Patent Application
20110286877
The present invention relates to the use of molybdenum-containing binder alloy powders for producing sintered hard metals based on tungsten carbide.
Hard metal is a sintered composite composed of hardness-imparting materials such as carbides and a continuous binder alloy. Sintered hard metals are widely used and are employed for machining of virtually all known materials such as wood, metal, stone and composites such as glass-epoxy resin, chipboards, concrete or asphalt-concrete. Localized temperatures up to above 1000° C. occur here as a result of cutting, deformation and friction processes. In other cases, deformation of metallic workpieces is carried out at high temperatures, for example, in forging, wire drawing or rolling. In all cases, the hard metal tool can be subject to oxidation, corrosion, and diffusive and adhesive wear, and is at the same time under high mechanical stress, which can lead to deformation of the hard metal tool.
The term “adhesive wear” refers to any phenomenon which occurs when two bodies are in contact with one another and at least briefly form a strong welded bond, with the material of one body adhering to the other body, which is released again by means of an external force. The term “diffusive wear” refers to any phenomenon which 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.
WO 2007/057533 describes alloy powders based on FeCoCu and containing from 15 to 35% of Cu and from 1.9 to 8.5% of Mo for producing 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 by the granulometric method of Fisher or in accordance with the standard ISO 10070, and because of the Cu content of over 500 ppm. The molybdenum is added as water-soluble ammonium salt to the oxide before the latter is reduced by means of hydrogen to the metal powder.
EP 1 492 897 B1 describes alloy powders based on FeCoNiMoWCuSn for producing diamond tools, where the sum of the contents of Cu and Sn is in the range of from 5 to 45%. Both elements are, however, detrimental to hard metals since Cu “sweats out” during sintering and Sn leads to pore formation. These alloy powders are therefore not suitable for producing hard metals.
EP 0 865 511 B9 describes alloy powders which are based on FeCoNi and have an FSSS value of not more than 8 μm and can contain up to 15% of Mo, although this is at least partly present as an oxide. These powders furthermore contain from 10 to 80% of Fe, up to 40% of Co and up to 60% of Ni, and are used for producing diamond tools. Powders which are similar but contain up to 30% of Co and up to 30% of Ni are also described.
Alloy powders as described in WO 98/49 361, EP 1 042 523 B1 and KR 062 925 are also not suitable because of the copper content.
EP 1 043 411 B1 describes carbide-Co—(W,Mo) composite powders in which the binder alloy is produced by the pyrolysis of organic precursor compounds. The formation of an alloy of cobalt with Mo and/or W avoids the occurrence of porosity as occurs on addition of metals. However, the process described has the that the carbon content of the composite powder changes (carbon deposition or removal by methane formation) during the pyrolysis of the organic precursor compounds, so that the carbon content must again be analysed and adjusted before sintering. The form in which the Mo or W is present after sintering also remains unclear since neither comparative experiments nor indications of the alloying state of Mo and W before sintering nor values for the magnetic saturation are given. The process described produces a fixed formulation in respect of the content and the composition of the carbide and binder alloy phases and is therefore too inflexible in practice since an uncomplicated and quick change of the formulation depending on the use of the hard metal produced is awkward.
DE 10 2006 057 004 A1 describes alloy powders based on FeCoMo having an FSSS value of <8 μm and a specific surface area of greater than 0.5 m2/g which are employed for producing carbon-free high-speed steels via a powder-metallurgical process. These can optionally contain up to 10% or 25% of Ni but particularly advantageously do not contain any nickel beyond the level of unavoidable contamination. They preferably comprise from 20 to 90% of Fe, up to 65% of Co and from 3 to 60% of Mo. Pure FeCo alloys without additional alloying of Ni are, however, not suitable for hard metals because of their brittleness and poor corrosion and oxidation resistance. The preferred range also includes high Mo contents, and a use for producing liquid-phase-sintered, carbon-containing hard metals having a hard material phase as a hardness imparter, such as carbides, is not described.
Metal cobalt represents a health hazard when used as sole binder metal, for example, for tungsten carbide. An aspect of the present invention is therefore to provide an additional alloying element for the production of sintered hard metals which allows the use of FeNi and FeCoNi binders in place of Co at high working temperatures of from 400 to 800° C. without suffering disadvantages such as binder lakes, the lack of interpretability of the magnetic saturation or an unknown proportion of the element concerned in the binder phase, with the element concerned leading to an increase in the hot hardness in the range from 400 to 800° C. Alternative aspects of the present invention are that the content of the element concerned should be as low as possible and, in order to improve the effectiveness, should be distributed as well as possible.
In an embodiment, the present invention provides a process of using a molybdenum-containing binder alloy powder to produce a sintered hard metal based on a tungsten carbide which includes providing a molybdenum-containing binder alloy powder with a FSSS value as determined in accordance with an ASTM B 330 standard of from 0.5 to 3 μm and comprising from 0.1 to 10% by weight of a molybdenum in at least one of an alloyed form and a prealloyed form, less than 60% by weight of an iron, up to 60% by weight of a cobalt, and from 10 to 60% by weight of a nickel. The molybdenum-containing binder alloy powder is incorporated into a hard metal. The hard metal is sintered so as to provide the liquid-phase-sintered hard metal based on a tungsten carbide.
The present invention is described in greater detail below on the basis of embodiments and of the drawings in which:
In an embodiment of the present invention, the molybdenum can, for example, be present entirely in metallic form. The binder alloy powder used can contain at least 10% by weight of nickel, based on the total binder alloy.
The binder alloy powder used can contain not more than 20% by weight, for example, not more than 10% by weight, of tungsten, based on the total binder alloy.
In an embodiment of the present invention, at least one constituent of the binder alloy can be present as a pulverulent alloy of at least one metal with molybdenum, and the remaining constituents of the binder alloy can be present as elements or alloys which each do not contain any molybdenum. Use can, for example, be made of a powder mixture of at least one alloyed or prealloyed molybdenum-containing alloy powder with at least one alloyed or prealloyed alloy powder or element powder, with the latter containing molybdenum only in the range of unavoidable contamination.
The molybdenum-containing binder alloy powder can, according to the present invention, be used to produce sintered hard metals, with sintering being carried out in the form of liquid-phase sintering.
The molybdenum-containing binder alloy powder can, according to the present invention, contain up to 30 percent by weight of an organic additive.
In an embodiment, the present invention provides the use of an iron-, cobalt- or nickel-containing binder metal powder comprising iron in an amount of from 0.1 to 65% by weight, cobalt in an amount of from 0.1 to 99.9% by weight and nickel in an amount of from 0.1 to 99.9% by weight.
The binder alloy powder used can additionally contain from 0.1 to 10% by weight of molybdenum, based on the total binder metal powder, in alloyed form. The binder alloy powder used can, for example, contain from 0.10% by weight to 3% by weight of molybdenum, for example, from 0.5% by weight to 2% by weight of molybdenum, or for example, from 0.5% by weight to 1.7% by weight of molybdenum, in each case based on the total binder metal powder.
In an embodiment of the present invention, the binder alloy powder used has an FSSS value measured using a “Fisher Sub Sieve Sizer” in accordance with the standard ASTM B330 of from 0.5 to 3 μm, for example, in the range of from 0.8 to 2 μm, or for example, from 1 to 2 μm.
In an embodiment of the present invention, the elements Mn and Cr can, for example, be present in contents of less than 1%. The binder alloy powder used can, for example, contain the molybdenum completely in nonoxidized form or completely in alloyed metallic form.
The binder alloy powder used can, for example, contain at least 20% by weight of nickel, based on the total binder alloy. The binder alloy powder used can, for example, contain not more than 20% by weight of tungsten, for example, not more than 10% by weight of tungsten, based on the total binder alloy. The alloy powder can, for example, be virtually free of tungsten and, for example, can have a tungsten content of less than 1 percent by weight.
In the binder alloy powder used, at least one constituent of the binder alloy as a pulverulent alloy of at least one metal can, for example, be introduced with molybdenum and the remaining constituents of the binder alloy can be introduced as elements or alloys which do not contain any molybdenum.
In an embodiment of the present invention, sintering of the binder alloy powder together with the hard materials can occur as liquid-phase sintering. This means that the appearance and disappearance of a liquid, metallic phase is due solely to the change in the temperature employed and the hard materials are dissolved and reprecipitated in the binder alloy and thus undergo an increase in particle size (Ostwald ripening). This is in contrast to solid-state sintering in which either no melt is formed or any melt formed temporarily is due to transient, local changes in composition but any hard materials present, such as diamonds, are not dissolved and reprecipitated in the melt so as to undergo an increase in particle size.
The hard metals produced by the process of the present invention need sufficient stability in respect of plastic deformability and the temperature-dependent creep behaviour in order to be used for their intended purpose. Creep of a material, for example, plastic deformation, is a failure mechanism for a material and should be avoided. The mechanisms of deformation are subject to the known time laws of load-dependent creep, with the creep rate being dependent not only on the load but also to a great extent on the temperature. The creep mechanism prevailing in each case also changes as a function of temperature. In the case of hard metals, the creep rate at temperatures up to about 800° C. is determined mainly by deformation of the metallic binder phase, while above about 800° C. the binder phase is so soft that it is of virtually no significance for the creep resistance. At temperatures above 800° C., for example, the load-bearing strength of the hard material phase is the determining factor. This load-bearing capacity depends in turn on the particle shape and size distribution of the hard material phase and on the proportion of heat-resistant, cubic carbides. For this reason, all hard metals used for the cutting of steels contain not only WC but also proportions 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 behaviour at high temperatures is experimentally very difficult, the determination of the hot hardness is employed instead. The hardness of a material is an indirect measure of its plastic deformability. The central idea is that plastic deformation processes predominate in the formation of the hardness indentation, so that the size of the hardness indentation at sufficiently high loading and loading duration is a measure of the plastic deformability of the material at a given compressed load.
During sintering, hard metals based on WC with Co as the binder alloy, tungsten, carbon and also small amounts of metals which form cubic carbides, such as V, Ta, Ti and Nb, dissolve in the binder phase during liquid-phase sintering. This also applies to Cr if Cr carbide is used as a “grain growth inhibitor,” for example, as a material which inhibits grain growth for the microstructural growth of WC which occurs during sintering.
The term “liquid-phase sintering” refers to sintering at temperatures which are sufficiently high for the binder alloy to at least partly melt. The liquid phase during sintering of hard metals is a consequence of the sintering temperatures, which are generally in the range of from 1100° C. to 1550° C. The molten phase, essentially the binder metal such as cobalt or the binder metal alloy or alloys used, is in equilibrium with the hard materials, with the principle of the solubility product applied. This means that the more tungsten present in the melt, the less carbon is dissolved in the melt, and vice versa. The tungsten content of the binder alloy is set via the overall W:C ratio in the hard metal, with W:C=1 always applying in the hard material phase and different concentrations with a W:C ratio which is not equal to one then being present in the binder metal melt. When the tungsten:carbon ratio in the melt reaches a critically low value, carbon-deficient carbides such as Co3W3C, known as eta phases (η phases), precipitate on cooling. These η phases are very hard but also very brittle and are therefore regarded as a quality defect in hard metals.
It has generally been found that the achievable content of a particular metal in the binder alloy is lower, the higher the chemical stability of the corresponding carbide. The chemical stability of the corresponding carbides is known and can be expressed in the form of the free enthalpy of formation of the carbides. If these values are ordered in the unconventional representation, namely based on one mole of metal content, then the order at 1000° C. is:
Cr3C2<Mo2C<WC<VC<NbC<TaC<ZrC<TiC<HfC.
It can be seen here that, as expected, chromium carbide as first the carbide liberates metallic chromium, which dissolves in the binder alloy, at an increasing carbon deficiency. Molybdenum is surprisingly the next most unstable carbide even before tungsten. It is therefore theoretically possible to alloy a hard metal binder with relatively large contents of molybdenum without the formation of eta phases (η phases) occurring as a result of a deficiency of carbon in the binder phase. The above series of the metal carbides is also a measure of the affinity of the metal for carbon. For example, titanium competes with Cr3C2 for carbon, so that chromium can, for example, be present as a metal and titanium can, for example, be present as a carbide. Tungsten carbide should be present as a hardness imparter in the material; all carbides which are to the left of tungsten carbide in the above series, such that they are less stable than tungsten carbide in respect of liberation of the metal from the corresponding carbide, are therefore suitable for increasing the hot hardness since they can go over into the metallic binder phase without formation of carbon-deficient carbides, that is to say without “η phases,” occurring.
Since the concentrations of all the abovementioned metals in the binder are governed by the rules of the solubility product, which is greater the more unstable the carbide, and since there is only one carbon potential in the equilibrium, the order also indicates the order in which the metals dissolved in the binder precipitate in the form of carbides with increasing carbon availability and are therefore no longer available to the binder to increase the hot hardness.
The content of chromium or tungsten is important for the high-temperature properties of the binder alloy since these elements lead to an increase in the hot hardness and thus to an increase in the deformation resistance. For this reason, types of hard metal which are to be used as tools (cutting inserts), for example, for turning steels, are sintered at such a carbon balance that the tungsten content of the binder alloy, which generally comprises cobalt, is maximized without formation of eta phases (η phases) occurring. In the case of tools containing Cr carbide for metal machining by drilling and milling, the carbon content is also set so that as much Cr as possible is present in the binder alloy. Since the magnetic saturation of cobalt decreases continuously with increasing Cr and W content, nondestructive testing of the state of the alloy can be carried out very simply by measuring the magnetic saturation. This is the standard method of measurement in industry.
Due to its antiferromagnetic character, however, chromium interferes in the determination of the carbon content in the hard metal and thus the determination of the content of chromium and tungsten because the relationship between magnetic saturation and content of chromium and tungsten is no longer unambiguous. Consequently, the absence of η phases cannot be ruled out merely on the basis of a measurement of the magnetic saturation.
Due to the health hazards associated with the combination of WC with cobalt as binder alloy, there is an interest in replacements for cobalt for which alloy powders based on FeCoNi or FeNi are possibilities. Although their suitability for wear parts and tools for working wood or stone has been demonstrated, it has not been demonstrated for applications associated with high temperatures. A significant reason for this is the lower hot hardness of the hard metals having a Fe(Co)Ni binder in the temperature range from 400° C. to 800° C. compared to cobalt.
The hot hardness of the binder alloy can be increased by means of precipitates or alloying-in of other metals. However, possible alloying elements are only metals which do not form stable carbides, for example, carbides which are not more stable than tungsten carbide, and therefore meet the prerequisites for appreciable solubility in the binder alloy. If, for example, Ta were to be alloyed into the binder, this would (depending on the carbon content of the hard metal) be virtually entirely present as an eta phase or as TaC after sintering and would thus not represent a high-hot-strength binder alloy of a high-quality hard metal since eta phases are undesirable in the hard metal because of their brittleness, leading to a decrease in the strength.
In principle, the metals W, Mn, Cr, Mo, Re and Ru are the main possibilities for increasing the hot hardness.
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 formation of eta phases, a distinction can be made between two cases in respect of the tungsten content: a) when the carbon content decreases and cobalt is used as binder metal, up to 20% by weight of tungsten dissolves in the cobalt binder; b) when the carbon content decreases and a FeCoNi binder alloy is used, significantly less tungsten, namely only up to about 5% by weight, dissolves in the FeCoNi binder alloy. Consequently, the solubility of tungsten in FeCoNi and FeNi alloys is even lower than in pure cobalt, which is a reason for the low hot hardness of hard metals bonded by means of FeCoNi.
Manganese has a comparatively very high vapour pressure, and for this reason concentration gradients and precipitates of pyrophoric Mn-metal condensates are obtained on sintering of manganese-containing hard metals. The concentration of Mn in sintered parts can therefore not be precisely set and is presumably lower close to the surface than in the core of the workpiece.
The metals rhenium, osmium and ruthenium have limited availability and are extremely rare, but are in principle suitable. Rhenium is, for example, used in high-hot-strength alloys for aircraft turbines in order to suppress the high-temperature creep of components. Ruthenium and rhenium are used commercially to a limited extent in special grades of hard metal based on cobalt.
Chromium is likewise suitable and has a high solubility in FeNi and FeCoNi alloys but has the disadvantage that owing to its antiferromagnetic character, it makes the interpretation of the magnetic saturation difficult. This is a disadvantage because hard metals for cutting metal machining are very close to the limit for formation of eta phases, but without appreciable amounts of the latter being present.
Molybdenum in the form of added molybdenum carbide (Mo2C, 5% by weight as additive to hard metals containing 10% of Fe-based binder) has also been shown (thesis by Prakash) to lead to an increase in hot hardness in FeCoNi alloys. However, since an unknown part of the Mo is present in carbidic form, formation of a mixed carbide between WC and the cryptomodification MoC dissolved therein occur, which leads to an unwanted and uncontrollable reduction of the intrinsic strength of the hard material. The mixed carbide formation in the case of molybdenum can be described by the reaction equation:
Mo2C=>Mo (alloyed in the binder)+(W,Mo)C.
The solubility of molybdenum in Fe- and Ni-containing alloys is higher than that of tungsten. The curve for the effectiveness of Mo in increasing the creep resistance of pure iron at 427° C. is very much steeper than that for Cr (Trans. Amer. Inst. Min. Met. Eng. 162, (1945), 84), with only a very slow increase being observed above 0.5% of chromium. Even 1% of Mo leads to a creep resistance of 38 kpsi (262 MPa), while 1% of Cr gives only 16 kpsi (110 MPa) and even 4% of chromium does not give a value exceeding 18 kpsi (124 MPa). The hot hardness-temperature curve of Mn does not have a plateau but has a significantly lower ascent. Mo can therefore be used as the element to increase the hot hardness of, for example, iron-containing binders in sintered hard metals. L. Prakash described that even a few percent of molybdenum are sufficient to achieve a significant effect on the hot hardness of Fe-containing hard metals (thesis by Leo J. Prakash, Universität Karlsruhe 1979, Fakultät für Maschinenbau, KfK 2984). The proportion of Mo which is actually present in the binder remains unclear, however, since Mo2C was used.
The metals which are to 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 hot hardness of the hard metal below 800° C. Precautions therefore have to be taken to ensure 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 standard industry practice to use carbides, metals or nitrides and set the carbon content of the hard metal by means of the formulation and measures during sintering so that the hard metal is at the edge of the existence region to an eta phase (η phase) and the maximum possible proportion of W and Cr is present in the binder. Cr is therefore generally added as Cr carbide which disproportionates during sintering, for instance, according to the following equation:
Cr3C2=>Cr (alloyed in the binder)+2CrC (alloyed in the WC)
Thus, only a fraction, namely ⅓, of the Cr used is effective in the binder. The situation is similar for Mo2C:
Mo2C=>Mo (alloyed in the binder)+(W,Mo)C.
When Mo carbide is used, only a maximum of about 50% is therefore effective in the binder alloy. Elemental Mo2C metal powder is used instead of Mo for this reason. Even when very finely divided Mo metal powder is used, however, regions which consist exclusively of binder alloy phase and contain no hard material are formed after sintering. This behaviour can be attributed to agglomerates of the Mo metal powder being comminuted ineffectively during mixed-milling because of the high modulus of elasticity of molybdenum and the resulting deformed agglomerates dissolving during liquid-phase sintering the molten binder alloy which in turn fills the pores formed by dissolution of the Mo particles in the molten binder. This results in formation of “binder lakes,” which is a term for a particular region of the binder alloy which has dimensions greater than the particle diameter of the hard material phase, but does not contain tungsten carbide or hard material particles.
These are disadvantageous and unacceptable for both the strength and for local wear resistance. Owing to the limited diffusion time, corresponding to the time within 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 is achieved at all.
If the molten binder does not fill the secondary pores formed during sintering, these are visible in the sintered body, as described in EP 1 043 411 B1. These secondary pores reduce the strength.
In an embodiment, the present invention provides that iron-, cobalt- or nickel-containing binder metal powders comprising iron in an amount of from 0.1 to 65% by weight, cobalt in an amount of from 0.1 to 99.9% by weight and nickel in an amount of from 0.1 to 99.9% by weight are used for producing sintered hard metals. The percentages are percentages by weight and are based on the binder alloy powder, unless indicated otherwise.
In an embodiment of the present invention, the binder alloy powder used contains from 0.1 to 10% by weight of molybdenum, based on the total binder metal powder, in alloyed form. The binder metal powder used can, for example, contain from 0.10% by weight to 3% by weight of molybdenum, for example, from 0.5% by weight to 2% by weight of molybdenum, or for example, from 0.5% by weight to 1.5% by weight of molybdenum, in each case based on the total binder metal powder. An excessively high molybdenum content leads to excessive strengthening of the binder powder, so that the pressing forces in the production of the hard metal and the resulting sintering shrinkage become too high, while an excessively low content leads to an insufficient increase in the hot hardness.
Hard materials can, for example, be carbides, such as tungsten carbide (WC). Binders can, for example, be alloys of iron, cobalt and nickel, such as the combinations iron and nickel, iron and cobalt, cobalt and nickel and also iron, cobalt and nickel. It is likewise possible to use cobalt alone as a binder.
The binder metal powders which have been alloyed with molybdenum display, owing to their physical nature, good dispersion behaviour in mix-milling with carbides to produce hard metal powders. The FSSS values (measured using the “Fisher Sub Siever Sizer” in accordance with the ASTM standard B330) can therefore be in the range of from 0.5 to 3 μm, for example, in the range of from 1.0 to 2 μm. Finer powders are pyrophoric; coarser powders no longer have a satisfactory dispersion behaviour and once again lead to “binder lakes.” The size distribution of the agglomerates can, for example, be in the range from 0.5 to 10 μm for the same reason. The specific surface area can, for example, be in the range from 2.5 to 0.5 m2/g for the same reasons. The oxygen content can, for example, be below 1.5%.
Cobalt contents in the binder alloy can, for example, be up to 60% by weight. The nickel content in the binder alloy can, for example, be in the range from 10 to 80% by weight, or, for example, from 20 to 60% by weight or, for example, from 30 to 75% by weight.
In an embodiment of the present invention, subsequently added organic additives can also be present. To determine the abovementioned parameters, these may have to be removed again, for example, by washing with a suitable solvent. The organic additives include waxes, agents for passivation and inhibition, corrosion protection, and pressing aids. Examples include paraffin wax and polyethylene glycols. The organic additives are also intended to prevent aging of the powder which would result in an increase in the oxygen content. The additives can, for example, be present in an amount of 30% by weight, based on the sum of binder alloy powder and additive.
The Mo-containing binder powder can contain Fe, Ni and Co. Since the sinterability and the hot hardness decrease with increasing Fe content, the iron content can, for example, be less than 65%, for example, less than 60%. The balance to 100% is Mo plus Co and/or Ni. Alloys in the system can, for example, include FeCoNi which are stably austenitic in the sintered hard metal, for example FeCoNi 30/40/30 or 40/20/40 or 20/60/20 or 25/25/50, and also FeNi 50/50 or 30/70 or 20/80, or CoNi in the ratios 50/50, 70/30 or 30/70, as binder alloys. However, it is also possible to use element powders such as Co or Ni alloyed with up to 10% of Mo, which thus become alloy powders.
The molybdenum-containing alloy powders can, for example, be produced by the process described in DE 10 2006 057 004 A1 where a MoO2, which has been comminuted to reduce the agglomerate size distribution, serves as a molybdenum source. This MoO2 is added to an oxalic acid suspension as used according to EP 1 079 950 B1 for preparing FeNi or FeCoNi mixed oxalates which are subsequently fired under oxidizing conditions and reduced by means of hydrogen to alloy powders. The alloy powders obtained in this way are fully reduced after reduction with hydrogen, for example, MoO2 can no longer be detected by means of X-ray diffraction. If appropriate, the agglomerate size can be reduced further by means of deagglomeration in order to improve dispersion in the mix-milling with the carbides. The agglomerates consist of primary particles which are agglomerated with one another. The agglomerate size and the distribution of the agglomerates can be determined by means of laser light scattering and sedimentation.
Instead of MoO2, it is also possible to use other finely particulate Mo compounds which do not dissolve in oxalic acid, for example, sulphides or carbides. These are oxidized to oxides in the calcination of the precipitated oxalate in air. Molybdenum oxides such as MoO3 are formed during calcination and, owing to their high vapour pressure, very quickly form mixed oxides with the Fe(Co)Ni mixed oxide and display good transport properties, so that an FeCoNi alloy powder which is homogeneously alloyed with a small proportion of Mo is formed in the subsequent reduction with hydrogen.
Other known processes are, however, also suitable such as a precipitation with ammonium salts of oxalic acid instead of oxalic acid, with Na or K hydroxide, with formic acid, and maleic acid. In all cases, a MoO2 can, for example, be used which should be phase-pure and contain only traces of Mo or MoO3 or Mo4O11. MoO2 is used because, in contrast to MoO3, it is soluble neither in acids nor alkalis and therefore remains completely in the alloy metal powder during the entire process. MoO3 would dissolve in the alkali used for precipitation of the Fe(Co)Ni content or in complexing organic acids; elemental Mo would be too coarse and would not oxidize completely to MoO3 in the subsequent calcination and thus not alloy satisfactorily during reduction with hydrogen. A fine MoO2 having a high specific surface area oxidizes completely to MoO3 (which has a high vapour pressure) during calcination of the Fe(Co)Ni oxalate in air and, via the gas phase, forms molybdates and mixed oxides with these metal oxides, which results in very uniform distribution of the molybdenum which is retained during the subsequent reduction with hydrogen.
Mo-alloyed FeCoNi powders can, for example, be used which contain the Mo in entirely metallic form. In these powders, Mo oxides can no longer be detected by means of X-ray diffraction, and accordingly, the oxygen present should be predominantly present on the surface of the powder. Useful powders according to the present invention include powders whose FSSS value is in the range from 0.5 to 3 μm because this improves dispersion in mix-milling. In this case, they do not contain, for example, any further metals in oxidic form.
Since Mo oxide reacts with carbon to form CO in the sintering of hard metals and can thus lead to a local carbon deficit and thus to local eta phases, the alloy powders described in the above paragraph are suitable for hard metal production when precautions are taken during sintering of the hard metal to provide that the oxygen liberated predominantly in the form of carbon monoxide can escape from the sintered body. These powders are suitable for use according to the present invention when they have the physical properties as described herein but only contain the above-described elements Mn, Cr, V, Al and Ti in at least partially oxidic form only to the extent which is permissible from the point of view of microstructural defects (pores and binder lakes) in the hard metal.
According to the present invention, the Mo-alloyed powders based on FeCoNi or FeNi can be additionally alloyed with up to 20% of tungsten, for example, to shift the commencement of sintering shrinkage to higher temperatures or induce the formation of precipitates which reinforce the binder phase, but this is successful only in the case of very coarse tungsten carbides.
In an embodiment of the present invention, the alloy powders used can be within a wide composition range of FeCoNi. In the range of high Fe contents (from 90 to 60%), binder alloy systems which, after sintering, have proportions of martensitic phase and therefore have a high hardness and wear resistance at room temperature are found. Examples are FeNi 90/10, 82/18, 85/15, FeCoNi 72/10/18, 70/15/15 and 65/25/10. However, the abovementioned alloys have very low hot hardnesses in the sintered hard metal. In the range from about 80 to 25% of Fe, the binder alloys after sintering are austenitic and although they have a lower intrinsic hardness, they have a high fatigue strength and ability to undergo limited plastic deformation. 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 the hard metals in the range from 400 to 600° C. is inferior to those having pure Co as binder if Mo or other alloying elements are not additionally incorporated in the alloy. Although an aspect of the use according to the present invention is the production of hard metals having improved hot hardness, it is also suitable for the production of hard metals with other aspects, such as a hard metal having molybdenum-containing corrosion-resistant binder alloy systems with are at present produced using elemental or carbidic molybdenum, for example, as described in EP 0 028 620 B2, or cutter inserts for drill bits, as described in U.S. Pat. No. 5,305,840.
The binder alloy present after sintering of the hard metal can, according to the present invention, also be obtained using a plurality of different alloy powders and optionally elemental powders as described in WO 2008/034 903, with at least one of these powders being alloyed with molybdenum. The advantages of such a procedure are the pressability and control of the sintering shrinkage.
The hard metal part present after sintering and, if appropriate, finer machining by grinding or electroerosion, has a defined tool geometry. This can, for example, be elongated (such as ground from a sintered round rod), but can also, for example, be plate-like for machining of materials such as metals, stone and composites by turning or milling. In all cases, the hard metal tools can, for example, have one or more coatings selected from among nitrides, borides, oxides and superhard layers (such as diamond, cubic boron nitride). These can, for example, have been applied by PVD or CVD processes or combinations or variations thereof and can have been modified in terms of their residual stress state after application. However, they can, for example, also be hard metal parts of any further geometries for any further uses, such as forging tools, forming tools, countersinks and counterbores, components, knives, scrappers, rollers, stamping tools, pentagonal drill bits for soldering-in, mining cutters, milling tools for milling machining of concrete and asphalt, rotating mechanical seals and also any further geometries and applications.
The present invention is illustrated by the following examples.
462.5 g of tungsten carbide 0.6 μm were mix-milled with 37.5 g of a FeCoNi alloy powder 40/20/40 (Ampersint® MAP A 6050; manufacturer: H.C. Starck, Germany) together with 0.57 1 of 94% strength ethanol at 63 rpm in a ball mill for 14 hours. 5 kg of sintered hard material balls were used for this purpose. 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 m2/g, FSSS value=0.90. Two batches having different carbon contents (“high carbon” and “low carbon”) were produced, so that different carbon contents resulted after sintering. The results are set forth in Table 1 below.
The ethanol was separated off from the resulting suspension by distillation under reduced pressure and the hard metal powder obtained was pressed uniaxially at 150 MPa and sintered at 1450° C. for 45 minutes under reduced pressure. The plate-like hard metal pieces were ground, polished and examined to determine their properties. As sintered bodies, both batches displayed neither eta phases nor carbon precipitates, but relatively small binder lakes. In both cases, the room temperature hardness and the hot hardness at selected temperatures up to 800° C. were measured under protective gas.
A WC—Co having the same proportion by volume of binder phase as in Example 1 was produced in a manner analogous to Example 1. Since Co has a higher density than FeCoNi 40/20/40, the proportion by weight of cobalt was 8% by weight, based on the total hard metal. Pressing and sintering at 1420° C. for 45 minutes under reduced pressure resulted in a defect-free hard metal having a magnetic saturation of 133 G·cm3/g, corresponding to 82% of the theoretical magnetic saturation. The room-temperature hardness (HV30 1597 kg/mm2) and the hot hardness were determined and plotted in
Example 1 at room temperature.
Example 1 was repeated with 1% by weight of Mo metal powder being added in a first batch and 3% by weight being added in a second batch. (These contents relate to 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% by weight. The particle size distribution was determined by the following parameters: D50 3.2 μm, D90 6.4 μm. The carbon content was selected so that, on the basis of experience from Example 1, neither eta phases nor carbon precipitates were to be expected in the sintered hard metal. For the Mo addition, no additional carbon was included so that the molybdenum was present virtually entirely in metallic form in the binder alloy. The carbon contents of the formulation were therefore 5.94 and 5.94% (3% by weight of Mo, based on the binder). The results after sintering at 1420° C. are set forth in Table 2. The hot hardnesses were determined as before and are represented by circles in
Surprisingly, eta phases occurred neither at 1% by weight nor at 3% by weight of molybdenum, rather, carbon porosity even occurred at 1% by weight of molybdenum. The hardness surprisingly increased without the K1C value being reduced compared to Example 1, so that at room temperature, a property combination which is equal to that of the Co-bonded hard metal and is superior to that of the purely FeCoNi-bonded hard metals was obtained. Surprisingly, 1% by weight of molybdenum in the binder is sufficient; at 3% by weight of molybdenum, no great change in the K1C and the hardness compared to 1% of Mo was observed. The effect of the molybdenum alloyed in the binder thus not only increases the intrinsic hardness of the binder but also simultaneously increases the fracture toughness. The behaviour is in this respect different from the case of alloyed W. Although an increase in the intrinsic hardness of the binder was also found here, there was a simultaneous decrease in the K1C value, both in Co-based hard metals and in materials based on FeCoNi, see Example 1.
However, many binder lakes occur, which is evidence of the dissolution of Mo in the binder which then fills the resulting pore volume. However, these binder lakes are not acceptable in a hard metal.
The comparison of the hot hardnesses with those from Example 2 is shown in
Example 1 was repeated using the FeCoNi binder alloy alloyed with 1.5% by weight of Mo produced by the process described in DE 10 2006 057 004 A1. The powder was subsequently 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 area=2.17 m2/g, D50=3.46 μm, D90=5.84 μm MoO2 could no longer be detected by means of X-ray diffraction at its characteristic diffraction angles even after long-term exposure. 37.5 g of this powder were used together with 462.5 g of WC for producing a hard metal. The hard metal mixture had a carbon content of 5.92% by weight, which was set by addition of 1.14 g of carbon black. The pressed bodies were sintered both in an open crucible and in a closed crucible. This variation has effects on the carbon content of the hard metal after sintering. The properties of the hard metal sintered at 1420° C. are set forth in Table 3.
The hard metal from the open sintering is at the low-carbon end of the two-phase region since it has a very low magnetic saturation compared to Example 1. However, eta phases could not be detected. The maximum possible concentration of Mo in the binder results in an enormous strengthening of the binder alloy, which is reflected in the simultaneous increase in hardness and fracture toughness. The hard metal from the closed sintering is also in the 2-phase region in respect of the carbon content but contains more carbon, which can be seen from the high magnetic saturation. Since more Mo is present as carbide due to the higher carbon supply and is therefore not present in the binder, the fracture toughness, which is determined essentially by the binder, decreases very greatly to the level of the “high carbon” variant of Example 1.
Further pressed bodies were produced and sintered at 1420° C. under reduced pressure, but argon was injected at a pressure of 40 bar towards the end of sintering at the final temperature. Cooling was carried out under pressure. Hard metal pieces having a hardness of 1643 HV30, a fracture toughness of 8.2 MPa·M1/2 and a magnetic saturation of 123 G·cm3/g were obtained. Both the room-temperature hardness and the hot hardness were determined as a function of the temperature on the hard metal pieces on another hardness testing machine. The evaluation of the results for the room-temperature and hot hardness is shown in
It can be seen that the use according to the present invention of a binder powder which is (pre)alloyed with molybdenum makes it possible to produce a defect-free hard metal without binder lakes and with a hot hardness curve virtually the same as with a cobalt binder. The decrease in the hot hardness around 600° C. is virtually eliminated. In addition, there is both an enormous improvement in the room-temperature strength and an increase in the hardness compared to Example 1 when the carbon balance is set appropriately in comparison with Example 1, which likewise offers advantages for applications at or close to room temperature. An improvement in the corrosion resistance compared to Example 1 is also to be expected since corrosive attack on hard metals generally occurs via the binder phase.
The principle of improving the properties of hard metals by means of alloyed molybdenum in the binder can be applied not only to the FeCoNi 40/20/40 binders described, but also to pure cobalt and to pure Ni as the hard metal binder, to CoNi and FeNi alloys and to further FeCoNi alloys.
The present invention is not limited to embodiments described herein; reference should be had to the appended claims.
| Number | Date | Country | Kind |
|---|---|---|---|
| 10 2008 052 104.3 | Oct 2008 | DE | national |
| 10 2008 052 559.6 | Oct 2008 | DE | national |
This application is a U.S National Phase application under 35 U.S.C. §371 of International Application No. PCT/EP2009/062844, filed on Oct. 2, 2009 and which claims benefit to German Patent Application No. 10 2008 052 104.3 filed Oct. 20, 2008 and to German Patent Application No. 10 2008 052 559.6, filed on Oct. 21, 2008. The International Application was published in German on Apr. 29, 2010 as WO 2010/046224 A2 under PCT Article 21(2).
| Filing Document | Filing Date | Country | Kind | 371c Date |
|---|---|---|---|---|
| PCT/EP09/62844 | 10/2/2009 | WO | 00 | 4/11/2011 |
