WO2010046224A2 - Metal powder - Google Patents

Metal powder Download PDF

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Publication number
WO2010046224A2
WO2010046224A2 PCT/EP2009/062844 EP2009062844W WO2010046224A2 WO 2010046224 A2 WO2010046224 A2 WO 2010046224A2 EP 2009062844 W EP2009062844 W EP 2009062844W WO 2010046224 A2 WO2010046224 A2 WO 2010046224A2
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Prior art keywords
binder
wt
molybdenum
alloyed
alloy
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PCT/EP2009/062844
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German (de)
French (fr)
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WO2010046224A3 (en
Inventor
Benno Gries
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H.C. Starck Gmbh
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Priority to DE102008052104.3 priority Critical
Priority to DE102008052104 priority
Priority to DE102008052559.6 priority
Priority to DE200810052559 priority patent/DE102008052559A1/en
Application filed by H.C. Starck Gmbh filed Critical H.C. Starck Gmbh
Publication of WO2010046224A2 publication Critical patent/WO2010046224A2/en
Publication of WO2010046224A3 publication Critical patent/WO2010046224A3/en

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    • CCHEMISTRY; METALLURGY
    • C22METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22CALLOYS
    • C22C29/00Alloys based on carbides, oxides, nitrides, borides, or silicides, e.g. cermets, or other metal compounds, e.g. oxynitrides, sulfides
    • C22C29/02Alloys based on carbides, oxides, nitrides, borides, or silicides, e.g. cermets, or other metal compounds, e.g. oxynitrides, sulfides based on carbides or carbonitrides
    • C22C29/06Alloys based on carbides, oxides, nitrides, borides, or silicides, e.g. cermets, or other metal compounds, e.g. oxynitrides, sulfides based on carbides or carbonitrides based on carbides, but not containing other metal compounds
    • C22C29/08Alloys based on carbides, oxides, nitrides, borides, or silicides, e.g. cermets, or other metal compounds, e.g. oxynitrides, sulfides based on carbides or carbonitrides based on carbides, but not containing other metal compounds based on tungsten carbide
    • CCHEMISTRY; METALLURGY
    • C22METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22CALLOYS
    • C22C1/00Making alloys
    • C22C1/04Making alloys by powder metallurgy
    • C22C1/0433Nickel- or cobalt-based alloys
    • CCHEMISTRY; METALLURGY
    • C22METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22CALLOYS
    • C22C29/00Alloys based on carbides, oxides, nitrides, borides, or silicides, e.g. cermets, or other metal compounds, e.g. oxynitrides, sulfides
    • CCHEMISTRY; METALLURGY
    • C22METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22CALLOYS
    • C22C29/00Alloys based on carbides, oxides, nitrides, borides, or silicides, e.g. cermets, or other metal compounds, e.g. oxynitrides, sulfides
    • C22C29/005Alloys based on carbides, oxides, nitrides, borides, or silicides, e.g. cermets, or other metal compounds, e.g. oxynitrides, sulfides comprising a particular metallic binder
    • CCHEMISTRY; METALLURGY
    • C22METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22CALLOYS
    • C22C33/00Making ferrous alloys
    • C22C33/02Making ferrous alloys by powder metallurgy
    • C22C33/0257Making ferrous alloys by powder metallurgy characterised by the range of the alloying elements

Abstract

The invention relates to the use of binder alloy powders containing molybdenum to produce sintered hard metals based on tungsten carbide, wherein the binder alloy powder used has an FSSS value of 0.5 to 3 μm measured with the "Fisher Sub-Sieve Sizer" device according to the ASTM B330 standard, and comprises 0.1 to 65 wt % of iron, 0.1 to 99.9 wt % of cobalt, and 0.1 to 99.9 wt % of nickel, and contains 0.1 to 10 wt % of Mo in alloyed form.

Description

 metal powder

Field of the invention

The present invention relates to the use of molybdenum-containing binder alloy powders for the production of tungsten carbide-based sintered hard metals. Cemented carbide is a sintered composite of hardeners, such as carbides, and a continuous binder alloy. Sintered cemented carbides are used in a wide variety of ways and are used to process virtually all known materials such as wood, metal, stone and composite materials such as glass / epoxy resin, chipboard, concrete or asphalt / concrete. In this case, by cutting, forming and friction processes locally limited temperatures up to about 1000 0 C. In other cases, forming operations of metallic workpieces are carried out at high temperatures, such as forging, wire drawing or rolling. In all cases, the cemented carbide tool may be subject to oxidation, corrosion and diffusive as well as adhesive wear, while being subject to high mechanical stress, which may lead to deformation of the cemented carbide tool. The term "adhesive wear" is understood to mean the phenomenon that occurs when two bodies touch each other and at least briefly enter into a welded and firm connection, which is released again by an external force, the material of one body adhering to the other body The term "diffusive wear" is understood to mean the phenomenon that occurs when two materials are in contact with one another and a component diffuses from one material into the other material, so that a crater is formed in the first material.

Background of the prior art

WO 2007/057533 (Eurotungstene Poudres) describes alloy powders based on FeCoCu with 15 to 35% Cu and 1, 9 to 8.5% Mo for the production of diamond tools. The FSSS value is typically 3 μm. These powders are not suitable for use in the field of hard metals because of the high FSSS value, measured according to the granulometric method of Fisher or according to the standard ISO 10070, and because of the content of Cu of more than 500 ppm. The molybdenum is added as a water-soluble ammonium salt to the oxide before it is reduced with hydrogen to the metal powder. EP 1 492 897 B1 (Umicore) describes FeCoNiMoWCuSn-based alloy powder for the production of diamond tools, the sum of the contents of Cu and Sn being in the range of 5 to 45%. However, both elements are detrimental to hard metals, since Cu is sweated out during sintering, and Sn leads to pore formation. These alloy powders are therefore not suitable for the production of hard metals.

EP 0 865 511 B9 (Umicore) describes FeCoNi-based alloy powders having a maximum FSSS value of 8 μm, which may contain up to 15% Mo, but which is at least partially present as an oxide. These powders also contain between 10 and 80% Fe, up to 40% Co and up to 60% Ni, and are used to make diamond tools. In addition, similar powders, but with Co and Ni are ever described up to 30%.

Due to the content of copper, it is also unsuitable for alloy powder according to WO 98/49 361 (Umicore), EP 1 042 523 B1 (Eurotungstene Poudres) and KR 062 925.

EP 1 043 411 B1 describes carbide-Co (W, Mo) composite powders, wherein the binder alloy is produced by pyrolysis of organic precursor compounds. The onset of alloying of cobalt with Mo and / or W avoids the appearance of porosity, as occurs with the addition of metals. However, the method described is disadvantageous in comparison with the use of alloyed powders according to the invention, since the carbon content of the composite powder changes during the pyrolysis of the organic precursor compounds (carbon precipitation or removal by methane formation), so that the carbon content must be analyzed and adjusted again prior to sintering , It also remains unclear in which form Mo or W are present after sintering, since neither comparative experiments nor information on the alloy state of Mo and W before sintering nor values for magnetic saturation. The described method produces a fixed formulation with regard to the content and the composition of the carbide and binder alloy phase and is therefore too inflexible in practice since an uncomplicated and rapid change of the formulation is cumbersome depending on the application of the cemented carbide to be produced.

Furthermore, FeCoMo-based alloy powders having an FSSS value <8 μm and a specific surface area of greater than 0.5 m 2 / g have become known (DE 10 2006 057 004 A1), which are suitable for the production of carbon-free cutting steels via a serve powder metallurgical process. These may optionally contain up to 10% or 25% Ni, but most preferably do not contain nickel beyond the level of unavoidable impurities. They consist preferably of 20 to 90% Fe, up to 65% Co and 3 to 60% Mo. Since pure FeCo alloys without alloying of Ni are not suitable for hard metals because of their brittleness and poor corrosion and oxidation resistance these alloy powders suggested no solution to the problem. In addition, the preferred range is high Mo contents, and use for producing liquid phase sintered carbonaceous hard metals having a hard phase as a hard carrier such as carbides is not described.

Object of the invention

It is known that the metal cobalt, when used as the sole binding metal, in particular for tungsten carbide, brings with it a health hazard. It is therefore an object of the present invention to provide an additional

Alloy element and its provision for the production of sintered cemented carbide materials, which allows the use of FeNi and FeCoNi binders instead of Co at high working temperatures of 400 to 800 0 C, without the disadvantages such as binder lakes, the lack of interpretability of magnetic Saturation or an unknown portion of the element in question in the binder phase, wherein the element in question leads to an increase in the hot hardness in the range 400 to 800 0 C. On the other hand, the content of the element in question should be as low as possible and distributed as much as possible to improve its effectiveness.

The object is achieved by the use of a molybdenum-containing binder alloy powder for the production of tungsten carbide-based sintered hard metals, characterized in that a) the binder alloy powder used has an FSSS value of 0.5 to 3 μm determined according to ASTM B 330, and b) the binder alloy powder used contains iron in an amount of 0.1 to 65% by weight, cobalt in an amount of 0.1 to 99.9% by weight, and nickel in an amount of 0.1 to 99.9% by weight. %, and c) the binder alloy powder used contains 0.1 to 10% by weight of alloyed or pre-alloyed Mo. Advantageously, the molybdenum is completely in metallic form. The binder alloy powder used contains at least 10 wt .-% nickel, based on the total binder alloy.

The binder alloy powder used contains at most 20% by weight, in particular at most 10% by weight, of tungsten, based on the total binder alloy.

At least one component of the binder alloy is present as a powdered alloy of at least one metal with molybdenum, and each of the remaining components of the binder alloy are present as elements or alloys, each containing no molybdenum, i. a powder mixture is used comprising at least one alloyed or pre-alloyed molybdenum-containing alloy powder on the one hand with at least one alloyed or pre-alloyed alloy powder or element powder on the other hand, the latter containing molybdenum only in the region of unavoidable impurities. The molybdenum-containing binder alloy powder is used according to the invention for the production of sintered hard metals, wherein the sintering takes place in the form of a liquid phase sintering.

The molybdenum-containing binder alloy powder according to the invention may contain up to 30 percent by weight of organic additives.

Brief description of the drawings

FIG. 1 shows the course of the hot curing from Example 1 using FeCoNi binder (symbol triangle, solid line denotes the "low-carbon" variant, dashed line the "high carbon" variant) in comparison to the hot hardness of the cemented carbide from example 2 with cobalt Binder (symbol on the top standing square).

FIG. 2 shows the course of the hard hardness of hard metals from example 3 (FeCoNi binder, Mo used as element powder, symbol circles, 1% Mo = dashed line, 3% Mo = solid line) compared to that of example 4 (FeCoNi binder alloyed with Mo, symbol lying square) and to example 2 (cobalt standing as a Binder symbol on the top).

Preferred embodiments of the invention

This object is achieved by using an iron, cobalt or nickel-containing binder metal powder, which iron in an amount of 0.1 to 65 wt .-%, cobalt in an amount of 0.1 to 99.9 wt .-% and nickel in an amount of 0.1 to 99.9 wt .-% comprises.

The binder alloy powder used also contains 0.1 to 10 wt .-% of molybdenum, based on the total binder metal powder, in alloyed form. The binder alloy powder used preferably contains 0.10% by weight to 3% by weight of molybdenum, particularly preferably 0.5% by weight to 2% by weight of molybdenum, very particularly preferably 0.5% by weight to 1.7% by weight of molybdenum, in each case based on the total binder metal powder.

The binder alloy powder used has an FSSS value measured with the "Fisher Sub Siever Sizer" device according to the ASTM B330 standard of 0.5 to 3 μm, and preferably of in the range of 0.8 to 2 μm, in particular 1 up to 2 μm.

Preferably, the elements Mn and Cr are each contained in contents of less than 1%. Preferably, the binding alloy powder used contains the molybdenum completely in non-oxidic form or completely in alloyed metallic form.

Preferably, the binder alloy powder used contains at least 20 wt .-% nickel, based on the total binder alloy. The binder alloy powder used preferably contains at most 20% by weight of tungsten, preferably at most 10% by weight of tungsten, based on the total binder alloy. In particular, the preferred alloy powder is substantially free of tungsten, and has a tungsten content of less than 1 percent by weight.

Preferably, in the binder alloy powder used, at least one constituent of the binder alloy is used as a powdery alloy of at least one metal with molybdenum, and the respective remaining constituents of the binder alloy are used as elements or alloys respectively containing no molybdenum.

According to the invention, the sintering of the binder alloy powder takes place together with the hard materials as liquid phase sintering. This means that the occurrence and disappearance of a liquid, metallic phase is due solely to the change in the applied temperature, and the hard materials in the Reduce binder alloy and thereby increase its particle size (Ostwald ripening). This is in contrast to the solid phase sintering, where either no melt does not occur or a possible, intermediately occurring melt due to temporary, local composition changes, but any existing hard materials, such as diamonds, do not dissolve in the melt with increasing their particle size.

Description of the invention

The hard metals produced by the process of the invention require for their intended use of sufficient stability with respect to the plastic deformability and the temperature-dependent creep behavior. The creep of a material, such as plastic deformation, is a major failure mechanism of a material and must be avoided at all costs. Deformation mechanisms are subject to the known laws of time-dependent creep, with the creep rate not only depending on the load but also strongly on the temperature. In addition, the prevailing creep mechanism changes - activated by the temperature. With hard metals it is known that at temperatures up to about 800 0 C, the creep rate is determined mainly by deformation of the metallic binder phase, above about 800 0 C, the binder phase is so soft that it is virtually meaningless for the creep resistance, ie for temperatures of more than 800 ° C., the structural strength of the hard material phase becomes decisive. This load-bearing capacity, in turn, depends on the particle shape and size distribution of the hard material phase as well as on the proportion of heat-resistant, cubic carbides. Therefore, all cemented carbide materials used for cutting steels contain, in addition to WC, shares of cubic carbides such as TiC, TaC, NbC, VC, ZrC or mixed carbides such as TaNbC, WTiC or WVC.

Since the temperature-dependent determination of the creep behavior at high temperatures is experimentally very expensive, the determination of the hot hardness is used as a substitute. The hardness of a material is indirectly a measure of its plastic deformability. The central consideration is that in the emergence of the hardness impression plastic deformation processes predominate, so that the size of the hardness impression at sufficiently high load and load duration is a measure of the plastic deformability of the material at a given pressure load. When sintering WC-based cemented carbide with Co as the binder alloy, liquid-phase sintering in the binder phase dissolves both tungsten, carbon, and small amounts of metals which form cubic carbides, such as V, Ta, Ti, and Nb. This also applies to Cr if Cr carbide is used as a so-called "grain growth brake", that is, as a grain growth inhibiting agent, for the growth of the WC entering during sintering.

Is meant by the term "liquid phase sintering" sintering at such high temperatures that the binder alloy is at least partially melts. The liquid phase during sintering of hard metals is due to the sintering temperatures, which are generally between 1100 0C and 1550 0C. The molten phase - essentially the binder metal such as cobalt, or the binder metal alloy or alloys used or used - is in equilibrium with the hard materials, the principle of the solubility product being true, which means that all the more tungsten is contained in the melt The tungsten content of the binder alloy is set by the overall ratio W: C in the cemented carbide, where in the hard material phase always W: C = 1, then in the binder metal melt different concentrations are included with a W: C ratio not equal to 1. If the tungsten: carbon When the melt in the melt reaches a critically small size, carbon deficient carbides, such as Co 3 W 3 C, the so-called eta-phases (η-phases) precipitate on cooling. These η-phases are very hard, but also very brittle and are therefore rated in hard metals as a quality defect.

In general, it has been found that the higher the chemical stability of the corresponding carbide, the lower the achievable content of a certain metal in the binder alloy. The chemical stability of the corresponding carbides is known and can be expressed in terms of the free formation enthalpy of the carbides. If one assigns these values but in the unusual representation, namely based on one mol of metal content, then the order results at 1000 0 C:

Cr 3 C 2 <Mo 2 C <WC <VC <NbC <TaC <ZrC <TiC <HfC.

It can be seen here that, as expected, chromium carbide liberates metallic chromium as the first carbide, which is found in the binding alloy. tion, but surprisingly molybdenum is already the next unstable carbide, even before tungsten. Therefore, there is the theoretical possibility of alloying a cemented carbide binder with larger contents of molybdenum without the formation of eta phases (η phases) due to a lack of carbon in the binder phase. The above series of metal carbides is also a measure of the affinity of the metal for carbon. For example, titanium competes with Cr 3 C 2 for the carbon, so that chromium is preferably present as metal and titanium as carbide. Tungsten carbide must be present as a hardness carrier in the material; Therefore, all carbides that are in the above row to the left of tungsten carbide, ie less stable than tungsten carbide with respect to the release of the metal from the corresponding carbide, are suitable for increasing the hot hardness since they can transition to the metallic binder phase without it to the formation of carbon deficient carbides, the so-called. η-phases comes.

Since the rules of the solubility product apply to the concentrations of all of the abovementioned metals in the binder, which is the greater the more unstable the carbide is, and since there is only one carbon potential in equilibrium, it also becomes clear from the order in which order the Binder excrete dissolved metals with increasing carbon supply in the form of carbides, and thus the binder to increase the hot hardness are no longer available.

The content of chromium or tungsten is very important for the high-temperature properties of the binder alloy, since these elements lead to an increase in the heat resistance and thus to an increase in the deformation resistance. Therefore, carbide grades to be used as tools (inserts), for example for turning steels, are so sintered with respect to the carbon budget that the tungsten content in the binder alloy, which generally comprises cobalt, becomes maximum without causing formation comes from eta-phases (η-phases). Even with tools for drilling or milling metal processing, which contain Cr carbide, the carbon content is adjusted so that as much as possible Cr is contained in the binder alloy. Since the magnetic saturation of the cobalt steadily decreases with increasing Cr and W content, a non-destructive examination of the alloy state via the measurement of the magnetic saturation is very easily possible, which measurement method represents the industrial standard. However, chromium, because of its anti-ferromagnetic character, makes it difficult to determine the carbon content in the cemented carbide, and thus the chromium and tungsten content, because the uniqueness of the relationship between magnetic saturation on the one hand and chromium and tungsten content on the other hand is lost. Consequently, the absence of η-phases can not be excluded only because of the measurement of the magnetic saturation,

Due to the health risks associated with the combination of WC with cobalt as a binder alloy, there is an interest in cobalt substitutes, which may be based on FeCoNi or FeNi based alloy powders. Their suitability has been proven for wearing parts and woodworking or stone working tools, but not for applications involving high temperatures. A major reason for this is the lower hot strength of the hard metals compared to cobalt with Fe (Co) Ni binder in the temperature range between 400 0 C and 800 0 C.

The hot hardness of the binder alloy can be increased by precipitation or alloying of other metals. As alloying elements, however, only those metals are suitable which do not form stable carbides, that is to say those carbides which are not more stable than tungsten carbide, and which therefore have the prerequisites for a significant solubility in the binder alloy. If, for example, the binder were to be alloyed with Ta, this would (depending on the carbon content of the cemented carbide) be practically completely present as eta phase or TaC after sintering and thus not represent a highly heat-resistant binder alloy of a high-quality hardmetal, because eta phases are in carbide because of their Brittleness not desirable because they reduce the strength.

In principle, to increase the hot hardness in particular the metals W, Mn, Cr, Mo, Re and Ru in question.

The solubility of tungsten in the binder alloy is limited by the solubility product of tungsten carbide in the binder alloy. At the limit of the formation of eta phases, two cases are to be distinguished with respect to the tungsten content: a) when the carbon content decreases and cobalt is used as the binding detail, up to 20% by weight tungsten dissolves in the cobalt binder; b) if the carbon content decreases and a FeCoNi-binder alloy is used, dissolves much less tungsten, namely only up to about 5 wt .-%, in the FeCoNi-binder alloy. Consequently, the solubility of tungsten in FeCoNi and FeNi alloys is even lower than in pure cobalt, which is one of the reasons for the low hot hardness of FeCoNi bonded hard metals.

Manganese has a comparatively very high vapor pressure, therefore it comes to the sintering of manganese-containing hard metals concentration gradients and precipitation of self-igniting Mn-metal condensates. The concentration of Mn in sintered parts is therefore not precisely adjustable, and presumably nearer the surface than in the core of the workpiece.

The metals rhenium, osmium and ruthenium are only limited available and extremely rare, but are in principle suitable. For example, rhenium is used in high temperature alloys for aircraft turbines to suppress high temperature creep of components. Ruthenium and rhenium are already being used on a small scale in special cobalt-based carbide grades.

Chromium is also suitable and has high solubility in FeNi and FeCoNi alloys, but has the disadvantage of making it difficult to interpret the magnetic saturation due to its anti-ferromagnetic character. This is disadvantageous because carbide grades for metal cutting are as close as possible to the limit for the formation of eta phases, but without having appreciable proportions thereof.

Molybdenum in the form of added molybdenum carbide (Mo 2 C, 5 wt .-% as an addition to cemented carbides with 10% Fe-based binder) has been shown (Prakash thesis) also to increase the heat hardness in FeCoNi alloys. However, since an unknown part of the Mo is carbidic, it comes to a

Mixed carbide formation between WC and the cryptomodification MoC dissolved in it, which leads to an unwanted and uncontrollable reduction of the inherent hardness of the hard material. The mixed carbide formation in the case of molybdenum can be described by the following reaction equation: Mo 2 C -> Mo (alloyed in the binder) + (W 1 Mo) C. The solubility of molybdenum is higher than tungsten in Fe and Ni containing alloys. The efficiency curve of Mo on increasing the creep resistance of pure iron at 427 ° C is much steeper than that of Cr (Trans. Amer. Inst. Min. Met. Eng. 162, (1945), 84) where from 0.5 % Chromium only a very slow increase is observed. Already 1% Mo leads to a creep resistance of 38 kpsi (262 MPa), while at 1% Cr reaches only 16 kpsi (110 MPa) and will not be exceeded even at 4% chromium 18 kpsi (124 MPa). Without plateau, the hot-hardness temperature curve of the Mn runs, but with a much lower slope. Mo is therefore the preferred element of choice to increase the hot hardness, especially of ferrous binder in sintered hard metals. L. Prakash found that just a few percent molybdenum suffice to achieve a significant effect in the hot hardness of Fe-containing hard metals (PhD thesis Leo J. Prakash, University of Karlsruhe 1979, Faculty of Mechanical Engineering, KfK 2984). However, it remains unclear what proportion of the Mo is actually in the binder, since Mo 2 C was used.

The metals, which should lead to an increase in the hot hardness of the binder, must be present in the binder and not in the hard material, so that they can lead to an increase in the hardness of the hard metal below 800 0 C. It must therefore be made arrangements that the metals are actually present in the binder metal alloy, and not in the hard material. In the case of W and Cr, it is the industrial standard to use carbides, metals or nitrides and to adjust the carbon content of the cemented carbide by formulation and measures during sintering so that the cemented carbide is at the edge of the existence area to the eta-phase (η-phase) , and the maximum possible proportion of W and Cr in the binder is. Cr is therefore generally added as Cr carbide, which disproportionates during sintering approximately according to the following equation:

Cr 3 C 2 -> Cr (alloyed in the binder) + 2 CrC (alloyed in the WC) This makes only a fraction, namely 1/3, of the Cr used in the binder effective. Similarly, the ratios for Mo 2 C are:

Mo 2 C -> Mo (alloyed in the binder) + (W 1 Mo) C.

When using Mo carbide only a maximum of about 50% in the binder alloy are therefore effective; therefore, instead of Mo 2 C elemental Mo metal powder is used. Even with the use of very finely dispersed Mo metal powder, however, it comes after sintering to areas that consist exclusively of binder alloy phase, and contain no hard material. This behavior is due to poor crushing of agglomerates of the Mo metal powder due to the high modulus of elasticity of molybdenum in the mixed grinding, and the resulting reformed agglomerates to dissolve in the molten binder alloy during liquid-phase sintering, which in turn results from dissolution of the molten binder alloy Mo particles formed in the molten binder filled pores. It comes to the formation of the so-called "binder lakes", which term refers to a specific range of the binder alloy, which is greater in terms of the dimension than the particle diameter of the hard material phase, but without containing tungsten carbide or hard particles.

These are detrimental to strength as well as local wear resistance and unacceptable. Due to the limited diffusion time, corresponding to the time during which molten binder phase is present during sintering, it is unclear whether complete dissolution of the Mo metal powder and a homogeneous alloy of the Mo in the binder alloy will even be achieved.

If the molten binder does not fill the resulting secondary pores during sintering, they are visible in the sintered body, as described in EP 1 043 411 B1, Spalteni, lines 29/30. These so-called secondary pores reduce the strength.

According to the invention, iron, cobalt or nickel-containing binder metal powders are used for the production of sintered hard metal materials, which iron in an amount of 0.1 to 65 wt .-%, cobalt in an amount of 0.1 to 99.9 wt .-% and nickel in an amount of 0.1 to 99.9 wt%. The percentages are by weight and generally refer to the binder alloy powder unless otherwise specified.

The binder alloy powder used contains 0.1 to 10 wt .-% of molybdenum, based on the total binder metal powder, in alloyed form. Preferably, the binder metal powder used contains 0.10 wt .-% to 3 wt .-% molybdenum, particularly preferably 0.5 wt .-% to 2 wt .-% molybdenum, most preferably 0.5 wt .-% bis 1, 5 wt .-% molybdenum, each based on the total binder metal powder. One too high molybdenum content leads to excessive solidification of the binder powder, so that the pressing forces in the production of the cemented carbide and the resulting sintering shrinkage are too high, too low content leads to an insufficient increase in the hot hardness.

Preferred hard materials are carbides, in particular tungsten carbide, WC. Preferred binders are alloys of iron, cobalt and nickel, in particular the combinations iron and nickel, iron and cobalt, cobalt and nickel, and iron, cobalt and nickel. Likewise, cobalt alone can be used as a binder.

Due to their physical properties, the binder metal powders alloyed with molybdenum are distinguished by good distribution behavior in the case of mixed grinding with carbides for the production of hard metal powders. The FSSS values (measured with the "Fisher Sub Siever Sizer" device according to the ASTM B330 standard) are therefore in the range of 0.5 to 3 μm, preferably in the range of 1 to 0 to 2 μm, even finer powders are self-igniting; Coarser powders no longer have a sufficient distribution behavior and lead again to so-called "binder lakes". The size distribution of Agg lomerate is in the range of 0.5 to 10 microns with the same reason. The specific surface area is preferably between 2.5 and 0.5 m 2 / g for the same reasons. The oxygen content is preferably less than 1.5%.

Preferred contents for cobalt in the binder alloy are up to 60% by weight. The preferred content of nickel in the binder alloy is in the range of 10 to 80 wt%, or 20 to 60 wt%, or 30 to 75 wt%.

It is also possible to add subsequently added organic additives. To determine the aforementioned parameters, they may need to be removed again, which is done, for example, by washing with a suitable solvent. The organic additives include waxes, agents for passivation and inhibition, corrosion protection, pressing aids. As an example, paraffin wax and polyethylene glycols come into consideration. By the organic additives also an aging of the powder is to be prevented, which brings an increase in the oxygen content with it. The additives may be contained in an amount of 30% by weight, based on the sum of binder alloy powder and additive. The Mo-containing binder powder may contain Fe, Ni and Co. Since the sinterability and the hot hardness decrease with increasing Fe content, the iron content is less than 65%, preferably less than 60%. The remainder to 100% is Mo and Co and / or Ni. Preferably, such alloys are selected in the FeCoNi system as binder alloys which are stably austenitic in the sintered cemented carbide, such as FeCoNi 30/40/30 or 40/20/40 or 20/60/20 or 25/25/50, but also FeNi 50/50 or 30/70 or 20/80, or CoNi in Councils 50/50, 70/30 or 30/70. However, it is also possible to use element powders, such as Co or Ni, alloyed with up to 10% Mo, which thus become alloy powders.

The molybdenum-containing alloy powders are preferably prepared by the following process (DE 10 2006 057 004 A1): a MoO 2 , which was comminuted to reduce the agglomerate size distribution, serves as molybdenum source. This MoO 2 is added to an oxalic acid suspension, as used according to EP 1 079 950 B1 for the preparation of FeNi or FeCoNi mixed oxalates, which are subsequently oxidatively annealed, and reduced with hydrogen to form alloy powders. The resulting alloy powders are completely reduced after reduction with hydrogen, ie it is no longer detectable by X-ray diffraction MoO 2 . It may also be reduced in agglomerate size by deagglomeration in order to improve the distribution in the mixed grinding with the carbides. The agglomerates consist of primary particles which are agglomerated together. Agglomerate size and distribution can be determined by laser diffraction and sedimentation.

Instead of MoO 2 , it is also possible to use other fine-grain Mo compounds which do not dissolve in oxalic acid, for example sulfides or carbides. These are oxidized to oxides in the calcination of the precipitated oxalate in air. During the calcination, molybdenum oxides such as MoO 3 are formed , which very quickly form mixed oxides with the Fe (Co) Ni mixed oxide due to their high vapor pressure and thereby show good transport properties, so that in the subsequent reduction with hydrogen forms a FeCoNi alloy powder with a small part of Mo is homogeneously alloyed. However, other known methods are suitable, for example, instead of oxalic acid, a precipitation with their ammonium salts, with sodium or potassium hydroxide, with formic and maleic acid. In all cases, preferably MoO 2 is used, which should be as pure as possible phase, and should contain Mo or MoO 3 or Mo 4 O 11 only in traces. MoO 2 is used because, in contrast to MoO 3, it is neither soluble in acids nor in alkali, and therefore remains completely in the alloy metal powder after the entire process. MoO 3 would dissolve in the alkali used to precipitate the Fe (Co) Ni content or in complexing organic acids; elemental Mo would be too coarse and would not fully oxidize to MoO 3 in the subsequent calcination and thus would not alloy sufficiently upon reduction with hydrogen. A fine MoO 2 having a high surface area completely oxidizes to MoO 3 (which has a high vapor pressure) upon calcination of the Fe (Co) Ni oxalate in air and forms molybdate and mixed oxides with these metal oxides through the gas phase, thereby providing a very uniform Distribution of the molybdenum is achieved, which is maintained in the subsequent reduction with hydrogen.

It is already known to use powders according to the invention which contain alloyed Mo for the production of sintered parts by means of solid phase sintering, as in the diamond tool industry, but not for the cemented carbide industry with intermediate formation of a molten phase during sintering.

However, those with Mo alloyed FeCoNi powder which contain the Mo in completely metallic form are particularly preferred. By means of X-ray diffraction, Mo oxides can no longer be detected in these powders, and consequently the oxygen present must predominantly be present on the surface of the powder. According to the invention, powders whose FSSS value is in the range from 0.5 to 3 μm are particularly suitable because the distribution during mixed grinding is better. In this case, if possible, they contain no further, oxidic metals.

Since Mo oxide reacts with carbon in carbide sintering to form CO, thus leading to a local carbon deficit and thus to local eta phases, the alloy powders described in the previous paragraph are then suitable for hard metal fabrication when providing for carbide sintering be taken that the predominantly released in the form of carbon monoxide oxygen from the sintering can escape. These powders are suitable for use in accordance with the invention if they have the physical properties preferred according to the invention, but the elements Mn, Cr, V, Al and Ti described contain at least partially oxidic form only to the extent that it is from the viewpoint of microstructural defects (pores and binder lakes). of the carbide is allowed.

According to the invention, the Mo alloyed FeCoNi or FeNi based powders may additionally be alloyed with up to 20% tungsten, for example to shift the onset of sintering shrinkage to higher temperatures or to provoke the formation of precipitates which reinforce the binder phase, however only possible with very coarse tungsten carbides.

According to the invention, the alloy powders used can occupy a wide range in the composition space FeCoNi. In the range of high Fe contents (90 to 60%), such binder alloy systems are to be found which, after sintering, have proportions of martensitic phase and therefore have a high hardness and wear resistance at room temperature. Examples are FeNi 90/10, 82/18, 85/15, FeCoNi 72/10/18, 70/15/15 and 65/25/10. However, the abovementioned alloys are distinguished by very low thermal hardness in the sintered hard metal. In the range between about 80 and 25% Fe, there is the range of austenitic binder alloys after sintering, which are characterized by a lower intrinsic hardness, but by high fatigue strength and limited plastic deformation capability. Examples are FeNi 80/20, 75/25, FeCoNi 60/20/20, 40/20/40, 25/25/50, 30/40/30, 20/60/20. In most cases, the hot hardness of carbides between 400 and 600 ° C is inferior to that of pure Co as a binder, unless alloyed with Mo or other alloying elements. Although the particularly preferred objective of the use according to the invention is the production of hard metals with better hot hardness, it is also well suited for the production of hard metals with other objectives, such as cemented carbide with molybdenum-containing corrosion-resistant binder alloy systems, which are produced today using elemental or carbidic molybdenum as described, for example, in EP 0 028 620 B2, or also bit bits for drill bits, as described in US Pat. No. 5,305,840. According to the invention, the binder alloy present after the sintering of the cemented carbide can also be obtained by using a plurality of different alloying powders and optionally elementary powders, as described in WO 2008/034903, wherein at least one of these powders is alloyed with molybdenum. The advantages of such a procedure lie in the compressibility and the control of the sintering shrinkage.

The hard metal part present after sintering and possibly the grinding or electro-eroding finish has a defined tool geometry. This may most preferably be elongated (for example, ground out of a sintered round bar), but more preferably also plate-shaped for turning or milling of materials such as metals, bricks and composite materials. In all cases, the cemented carbide tools may preferably have one or more coatings from the classes of nitrides, borides, oxides or superhard layers (e.g., diamond, cubic boron nitride). These can be applied by PVD or CVD methods or their combinations or variations and still be changed after application in their residual stress state. It may be in a preferred manner but also 'carbide parts of further and arbitrary geometry and application, such as forging tools, forming tools, countersinks, components, knives, peeling plates, rollers, stamping tools, pentagonal drill bits for soldering, mining chisel, milling cutters for milling of concrete and Asphalt, mechanical seals and any other geometry and application. The invention is explained in more detail by the following examples.

Examples

Example 1 (Comparative example, not according to the invention) 462.5 g tungsten 0.6 microns were mixed with 37.5 g of a FeCoNi alloy powder 40/20/40 (Amperesint MAP ® 6050 A, manufactured by HC Starck, Germany) in a ball mill with Mix 0.57 liters of 94% ethanol for 14 hours at 63 rpm. In this case, 5 kg of hard metal balls were used. The FeCoNi powder used had the following properties: Fe 38.8%, Co 20.22%, Ni 40.38%, O 0.71%, specific surface area 1, 63 m 2 / g, FSSS value 0.90. There were 2 approaches with different carbon contents ("high carbon" or "low carbon") produced, so that after the Sintering different carbon contents result. The results are shown in the table below.

From the resulting suspension, the ethanol was removed in vacuo off separated by distillation, and the pressed hard metal powder axially at 150 MPa and at 145 ° 0 C for 45 min sintered in vacuum. The plate-shaped hard metal pieces were ground, polished and examined for their properties. Both batches showed neither eta phases nor carbon precipitates as sinters, but smaller binder lakes. In both cases, the room temperature hardness and the hot hardness at selected temperatures up to 800 ° C. under protective gas were measured. FIG. 1 shows the results: both batches show a sharp drop in the hot hardness in the region around 600 ° C. Thus, this binder alloy for the production of carbide tools for metal cutting (turning) at higher stress compared to pure cobalt clearly inferior, since due to the low hot hardness, especially at 600 0 C, a plastic deformation of the cutting edge due to the cutting forces is expected.

Figure imgf000019_0001

Example 2 (Comparative Example, WC-Co, not according to the invention) Analogously to Example 1, a WC-Co with the same volume fraction as in Example 1 was prepared on binder phase. Since Co has a higher density than the FeCoNi 40/20/40, the weight fraction of the cobalt was 8 wt .-%, based on the total carbide. After pressing and sintering at 1420 0 C for 45 min in vacuo resulted in a perfect hard metal with a magnetic saturation of 133 G-cm 3 / g, corresponding to 82% of the theoretical magnetic saturation. The

Room temperature hardness (HV30 1597 kg / mm 2 ) and the hot hardness were determined and entered in FIG. It is evident that Co above which the carbide skeleton hot hardness largely determined, is superior to the FeCoNi binder of 350 up to 800 0 C. The KiC value (fracture toughness, determined from the crack lengths at the corners of the hardness impressions, calculated according to the formula of Shetty) of the cemented carbide Room temperature was 10.1 MPa ^ m 172 . Thus, the cobalt binder at room temperature has a better HaIIeZK 1 C ratio than the binder of Example 1 ..

Example 3 (comparative example, not according to the invention) Example 1 was repeated, but in a first batch 1% by weight and in a second 3% by weight Mo metal powder were added. (These contents are based on the Mo content of the binder alloy phase). The deagglomerated molybdenum metal powder had the following properties: FSSS value 1, 09, O content: 0.36 wt%. The grain distribution is determined by the following parameters: D 50 3.2 μm, Dg 0 6.4 μm. The carbon content was chosen so that according to the experience of Example 1 in the sintered hard metal neither eta-phases nor carbon precipitations are to be expected. For the Mo addition, no additional carbon was included so that the molybdenum is as completely as possible in metallic form in the binder alloy. Therefore, the carbon contents of the recipe were 5.94% and 5.94%, respectively (3% by weight of Mo based on the binder). The results after sintering at 1420 ° C. are reproduced in the following table. The thermal hardnesses were determined as before and are represented by circles in FIG. 2:

Figure imgf000020_0001

Surprisingly, eta phases did not occur at 1 or at 3% by weight of molybdenum, but on the contrary at 1% by weight of molybdenum, even a carbon porosity. Surprisingly, the hardness increases without the K 1 C strength being reduced in comparison to Example 1, so that a combination of properties is obtained at room temperature which is equal to that of the Co-bonded hard metal and clearly superior to that of the purely FeCoNi-bonded hard metals is. Surprisingly, 1% by weight of molybdenum in the binder is already sufficient; with 3 wt .-% molybdenum no strong change of K 1 C and the hardness compared to 1% Mo is observed more. The effect of molybdenum alloyed in the binder is thus not only an increase in the inherent hardness of the binder, but also in a simultaneous Increase in crack resistance. In this respect, the behavior is different than for alloyed W: here, too, an increase in the inherent hardness of the binder is found, but a simultaneous decrease of K 1 C-We (IeS, this in both Co-based carbides as well as those based on FeCoNi See example 1.

However, many binder lakes appear, which is evidence of the dissolution of Mo in the binder, which in turn fills the resulting pore volume. However, these binder lakes are unacceptable in a cemented carbide. The comparison of the hot hardness with those of Example 2 is shown in FIG. 2. The hot curing at all temperatures up to 800 ° C. are surprisingly even lower than those of Example 1.

Example 4 (according to the invention)

Example 1 was repeated by using the FeCoNi binder alloy alloyed according to the process described in DE 10 2006 057 004 A1 and containing 1.5% by weight of Mo. The powder was then deagglomerated. The analyzed properties of this powder were: Fe 38.23% by weight, Co 19.96% by weight, Ni 39.10% by weight, Mo 1, 55% by weight, O 0.8565% by weight. %, FSSS value: 1.21, specific surface 2.17 m 2 / g, D 50 3.46 μm, D 90 5.84 μm. It was by X-ray diffraction even in long exposure no MoO 2 in its characteristic

More evidence of diffraction angles. 37.5 g of this powder with 462.5 g of WC were used to produce a cemented carbide. The cemented carbide mixture had a carbon content of 5.92% by weight, which was adjusted by adding 1.14 g of carbon black. The compacts were sintered in both an open and a closed crucible. This variation has an effect on the

Carbon content of the cemented carbide after sintering. The properties of the cemented carbide sintered at 1420 ° C. were as follows:

Figure imgf000021_0001
The cemented carbide from the open sintering is located at the low carbon end of the two phase region because it is characterized by a very low magnetic saturation compared to Example 1. However, eta phases were not detectable. Due to the maximum possible concentration of Mo in the binder, an enormous strengthening of the binder alloy is achieved, which is expressed by a simultaneous increase in hardness and fracture toughness. The cemented carbide from the closed sintering is also in the 2-phase region in terms of carbon content, but contains more carbon, which is indicated by the high magnetic saturation. Because of the higher carbon supply apparently more Mo is present as carbide and therefore is not present in the binder, the fracture toughness - which is largely determined by the binder - drops very much to the level of the "high carbon" variant of Example 1. This example confirms the theoretical considerations made in the description.

Other pellets were prepared and sintered at 1420 0 C in a vacuum, but was carried out towards the end of the sintering at the end temperature, an application of argon at 40 bar pressure. It was cooled under pressure. There were hard metal pieces with a hardness of 1643 HV30 obtained, a crack resistance of 8.2 MPa-m 1/2 and a magnetic saturation of 123 G-cm 3 / g. On the hard metal pieces on another hardness testing machine both the room temperature and the

Hot hardness determined as a function of temperature. The evaluation of the determination of the room temperature and hot hardness shows the figure 2, represented by squares, for comparison, the curves of Examples 2 and 3 is plotted: the decrease in the hot hardness at 600 0 C compared to a cobalt-bonded carbide is for the hard metals of example 4 compared to those of Example 2 significantly reduced. The hot hardness is now above that of the cemented carbide prepared from the non-Mo alloyed binder alloy powders (Example 3). (Due to the other hardness testing machine, there is a discrepancy in the room temperature hardness).

It can be seen that the use according to the invention of a molybdenum (pre-alloyed) binder powder makes it possible to produce a perfect hard metal without binder lakes and with a course of heat history virtually as with a cobalt binder. In particular, the drop in the hot hardness is virtually eliminated by 600 0 C. In addition, with an appropriate setting of the carbon budget compared to Example 1, a tremendous improvement in both room temperature and solidity results an increase in hardness compared to Example 1, which also offers advantages for applications at or near room temperature. In addition, an improvement in the corrosion resistance is expected over Example 1, because the corrosive attack on hard metals is generally carried out via the binder phase.

The principle of improving the properties of hard metals by alloyed molybdenum in the binder is applicable not only to the binder described FeCoNi 40/20/40, but also to pure cobalt as well as pure Ni as a carbide binder on CoNi and FeNi alloys as well as others FeCoNi alloys.

Claims

claims
1. Use of a molybdenum-containing binder alloy powder for the production of liquid-phase sintered hard metals
Tungsten carbide base, characterized in that a) the binding alloy powder used has an FSSS value of 0.5 to 3 [mu] m, determined according to ASTM B 330, and b) the binding alloy powder used is iron in an amount of less than 60 Wt .-%, and cobalt in an amount of up to 60
Wt .-% and nickel in an amount of 20 to 60 wt .-% comprises, and c) the binder alloy powder used 0.1 to 10 wt .-% Mo in alloyed or pre-alloyed form.
2. Use according to claim 1, wherein the molybdenum is completely in metallic form.
3. Use according to one or more of claims 1 to 2, wherein the binding alloy powder used contains at least 10 wt .-% nickel, based on the total binder alloy.
4. Use according to one or more of claims 1 to 3, wherein the binder alloy powder used contains at most 20 wt .-% tungsten, based on the total binder alloy.
5. Use according to one or more of claims 1 to 4, wherein at least one component of the binder alloy is used as a powdered alloy of at least one metal with molybdenum, and the remaining constituents of the binder alloy are used as elements or alloys, each of which is not molybdenum contain.
6. Use according to one or more of claims 1 to 5 for the production of sintered hard metals, wherein the sintering takes place in the form of a liquid phase sintering.
7. Use according to one or more of claims 1 to 6, characterized in that it contains up to 30 weight percent of one or more organic additives.
8. Use according to one or more of claims 1 to 7, wherein the binder alloy powder used contains at most 10 wt .-% tungsten, based on the total binder alloy.
9. Pre-alloyed powder containing 0.1 to 65% by weight of iron, 0.1 to 60% by weight of cobalt, 10 to 80% by weight of nickel and 0.1 to 20% by weight of molybdenum in metallic form, wherein the FSSS value according to ASTM B 330 is a maximum of 3 μm and the remaining constituents of the powder are unavoidable impurities.
10. Pre-alloyed powder according to claim 9, which additionally contains up to 10 wt .-% tungsten in alloyed or pre-alloyed form.
11. Pre-alloyed powder according to one or more of claims 9 to 10, characterized in that it contains 0.1 to 20 wt .-% molybdenum.
12. Pre-alloyed powder according to one or more of claims 9 to 11, characterized in that it contains 0.1 to 65 wt .-% iron and 10 to 60 wt .-% nickel.
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