintered alloy with a binding metal:

found that in the relationship

A _ / N atomic% . _n W atom% ,

^A " ^K (Mo + W) atom% ' ^{x {l} " (Mo + W) atom%'

the appropriate range of A is defined by 0.005 ^ A ^ 0, 5. If A is below the lower limit, the nitrogen has no effect, while above the upper limit it is difficult by sintering

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to achieve excellent properties. The most suitable range of A is defined by Q _P 01 = A = 0.4.

As for the effect of oxygen, it was found that in the relationship

R _ ( O atom% Y _n W atom% ,

^ö "* ■ (Mo + W) atom% ' ^{x K} (Mo + W) atom%'

the appropriate range of B is defined by 0.005 ^ B ^ 0.05 B below the lower limit, there is no beneficial effect of oxygen, while above the upper limit it is difficult to obtain excellent properties by sintering. The most suitable Range of B is defined by 0.01 = B = 0.04.

On the other hand, the ratio W / Mo is preferably 5/95 to 90/10 because below 5/95 the alloy is unstable and above 90/10 the advantages of replacing (light weight, low price) are essentially lost, those of replacing molybdenum or Tungsten amount of chromium used is 0.5 or less in atomic ratio to (W + Mo), because at values greater than 0.5 the alloy is brittle, although the corrosion resistance is increased.

As is known, it is advantageous for the production of cutting tools to use a solid solution of the type Bl au biidea _s which is composed of at least one of the metals of groups IVa, Va and VIa, such as titanium, zirconium, hafnium, vanadium, tantalum, chromium, Molybdenum and tungsten with at least one non-metallic component, such as carbon, nitrogen and oxygen, in addition to the simple hexagonal phase. The amount of the solid solution of the type B1 is advantageously selected in accordance with the cutting purpose.

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As for the amount of nitrogen in this case, the result was from research found that when changing the definition of A to

, N atom% . ₍₁ W atomic%

* Size IVa, Va, VIa metals atom-% ' ^K "Gr. IVa, Va, VIa metals atom-%'

the appropriate range of Α values is represented by the equation 0.005 = A = 0.5, although some of the nitrogen in the solid solution of the type B1 is included. The optimal range of A values is given by the equation 0.01 = A = O ₃ 4.

As far as the amount of oxygen is concerned, research has found that when changing the definition of B to

/ O atomic% . _n W atom%

¹ gr. IVa, Va, VIa metals atom-% ' ^l "Gr. IVa, Va, VIa metals atom-%'

the appropriate range of B-values is represented by the equations

nale B
given by the equation 0.01 ^ B = 0.04.

chung 0.005 = B = O ₅ Oi. The optimal range of B values is again-

A metal of the iron group in proportion is preferably used as the binding metal of 3-50% by weight, based on the total composition, is used because the alloy is at a ratio of less than 3% by weight Is brittle and, if the ratio is more than 50% by weight, the alloy is too soft.

To produce the starting material, a (Mo, W) powder is carburized with carbon, reduction and carburization of oxide powders with carbon or a combination thereof, the reaction is carried out at high temperature in a hydrogen reactor. As a result of studies of the nitrogen pressure Decomposition of (Mo, W) (C, N) it was found that the outer stick-

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substance pressure, depending on temperature, at a temperature of 700 ° C or more at which the carbonitride-forming reaction takes place should be 300 Torr or more. The simultaneous presence of Hydrogen is not always harmful, but it is desirable to adjust the amount of hydrogen so that it is at most twice in particular is at most the same as the amount of nitrogen, so that the nitride-forming reaction is not inhibited. Becomes a Ammonia decomposition gas is used, so it is necessary to use this with Enrich with nitrogen.

In the production of raw materials containing oxygen, the simultaneous presence of carbon monoxide and carbon dioxide in the atmosphere is required. In this case, the amount of hydrogen is not limited as described above, but should not exceed 50% of the atmosphere. The heating and sintering in an atmosphere containing nitrogen or carbon monoxide is effective to prevent denitrification or deoxidation of the sintered alloy.

In the five embodiments described above, the treatment for the dispersion of the phase of the type M «C can be omitted, in which case also considerable and good effects can be achieved are.

In the case of cemented carbide alloys, which mainly consist of WC, are for the purpose of use for drilling tools, bearings, thread cutting tools, etc. excellent properties can be achieved in that with a binder metal content of up to 15% by weight, the particle size of the carbide is decreased. However, if the alloy has a higher Contains binding metal, this measure has no effect. at

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Tools made from these alloys, especially in drills that are used for cutting at low speeds, are the cutting edges are deformed by frictional heat.

Further studies to develop an alloy which has great hardness against wear and strength have revealed that the strength against deformation at high temperatures can be remarkably improved by using a carbide composed of one in place of tungsten carbide solid solution of the three elements molybdenum, tungsten and chromium, ie a (Mo, W) C-Co alloy is harder than a WC-Co alloy at higher temperatures. In addition, when Cr is dissolved in this carbide, the hardness increases even more and the hardness at high temperature is also improved. The disadvantages of the known WC-Co alloy can thus be overcome in one go (see FIG. 4). It should be noted that the carbide phase consists of a solid solution of (Mo, W, Cr) C. It has also been found that when Cr is dissolved in a solid solution of (Mo, W) C, the carbide particles become finer and stabilized as a monocarbide of (Mo, W, Cr) C. In contrast, the known method in which only chromium is added to the binder phase has the disadvantages that it is not possible to obtain finer carbide particles and that the carbide phase does not stabilize as a monocarbide of a solid solution of (Mo, W, Cr) will. The amount of chromium to be added to the carbide (Mo, W) C in solid solution is preferably in the range of 0.3 - 10%, because with amounts below 0.3% the carbide particles cannot be made finer than 10%, Cr ₃ C _{2 is} separated and precipitated in the alloy, which leads to a reduction in hardness.

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In a further embodiment of the invention, part of the Carbon in the solid solution carbide (Mo, W, Cr) C is replaced by nitrogen, oxygen and / or hydrogen, i.e. it is It is believed that when the carbon contained in (Mo, W, Cr) C is added in solid form and reacts 100%, the crystal is stabilized will. However, it has now been found that the inclusion of not only carbon but also nitrogen provides stabilization of the monocarbide as (Mo, W, Cr) (CN) and that the further inclusion of oxygen and hydrogen leads to the monocarbide as (Mo, W, Cr)

(C N. O HJ, where a + b + c + d = l, stabilized even more because, at abcd ^p '

Presence of defects in the carbide, the carbide is unstable during sintering and a mixed carbide of the type M ₀ C is precipitated in a needle-like form, thus reducing the strength.

When the amount of chromium contained in (Mo, W, Cr) C is 0.3 to 10 wt.% Is limited in order to obtain a finer carbide, and if one or more of the metals of the iron group, such as iron, Nikkei and cobalt, as a binding phase in a ratio of 15 - 30% by weight are added, the alloy thus obtained is useful as a cemented carbide alloy used for cutting low speeds, e.g. in the form of drilling tools, thread cutters and dies, with excellent results can be used. When the amount of the binder metal is within the range of 3-15% by weight, the alloy is also more effective Way as a corrosion-resistant alloy. Application examples of the corrosion-resistant alloy are corrosion-resistant sealing rings, Clockwork frames, ends of slide gauges, mechanical seals, etc.

The material of a cemented carbide alloy is an alloy with selected relatively high cobalt content, which has a high deformation resistance. As Fig. 5 shows, the deforma-

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tion caused tendency to break when increasing the amount of cobalt. at However, as the amount of cobalt increases, the alloy has decreased resistance to deformation. This tendency to deform is disadvantageous when the alloy is in the form of a forging tool, e.g. B. as an upsetting punch, should be used, although a There is hardly any crack formation.

According to the invention, an alloy of high deformation resistance is obtained and sufficient elastic strength without reduced hardness is provided, and the properties are compared to those of the known Alloys of the type WC-Co improved, i.e. as a result of the detailed investigations carried out on alloys of the WC-Co type, it has been found that among the WC-Co type alloys, those in the WC-γ zone between the precipitation zone of the free carbon and the precipitation zone of the δ phase (η -Co "W" C-phase) lying alloys have excellent mechanical properties, so that the alloys lying in the WC-y zone mainly be used. In the γ phase, tungsten is dissolved in cobalt and, as is known, the properties of the alloys change with the amount of this solid solution. The deformation resistance depends on the amount of tungsten dissolved in the cobalt binder phase. If no tungsten is dissolved in the cobalt, the deformation resistance can decrease of the alloy are considered to be still further enlarged; however, this is meaningless if no free carbon is excreted.

An alloy was sought in which the binder phase was so pure can be kept as possible without free carbon being present at the same time, i.e. one aim of the invention is to to keep the deformation strength of the binding phase as high as possible, without any substantial amount of tungsten and molybdenum being dissolved in the binding phase, even if the amount of carbon is changed,

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and additionally to uniformly disperse a compound of the type MgC, which exists due to the lack of carbon, so that a concentration of the acting forces or loads is avoided. In general, alloys consisting of carbides of molybdenum and tungsten, has not been used in practice because the precipitation of a carbide (Mo, W) takes place ₀ C in the form of needle-shaped particles, leading to a considerable decrease in hardness or strength Alloy leads. The invention succeeded in increasing the deformation strength of the alloy without reducing the hardness by means of a good dispersion of the (Mo, W) “C”.

Fig. 6 is a graph showing a comparison of the relationship between compressive stress and deformation, as in the case of (Mo, W) C-Co alloys and WC-Co alloys according to the invention the state of the art is given. It turns out that the alloy of the type WC-Co according to the prior art have a deformation of approximately 2-4% under pressure, while the inventive Alloys show a deformation of 4 - 5%. An example of a WC-24 vol.% Co alloy is the crushing stress

ο
400 kg / mm and the deformation about 4%, while the invention

moderate alloys result in higher crushing stresses, ie 500 kg / mm, and a deformation of about 5 % .

In an alloy according to the invention in which the composition ratio of molybdenum and tungsten by the formula (MoW) C like-

x y

is given, the composition of the (Mo W) C in the alloy

x y

tion is not always limited to one, but two or more combinations can be used to change the properties of the alloys. A carbide of the type M "C, i.e. (Mo, W) "C, be uniformly dispersed to the desired To make an impact. The advantages of the invention are retained

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even if the alloys contain carbides, nitrides and carbonitrides of metals of groups IVa, Va and VIa. Further, the advantages of the invention can also be preserved when C in (Mo, W) C or (Mo, W) is replaced by ₀ C. N, O and / or H. These alloys, which in particular have excellent impact strength, can be used with advantage in tools for the production of metal containers, such as cans, matrices, mining tools, rollers, etc., as well as in upsetting punches and other known shock-absorbing tools.

In a further embodiment of the invention, one or Several of the metals manganese, rhenium, copper, silver, zinc and gold are included in the binder phase to make up the microstructure of the binder phase to change and create non-magnetic properties. It has been found that when these elements are added, the binder phase will alloy which leads to an improvement in the corrosion resistance of the alloy. The hardness and wear resistance of the alloy will be deteriorated when the amount of the binder phase exceeds 30% by weight. The wear resistance of the alloy is not reduced unless the amount of these elements exceeds 5% by weight.

It is known that in the case of alloys of the WC type, the alloys are designed as alloys with little carbon, see above that the non-magnetizability and the corrosion resistance are considerably improved. When the amount of alloyed carbon is less than in the WC-γ phase zone, tungsten dissolves in larger quantities in the binding phase and thereby reduces the magnetizability. On the other hand, in alloys of the (Mo, W) C type, it is difficult to lower the magnetizability even if the alloys combined as alloys with little alloyed carbon

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sets are. However, when nickel is used as the binding metal, a remarkable effect can be achieved in combination with the above additives. In the case of alloys with a low carbon content, a M “C phase is precipitated, which is preferably subject to corrosion because it is electrochemically relatively basic in comparison with a (Mo, W) C phase. However, it has been found that if an M "C phase is uniformly dispersed in the ratio of 30 % by volume or less in the alloy, corrosion of the alloy base does not occur, and the corrosion resistance of the entire alloy body is improved because the MJC phase is finer , is spherical in shape, d. A low carbon alloy is obtained in which an M ^ C phase is uniformly dispersed in a proportion of 30% by volume or less, and which is imparted with corrosion-resistant and non-magnetic properties by using nickel as a binder phase and adding at least one of the metals manganese, rhenium, copper, silver, zinc and gold.

In the production of the alloys according to the invention, these are inevitably contaminated by small amounts of impurities such as iron, cobalt, etc. As long as the total amount these impurities does not exceed 1% ^ will benefit the Invention retained.

In a further embodiment of the invention, the amount of iron present in the alloy is preferably controlled according to the relationship:

_{n nn1} L amount of iron (wt.%) ^ _{ft 1}

w, ^{UUi - amount of} binder metal (% by weight) "'

In this case, the hard phase consists of (Mo, W) C and the binding phase made of Co and Ni with the addition of Fe as an additional element. So it becomes

109827/0991

a carbide of the type (Mo, WLC precipitated in granular form and not in acicular form. The amount of Fe to be added as an additional element is preferably 0.1-10% by weight, because if the amount is less than 0.1%, the effect of the Fe is low and with an amount greater than 10 % the precipitated MC is too coarse to ensure the strength of the alloy. Fe can be added by adding it to the alloy or the reaction mixture during the production of the carbide. The reactivity in carburizing the carbide can be controlled by changing the amount of Fe, and if the amount of Fe is less than 0.1%, the carburization proceeds insufficiently and if the carbide thus obtained is used to make an alloy If the amount of Fe is more than 10%, the carbide is alloyed and crushed or ground is hardly dispersed in the alloy in a granular form becomes very difficult, resulting in a decrease in the yield of carbide useful for making a hard alloy. A (Mo, W) C alloy in which a carbide of the type M ₀ C according to a

Feature of the invention is excreted and dispersed, has as a so-called dispersion alloy has a high strength. When the particle size distribution and the dispersed state of the carbide of the type M "C not evenly or in the inner and outer Parts of the alloy are different, the alloy strength can not be so great be achieved. In the case of large alloy bodies or alloys with a high content of binder phase, e.g. Morgan rolls, Very high pressure pistons, anvils and the like often pose the following problem. If carbon diffuses into the alloy from the outside, so that there are considerable differences in the carbon content of the outer and inner layers of the alloy, or if the cooling rate on the surface and inside the alloy differ, the conditions for the precipitation of the

$ 09827/0991

Molybdenum and tungsten from the liquid phase unequal and the shapes and dispersion states of the M "C type carbide are at the Outside and inside the alloy are different. On the outside or surface of the alloy, the M "C tends to be a rough Taking shape and forming agglomerates, reducing strength.

This problem can be solved with an embodiment of the invention. If cobalt and nickel are used as the binder phase and iron is added in a ratio of 0.1-10 % , the carbide of the M "C type is stably precipitated and dispersed regardless of the amount of binder phase and the shape of the alloy so that a high strength of the alloy is achieved. If the binding phase consists only of cobalt, the M "C tends to form agglomerates. If the binding phase consists only of nickel, the alloy is of lower hardness and compressive strength. If the amount of iron in the binder phase is less than 0.1%, the effect of iron is poor, while if it is more than 10%, the corrosion resistance and mechanical strength of the alloy are deteriorated.

In this embodiment, it is also appropriate that a binder phase of an iron group metal as a main component at a ratio of 3-50 wt -% of the total composition is present.. If the ratio is less than 3% by weight, the alloy is too brittle, while if the ratio is more than 50% by weight, the breaking temperature properties are impaired. The iron group metal as a binder phase naturally dissolves the IVa, Va and VIa metals, and the advantages or effects of the invention are not lost when elements soluble therein such as aluminum, silicon, calcium, silver, etc. are added .

009627/0901

The basic concept of the invention is retained even if part of the Molybdenum and tungsten carbide is replaced by a mixed carbide of the type B1, the titanium, zircon, hafnium, vanadium, niobium, tantalum, Contains chromium, molybdenum and / or tungsten. Furthermore, the properties of the alloys according to the invention are not changed significantly, when part of Mo and W in (Mo, W) (CNO) or (Mo, W) C is replaced by other elements, as long as the simple hexagonal structure preserved.

As with this embodiment of the invention from the points of view As can be seen, a micro-amount of iron is required as a stabilizer in a (Mo, W) C alloy or a (Mo, W) (CNO) alloy. As a method of dispersing the iron, it is appropriate to add the iron during the formation of the carbide and the carburization reaction at a temperature of 1,500 C or more in a stabilizing atmosphere consisting of nitrogen or carbon monoxide perform.

A method for producing (Mo, W) C is known, which consists in adding a large amount of a diffusion aid such as iron or cobalt to Mo ₀ C and WC and carrying out the reaction at 2,000 C or more ( Japanese Patent Application '(OPI) No. 146 306/1976). In this process, iron is added to promote the reaction to form a solid solution of WC and Mo ₀ C.

In the embodiment of the invention, a small amount of iron can be added when the solid solution of Mo-W is ready, (Mo, W) alloy powder is carburized or when a (Mo, W) oxide is carburized directly. The added iron becomes the stabilization reaction used at a temperature of 1,500 ° C or more and has no adverse effect on the carbide.

© 09827/0991

In this embodiment, some of the carbon can be in the carbide can also be replaced with nitrogen and / or oxygen while achieving substantially the same effects.

In the last embodiment of the invention, the strength of the Alloy can be increased by two or more carbides, which have a simple hexagonal phase, but in the ratio Mo / W differ, can be used in combination with each other. The cause of the increase in strength is not known in detail, however it is believed that when (Mo, W) C is separated into two phases, the dissolution pressure of both phases is decreased, resulting in strength which is greater than in the case of a single phase. An alloy consisting of a (Mo W) C (y > x) phase, which has similar properties to WC and a (Mo W) C (x> y) phase, which has properties similar to MoC, has two properties, and the hardness of the toilet and the heat and deformation resistance of the MoC. For this reason, this embodiment is advantageous over the use of only one type of (Mo, W) C. Preferably, the carbide consists of WC or a solid solution of a small amount Amount of MoC dissolved in WC and a solid solution of WC dissolved in MoC. This corresponds to the case of X-ray diffraction analysis, at which the maximum of the plane (1, O, 3) splits into two parts. The question of whether two or more simple hexagonal phases of (Mo W) C exist or not can be clarified by observation with an optical microscope on a sample, that with an alkaline Solution of a hexacyan ferrate (III) has been etched or by XMA observation.

The scope of the alloys according to the invention is as follows. For example, the alloys according to the invention used for the production of wear-resistant tools, such as guide

909827/0991

rollers, rollers for hot wire rolling equipment etc. and cutting tools, because they have strength and hardness similar to or better than those of WC-Co alloys. Especially if a Alloy according to the invention as a substrate with one or more wear-resistant ceramic layers such as TiC, TiN, A1_O ", coated is, it can be used to produce cutting tools that, compared to the tools according to the prior art, with substrates made of alloys type WC-Co have excellent mechanical strength and wear resistance. As is well known, one is referred to as the η phase Carbon stripping layer is formed at the interface between the substrate and the coating. This arises in a similar way Way with the alloy according to the invention. The emergence of brittleness just below the coating due to the withdrawal of Carbon is effectively prevented by the presence of free carbon in a surface layer of about 300 μm thickness, without reducing the strength.

For use in the production of cases for watches, the alloys according to the invention have the same properties compared to alloys of the type WC-Co has excellent properties, which are summarized below:

(1) An attractive appearance is given by polishing to a mirror finish.

(2) Grindability and polishability.

(3) Excellent corrosion resistance, especially against body sweat when used as a material for jewelery.

(4) The mechanical strength is great.

According to the invention, the strength of an alloy can be compared with

POS 827/0991

Increase deformation without deteriorating mechanical strength by _{dispersing (Mo, W) 0} C well or evenly. A carbide of the type M ₀ C alone has a low hardness (Vickers

2
Hardness of Mo C: 1,500 kg / mm). However, if this M ₀ C is uniformly dispersed in an alloy, the alloy maintains a high strength without the overall hardness of the alloy decreasing because the soft M ₀ C can dampen an impact force acting on the alloy. Because of the high wear resistance and mechanical strength at a low price, the alloys according to the invention are suitable for producing spikes for shoes or spikes for ice. When an M ₀ C carbide is suitably or uniformly in an alloy

is dispersed, the alloy has good sliding properties on a Concrete surface and is able to absorb the shocks resulting from the roughness of the concrete surface.

The invention is explained in more detail with the aid of the following exemplary embodiments. For those skilled in the art it is apparent that the The following recipes, the proportions and components as well as the order of the process steps within the framework of the invention can be modified. Thus, the invention is not to be construed as being limited to the following examples. All shares, Percentages and similar data are to be taken as weight-related unless otherwise stated.

example 1

The starting material was (Mo _n "W _n ") C, which contains 0.2% Fe as an additive

U, ι U, 0

set included, used. The starting materials were determined by weight so that the overall composition can be given by the formula (Mo _n ^ W _n JC -15% Co, in which ζ (carbon content

U, I Uj O Z

in the alloy / theoretical carbon content) 100, 98 and 96 Atom%. The materials were wet in an or-

109827/0998

Chemical solvents mixed, dried, compressed or compressed and sintered at 1,450 C in a vacuum. For comparison, this procedure was repeated with the exception that one Fe-free starting material was used. The properties of the alloys obtained are shown in Table 1.

Transverse fracture
strength
(kp / mm) Tabel 1 Total-
carbon Suitor
carbon Z Fe-containing ends 190 hardness
(HRA) 7.61 0.02 100 jr I UUc
(A) 290 86.8 7.43 CLOO 98 (B) 260 87.1 7.28 0.00 96 (C) 87.2 Fe-free
sample 175 7.59 0.00 100 (D) 140 86.8 7.43 0.00 98 (E) 120 87.0 7.28 0.00 96 (F) 87.1

As can be seen from these results, the alloys (B) and (C) of the present invention are better in strength than the alloys (A), (D), (E) and (F) of the prior art. In the case of the alloys according to the invention, the strength also did not decrease if the carbon content was less than the theoretical carbon content, while in the case of the alloys according to the prior art, a molybdenum-tungsten-Mix carbide of the type M ₀ C due to a lack of carbon in the form of acicular shapes Crystals precipitated, which led to a decrease in the strength of the alloy.

Fig. 1 shows a microscopic photo of the alloy (E) according to the prior art and FIG. 2 shows a microscopic photo of the Sea alloy (B) according to the invention,

I09827 / 099S

Example 2

A powdery solid solution of (Mo _{n Q} WJ with a particle

Uj ti Uj 1

size of 6 μηι was mixed with O, 2% Fe powder and there were various amounts of carbon were added in order to achieve the ζ values given in Table 2. The mix was Carburized for one hour at 1,600 ° C in nitrogen gas and pulverized. The carbide was heated in CO gas for 30 minutes and stabilized. The carbide obtained was a carbide in which MC and MC phases coexisted, as shown in Table 2. The carbide was mixed with 10% Co and 10% Ni and sintered at 1,300 ° C. The properties of the alloys obtained in this way are in shown in Table 2.

; Carbide
Total-
carbon O %
⁰ Ic 0.3 Tabel 2 180 0.4 (Mon W) (
ζ 10.29 0.2 260 0.7 1.0 9.74 0.2 0.2 280
N 0.8 0.95 9.22 0.2 0.05 Alloy properties
Transverse fracture Charpy
C Vol. - ⁰ I ₀ strength value
(kp.mm) 220 0.7 0.9 8.23 0.1 0.05 0 170 0.3 0.8 7.20 0.1 0.03 7th 140 0.2 0.7 6.17 0.1 0.02 14th 100 0.1 0.6 5.14 0.01 30th 0.5 0.01 45 55 75

As can be seen from this table, no practical one can be found Alloy strength obtained in cases where the values of ζ in the formula (Mo, W) C are 1.0 or less than 0.7.

109827/0991

Example 3

There were 80-90% of a carbide which had the theoretical content of bound carbon (Mo _n "W _n ") C and a particle size

Uj I Uj O

ße of 5 μm, 0-10 % (Mo _n "W _{n Q} ) _O C with a particle size

Uj I Uj Ö Ci

of 3 μΐη, 9-10% Co and 0.1-0.5% Fe, Re, Si, B and Be mixed together and alloys produced therefrom in a manner analogous to Example 1. A carbide of the type M "C, ie (Mo, W)" C, as shown in FIG. 2, was uniformly dispersed in the structure of the alloys produced in this way, while the alloy (Mo, W) C- (Mo , W) ₂ C-Co acicular crystals were precipitated according to the prior art, as shown in FIG. For comparison, the properties of these alloys are shown in Table 3:

Table 3

attempt ( ^Mo O, 7 ^W 0.3> ^C ( ^Mo 0.7 ^W 0, ₃ ) ₂ C Co Dispersions Transverse
fracture-
strength
(kp / mm) 1 85 5 9.9 Fe 0.1 260 2 82 8th 9.5 Re 0.5 190 3 88 2 9.8 Si 0.2 240 4th 88 1 9.9 B 0.1 270 5 87 3 9.7 Be 0.3 200 6th 87 3 10 - 150 7th 82 8th 10 - 120 8th 89 1 10 _ 160

Note: Experiments 1-5 are according to the invention. Experiments 6-8 are according to the state of the art

0S827 / 099

As can be seen from this table, the alloys according to the invention are in which {Mo, W) “C is dispersed by adding an impurity has been of high strength.

Example 4

There were 86% (Mo _n "W _n") C powder with a particle size of

Uj IU, O

5 μηα, 5% (Mo _n "W _{n") O} C Powder μπι having a particle size of 2, U, ι Vf ο Δ

9% Co powder and O, 2% Fe powder balanced, in the wet process in mixed with an organic solvent, dried, compacted and sintered at 1,450 C in a vacuum, resulting in an alloy,

2 has a transverse breaking strength of 260 kp / mm and a hardness (Hv) of

1,400 kgf / mm. Thereafter, the alloy was subjected to carburization to free carbon in the surface region of 300 meters depth and excrete with a layer of TiC, a double layer of TiC and TiN or a double layer of TiC and Al ₀ O "

overdrawn. For comparison, a sample of a commercially available WC type alloy plated in a similar manner. The insoles thus obtained, designated Samples 1-6, were a cutting test under the following conditions (form no. SNU 432) subjected:

Workpiece SCM 3 (H _D = 280)

Cutting speed ν = 170 m / min

Feed f = 0.86 mm / rev

Cutting depth d ■ = 1.5 mm

Cutting time 30 min

The test results are shown in Table 4.

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Table 4 V _B (mm) K _T (mm) sample Coating layer In accordance with the invention 0.16 0.09 1 TiC 0.14 0.03 2 TiC, TiN 0.12 0.02 3 TiC, Al O _q Toilet alloy 0, 21 0.10 4th TiC 0.17 0.04 5 TiC, TiN 0.15 0.03 6th TiC, Al ₉ CL

As can be seen from these results, the alloy according to the invention is similar to or better than with respect to the wear of the front cutting surface (flank wear V-.) And the crater depth (K_,)

Ϊ5 1

Prior art WC-type alloys. In the inventive Alloy no decarburized layer (η phase) was found at the interface between the substrate and the TiC layer, while free carbon in a range of 300 μπα depth directly was found below the coating layer.

Example 5

Mo_C powder with a particle size of 2 μτα, WC powder with a particle size of 2 μπι and carbon powder with Co powder as a diffusion aid were mixed in order to achieve the overall composition (Mo "gW _Q 2 ^ ^C 0 95 ^N 0 05 ^ 1 0 ^{tO he just>} after which the mixture hydrogen stream of nitrogen was contacted with a nitrogen partial pressure of 0, 5 bar of reacting S for 30 minutes at 1,800 ° C in one. an X-ray diffraction analysis showed the formation of a simple hexagonal crystal of the type WC.

109827/099

This powder was mixed with Co powder to give the final alloy composition (Mo _n gW "J (C, N) -IO % Co, and the mixture was compacted into a suitable shape and sintered. Sintering was carried out by heating in vacuo at 10 Torr up to 1,000 ° C. and in a Co atmosphere under reduced pressure of 10 Torr from 1,000 ° C. to 1,400 ° C. For comparison, however, an alloy was produced in a similar manner, but without the use of nitrogen in the Production of the carbide and without the use of carbon monoxide in sintering The results are shown in Table 5.

Table 5

Composition of the hard phase A value B value

structure

According to the invention

), 98 0.04 0.01 WC phase + Co phase

State of the art

( ^M 0.8 ^W 0.2 ^ ^C 0.976 ^N 0.004> 0.98 ° '° ⁰⁴

0.00 WC phase + needle-shaped M “C phase + free carbon + Co phase

The alloy of the present invention is light in weight and excellent in impact resistance and high-temperature hardness as seen from test results. Consequently, the alloy according to the invention is suitable for production of various tools, especially wear-resistant tools.

Example 6

A previously prepared, pulverized, solid solution of (Mo _n JW _{n Λ κ}

U, ι U, Iu

Cr

9103

with a particle size of 2 μπι was with carbon and

909827/0990

0.2% Fe mixed as a diffusion aid, carburized in hydrogen at 1,800 ° C and then reacted for 30 minutes at 1,300 C in a mixed gas of nitrogen and carbon monoxide. The hard phase obtained in this way, which has the overall composition (Mo _n -W _{n 1K}

U, ι U, ID

Cr _{n 15} ) (Cq qqNq _n g0 _{n ol} ) was mixed with 9.5% of a binder metal consisting of Co / Ni (1/1) and a micro amount of Fe, and sintered. An analysis by X-ray diffraction showed that the obtained alloy was composed of a hexagonal monocarbide (Mo, W, Cr) C and a (Mo, W, Cr) ₂ C phase together with the binder phase. Due to the structure, a granular carbide of the M "C type was uniformly dispersed in the alloy.

This alloy had better corrosion resistance, in particular to body sweat, compared to prior art WC-type alloys and is suitable for the production of Jewels, such as watch cases, due to its light weight, and as a non-magnetic alloy.

Example 7

There were 60% (Mo _{n 7} W _{Q 3} ) (C _ng N _{n Q8} ), produced by mixing (Mo, W) O ₃ ^with carbon and 0.05% Fe and react at 1,700 ° C in nitrogen at 1 , 1 bar, 30% (Ti _{n 7} W _{n 3} ) (C _{Q 8 &} Ν _{η 15} ), 5% Co and 5% Ni mixed and ground in a ball mill in the wet process. The powder mixed in this way was compacted, then heated to 800 ° C. under vacuum, heated to 1,200 ° C. in H ₂ at 30 torr and heated in CO at 30 torr at 1,200-1,400 ° C. and for one hour at Held at 1,400 ° C to complete the sintering.

An analysis of the alloy showed the result: A = 0.11 and B = 0.02. An examination of the structure revealed the existence of a (Mo, W) (CN) phase and a (Mo, W) “C phase, which are well in a spherical shape were dispersed.

§09827 / 0991

For comparison, a nitrogen-free alloy was sintered in Vacuum created.

These samples were subjected to a cutting test under the following conditions subjected:

Workpiece simple steel (hardness: RC 20-29)

Cutting speed 150 m / min

Feed 0.381 mm / rev

Cutting depth 0.13 mm

The results are shown in Table 6.

According to the invention
Alloy

alloy
after this
State of
technology

Table 6

Cutting time flank crater -

wear wear (min) (mm) (mm)

0.15

0.20

0.010

0.017

Edge deformation (mm)

0.006

0, 015

As can be seen from these results, the alloy of the present invention is of better wear resistance as well as better resistance to edge deformation than the alloy according to the state of the art of the technique. After coating the alloy according to the invention with one or more carbides, nitrides, oxides and borides in Form of a single or multiple layer; to so-called coated Forming deposits remained the excellent strength compared to

109827/099

Edge deformation obtained in the alloy.

Example 8

Amounts of MoO "powder and WO" powder were weighed out to to give a value of 8/2 for the ratio Mo / W, and with a Amount of carbon sufficient to remove the oxygen in the oxides and 0.2% Fe as a catalyst to fix nitrogen mixed during the reaction. The mixture was reacted for one hour at 1,500 ° C. in a gaseous stream of (NEL + 10% by volume CO)

greed to complete the reducing reaction.

An X-ray examination of the compound obtained showed the formation a compound of (Mo, W) (CNO) of a hexagonal type.

This carbide was mixed with 10% Co and Ni, and an alloy was produced therefrom in a manner analogous to Example 5. An analysis of the alloy obtained showed: A = 0.2 and B = 0.03. An examination of the structure of the alloy showed that a granular carbide of the type M ₀ C in an amount of 2% by volume dis-

Cl

was pergated. This alloy was particularly excellent Shock resistance on.

Example 9

527 g of WC powder with a grain size of 1 μm, 430 g of Mo ₀ C powder with a grain size of 2 μm and 13 g of Cr _Q C "were mixed,

German O Ct-

after which 5 g of Co powder as a diffusion aid and 27 g of carbon black for replacement of the missing carbon were added. The mixture was ball milled in a dry process for about 30 hours. The powder mixed in this way was reacted at 2,000 ° C in a stream of hydrogen to form a primary carbide.

909827/0391

2855513

the. The primary carbide was ground well, and then a secondary Carburized. In the secondary carburization, the carbide was each under one of the following conditions reacts further, whereby part of the carbon in the carbide (Mo, W, Cr) C is replaced by oxygen, nitrogen and hydrogen and the carbide was further stabilized:

I) in a _{stream of NH 3} , 1 hour at 1,400 ° C, II) in a stream of CO, 2 hours at 1,600 ° C,

III) in a _{stream of H 2} , 1 hour at 1,500 ° C and

IV) in a vacuum, 1 hour at 1,500 ° C.

Various carbides were obtained as shown in Table 7 is shown.

Table 7

Condition Total Free O ₉ N ₉ H ₉ Co

carbon carbon

D. 8.33 Π) 8.32 in) 8.45 IV) 8.53

0.08 0.03 0.25 0.01 0.5

0.05 0.21 0.05 0.001 0.5

0.03 0.003 0.01 0.02 0.5

0.30 0.001 0.001 0.001 0.5

The carbides produced by heating in gaseous atmospheres were all monocarbides, while the carbides produced by heating in vacuum free carbon in larger quantities and excreted Mo “C contained, as from the results of an X-ray analysis emerged.

9103

§09827 / 099

A (Mo, W, Cr) C-16% (Co + Ni) alloy was then prepared using the carbide (II) of the carbides described above. An MpC phase was found uniformly dispersed in the alloy in a ratio of 10 % by volume.

The properties of the alloy according to the invention and the known WC-Co alloy were made using an end mill compared under the following conditions:

Workpiece SCM 3, HRC 8-13, length 385 mm,

Face milling machine 8 mm j3, two cutting edges,

right cutting edge right twist 25 °, fixed

Cutting system Cutting a groove 5 mm deep

in the workpiece specified above, comparison of the service life by measuring time or cutting length up to V _D = 0.3 mm or up to crack or chip formation

Device simple milling machine no.4

Cutting conditions V = 26.5 m / min

f = 0.0285 mm / edge water-insoluble cutting oil.

The results are shown in Table 8.

909827/0991

Table 8

Cutting length number of
milled grooves 36.3 m 92.6 24.5 62.5 1.54 4th 2.55 67

Alloy according to the invention

WC fine particle alloy
according to the state of the art

Sintered carbide alloy KlO
according to the state of the art

High speed steel

As can be seen from these results, the alloy of the present invention is superior to other known alloys in terms of wear resistance and chip formation resistance, because the inventive Alloy has a high high temperature hardness and consequently the strength can be increased without reducing the wear resistance, even if the amount of binding phase is increased.

Example 10

The solid solution carbide (Mo, W, Cr) C obtained by following the procedure of Example 9 was mixed with 10% Ni powder and 100 hours in a wet process in an organic solvent for a sufficient treatment in one Subjected to ball mill. The powder mixed in this way was compacted under a pressure of 100 MPa and alloyed at a temperature of 1,400 ° C. A (Mo, W) “C phase was found in the alloy obtained, in which particles with a size of 1 μm or less were uniformly dispersed in a ratio of 5% by volume.

The obtained alloy was polished using a diamond paste to give a mirror finish. The physical properties

shafts of this alloy are shown in Table 9, which shows that the alloy according to the invention is superior to an alloy of the type WC-Co according to the prior art with regard to density, Strength and corrosion resistance to body sweat.

Table 9

Transverse fracture density hardness strength (kp / mm)

Corrosion resistance

Alloy according to the invention

Prior art WC-Co alloy

10.8 91.0 180

14.5 90.5 160

good not good

The corrosion resistance given in Table 9 was measured by immersing the sample in a man-made sweat for 48 hours.

A watch frame made from this alloy was subjected to a property test. The invention proved to be Alloy as more useful due to its light weight and its scratch resistance and resistance to body perspiration.

Example 11

80 % of a carbide of molybdenum and tungsten with a molar ratio of 7: 3, (Mo _n "W _n JC, with 10% Co and 10 % Ni

U, ι U, ο

mixes, after which the amount of carbon was controlled so that the carbon content in the alloy was 98 atomic percent based on a carbon content of 7.11% in the stoichiometric composition the alloy.

9103

909827/0998

As a comparative experiment, samples (I) to (ΠΙ) were as follows manufactured as described:

Sample (I): There was 0.1% Fe as an impurity of those described above Alloy added to (Mo, W) “C according to the Invention to disperse evenly.

Sample (II): The sample described above was used as in prior art Technology, no pollution added.

Sample (III): A WC-18 vol.% Co alloy for slotting tools according to the state of the art.

In the invention, (Mo, W) “C was uniformly dispersed in a granular form in a ratio of about 8% by volume, as shown in FIG. 2, while in the case of the (Mo, W) C-Co alloy according to the prior art Technique needle-shaped crystals were precipitated, as shown in Fig. 1.

Upsetting dies and theirs were made from these alloys Life span studied, by being subjected to plastic machining subjected the head to a screw consisting of an SCr-4 steel rod and the results shown in Table 10 were established.

Table 10

0 2 4 6 8 10 χ 10

_ | I I I I I

(Number of screws / die) 2

Sample (I)>

Sample (II) χ

Sample (ΙΠ) ■ χ

Note: χ indicates a broken tip 9103

909827/0998

As can be seen from the test results, shows the invention Alloy a maximum service life and sufficient usability, even if cracks or deformations occur.

Example 12

A (Mo _n -W _n JC powder, where ζ = 0.9, was prepared, with 15%

U, I U, «J Z

Ni powder and 2% Mn powder mixed, treated sufficiently in a wet process in a ball mill, compacted and sintered at 1,350 ° C in a vacuum. An examination of the structure showed the presence of granular (Mo, W) “C with a particle size of 2μηι ₎ which was dispersed in a ratio of about 10 % by volume.

The properties of the alloy thus obtained were as follows:

Density: 9.88 g / cm ³

Hardness: H _RA = 88.5, H _c = 0.4πσ = 0

Transverse breaking strength: 170 kp / mm

This cemented carbide alloy was not magnetic.

Example 13

The amount of carbon in an alloy, consisting of 85% (Mo _n "W _n JC, 16 % Ni, 0.6% Mn and 3% Re, was determined in this way.

U, / U, u

controls that the content of alloyed carbon was 95 atomic%, based on the theoretical carbon content (7.59%) and it was 0.1% Fe added to the alloy. The mixture was an hour sintered at 1,450 C in vacuum, resulting in an alloy with the following properties:

309827/0998

Density: 10.25 g / cm ³
Hardness: H _RA 89.4
Transverse breaking strength: 165 kp / mm 4tux = 0

^H C = °.

About 1% by volume of an M “C phase was found in the structure of this alloy. The corrosion resistance of the alloy according to the invention and that of the WC-7% Co alloy according to the prior art are listed in Table 11.

Table 11

2 (unit: mg / cm / h)

Hot 10% hot 35% .. H ₂ SO ₄ solution HNO ₃ solution ^ALKaLl

WC-7% Co alloy 15 9 0

According to the invention

Alloy 0.4 4 0

Example 14

Using (Mo _n "W _n ") C, which is obtained by adding 0.2% Fe

U, ι U, <3

While the carbide was being produced, the composition (A) (Mo _n -W _n ") C-15 % Co, the composition (B) (Mo _n " W _n JC-7, 5% Co -7.5% Ni and the

Uj I OfO

Composition (C) (Mo _Q “W _{Q 3} ) C-15% Ni obtained. The raw materials were weighed so that the content of alloyed carbon ζ in (Mo _{Q 7} W _{Q 3} ) C _{z was} 98 atomic%. The raw materials were subjected to a ball mill treatment in an organic solvent by the wet method, dried, compacted and then sintered at 1,350 ° C. in vacuo to give alloys having the properties shown in Table 12.

109827/0998

285G513

For comparison are the alloy properties of compositions prior art, i.e. Composition (D) (Mo, W) C-15% Co and Composition (E) (Mo, W) C-7, 5% Co-7, 5% Ni as well listed in Table 12. In the case of the alloys according to the invention (A), (B) and (C), the total content of oxygen and nitrogen was 0.15.

Table 12

According to the invention

Composition (A)

Composition (B)

State of the art

Composition (D)

Composition (E)

! * u ττ " _t total suitors

fracture-hard. ,, __,. Value z

r Li -ι. ,, π,., coal- coal- ~ _λ

estigke ^ (HRA) _{stQff gtoff} (At-%)

(kp / mm)

240 87.0 7.43 0.00 98

290 86.9

Amount of iron in binder

ΓΛ I -v ^ H η ** * * *

175 86.8 7.59

140 86.6 7.43

1.1

100

As can be seen from this table, the alloys of the invention are (A), (B) and (C) of greater strength than the alloys (D) and (E) according to the prior art.

According to the invention, the alloy is stabilized and consequently the Strength of the alloy is not due to the presence of Fe, N and O in of the alloy, even if the alloyed carbon content is less than the theoretical carbon content. In the inventive Alloy was granular (Mo, W) «C uniformly

9103

909827/0990

In the case of the alloy according to the prior art (Mo, W) ₀ C was precipitated in the form of needle-like crystals due to the lack of carbon, which led to a reduction in strength.

Example 15

Starting materials were mixed in such a way that the hard phase of an alloy obtained from (Mo _n "W _n ") C with a particle size

Uj i Uj O

of 5 μm and (Mo _Q "W _Q J ₀ C with a particle size of 3 μm and the binder phase was composed of Co and Ni with 0.1 to 1.0% Fe. From these starting materials, as in Example 14 In the structure of the alloys produced by this procedure, carbides of the type M ₀ C were uniformly dispersed. The properties of the alloys according to the invention are shown in Table 13.

Table 13

(Mo W JC (Mo W IC Co Ni amount of iron transverse fracture

O- 7 U- 3 U, ι U, ο 2 · ι-ΐ · j j? ι · ι · ι_

, '~' , nt ^, n / x, otK in bond strength

Attempt <%> ft) (%) (%) _age ^B

^{(kp / mm}}

1 85 5 10 _ 0.1 260 2 ti Il Il - 1.0 275 3 Il II 6.6 3.4 0.5 300 4th Il Il 7.5 2.5 0.5 290

In the alloy according to the invention, the strength can be increased by adding Fe as an additional element to the binder phase consisting of Co and Ni, as a result of which a carbide of the type M ₀ C is well dispersed in the alloy.

909827/0993

Example 16

(I) WC powder with a particle size of 6 μηα was mixed with Mo ₀ C powder with a particle size of 2 μηα and carbon in amounts which led to the production of a carbide of the final composition (Mo _Q -W _{0 5} ) C. The mixture was subjected to a ball mill treatment by the wet method for 30 hours, and the mixed powder was reacted for one hour at 2,000 ° C. in a! "Stream. The carbide had a total carbon content of 7.81 % and a free carbon content of 0.03 %. The content of bound carbon was close to the theoretical value according to the formula (Mo _{n K} W _{n K} ) C. Investigations showed in the X-ray diffraction

Uj 0 Uj D

gram a disappearance of the Mo ₀ C maximum and the appearance of only the (Mo, W) C maxima. An examination of the cross section of the powder revealed that the powder had a core structure. An examination of the maxima on the side of the angles with high values showed that there were two separate phases.

(II) WC powder with a particle size of 1 μm was mixed with Mo ₀ C powder with a particle size of 2 μm and carbon in amounts which led to the production of a carbide with the final composition (Mo _{n K} W _{n K} ) C . After that, 0.5% Co was used as

U, U Uj 0

Diffusion aid added. The mixed powder was heated in a stream of _{H 0 at 2,000 ° C. for one hour.} The temperature was then lowered to 1,400 ° C. and the product was held at this temperature for 10 hours. An analysis of the product showed a total carbon content of 7.51%, a free carbon content of 0.05% and a Co content of 0.5%, the carbon content being close to the value of the theoretical carbon content exhibited. An examination by X-ray diffraction showed no separation into two phases but the presence of a single phase.

909827/0998

From the procedures (I) and (II) described above produced carbides, alloys were produced. The starting materials were weighed to obtain a composition according to the formula To give (Mo, W) C-15% Co, subjected to treatment in a ball mill by the wet process in an organic solvent, dried, compacted and sintered at 1,400 ° C. in a vacuum, whereby alloys with the properties shown in Table 14 were obtained.

Table 14

_j. ,, Hardness, transverse breaking strength

According to the invention (I) 11.6 88.6 280

State of the art (Π) 11.5 88.0 180

In the alloy according to the invention, there were two phases from that Art MC and a granular M "C phase that was well distributed, whereby the alloy has been given high strength while maintaining the properties of cemented carbide alloys of the type State-of-the-art WC-Co. In contrast, the alloy according to the state of the art consists of a uniform solid solution, which, however, has a lower strength.

9 0 9 8 2 7/099

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Claims (19)

Liedl, Nöth, Zeitler Patentanwälte 000 Munich 2 2 · Steinsdorfstrasse 2 1 - 2 2 · Telephone 089/22 94 41 856513 SUMITOMO ELECTRIC INDUSTRIES, LTD., 15, 5-chome, Kitahama, Higashi-ku, Osaka, Japan Hard alloy containing molybdenum and Wolfram claims: 1. Hard alloy containing molybdenum and tungsten, characterized by a hard phase composed of at least one compound with a simple, hexagonal crystal structure of the type MC, where M is metal and C is carbon, selected from the group consisting of mixed carbides, carbonitrides and carboxyznitrides of Molybdenum and tungsten as the predominant component, and a binder phase composed of at least one element selected from the Group consisting of iron, cobalt, nickel and chromium, in which a hard phase consists of a compound of the type M «C with a hexagonal Crystal structure is evenly dispersed. 2. Hard alloy according to claim 1, characterized in that the compound of the type M «C in a ratio of a maximum of 30% by weight, based on all hard phases, uniformly is dispersed. 3. Hard alloy according to claim 1, characterized in that the compound of the type M "C in granular or A 9103 - J / W 909827/0998 «" »· *. Goods» spherical shape with a particle size of 10 µm or less is. 4. Hard alloy according to claim 1, characterized in that the atomic ratio of carbon content in the hard phases of the alloy to the theoretical carbon content of the compound of the type MC is 0.98 to 0.8. 5. Hard alloy according to claim 1, characterized in that that part of the connection of type MC is replaced by a hard connection of type B1. 6. Hard alloy according to claim 5, characterized in that the hard connection of the type Bl at least contains one of the metals titanium, zirconium, hafnium, vanadium, niobium, tantalum, chromium, molybdenum and tungsten. 7. Hard alloy according to claim 5, characterized in that the amount of the replaced by the hard-type compound Bl compound of 30 wt -% by.. 8. Hard alloy according to claim 1, characterized in that at least one of the mixed carbides is a solid solution of (Mo, W, Cr) C. 9. Hard alloy according to claim 8, characterized in that the amount of Cr is 0.3 to 10% by weight. 10. Hard alloy according to claim 1, characterized in that part of the carbon in the hard phases ⁹¹⁰³ I09P27 / 099 forming carbides by at least one of the elements nitrogen and Oxygen is replaced. 11. Hard alloy according to claim 10, characterized in that that the amounts of nitrogen and oxygen are defined by the following relationships to the alloy composition: W atom% ., N atom% Group IVa, Va, VIa Metal Atom-% Group IVa, Va, VIa Metal Pixa-% «_ W atom% v, O atom% , Group IVa, Va, VTa Metal Atom-% Group IVa, Va, VIa Metal Atom-% ^ 0.05 12. Hard alloy according to claim 1, characterized in that that the binder phase is contained in a weight ratio of 3 - 50% of the alloy composition. 13. Hard alloy according to claim 1, characterized in that that the amount of iron in the alloy composition is defined by the following relationship: Amount of iron (% by weight) ^ ₀₁ Amount of binder metal (% by weight) 14. Hard alloy according to claim 1, characterized in that the dispersion of the hard phase consists of a compound §09827 / 099 28S6513 dung of the type M "C by adding an impurity element to the Binder phase is carried out. 15. Hard alloy according to claim 14, characterized in that the impurity element at least one of the Elements is beryllium, magnesium, calcium, boron, silicon, phosphorus, manganese, iron and rhenium. 16. Hard alloy according to claim 14, characterized in that the impurity element in a ratio from 0 to 3% by weight is added. 17. Hard alloy according to claim 1, characterized in that the dispersion of the hard phase consists of a compound of the M «C type is carried out by controlling the temperature during sintering. 18. Hard alloy according to claim 1, characterized in that at least one of the metals manganese, rhenium, Copper, silver, zinc and gold are added to the binder phase to make a non-magnetic alloy. 19. Hard alloy according to claim 1, characterized in that that the hard phase consists of a compound of the type MC of two or more simple hexagonal phases, which are a different Have Mo / W ratio.