GB2100283A - A hard alloy containing molybdenum - Google Patents

Abstract

A hard alloy comprising a hard phase consisting of at least one compound having a crystal structure of the simple hexagonal MC type (M: metal; C: carbon) which is a mixed carbide, carbonitride or carboxynitride of molybdenum and tungsten, and a binder phase consisting of at least one of iron, cobalt, and nickel, the hard phase being prepared by carburizing an (Mo,W) alloy obtained by reducing oxides of molybdenum and tungsten having a particle size of 1 micron or less and comprises coarse particles having a mean particle size of at least 3 microns and a uniform molybdenum to tungsten ratio and having a gross composition within the shaded portion ABCDEA of Figure 1 of the accompanying drawings. <IMAGE>

Description

SPECIFICATION A hard alloy containing molybdenum This invention relates to a hard alloy containing molybdenum and more particularly, it is concerned with a hard alloy comprising a hard phase consisting of at least one compound having a crystal structure of the simple hexagonal MC (M: metal, C: carbon) and a binder phase, which is suitable for use as a tool capable of resisting a high impact for a long time.

A (Mo, W) C base alloy is described in British Patent No. 635,221. The process described for producing the (Mo, W)C base alloy comprises nitriding oxides of molybdenum and tungsten in an ammonia stream, carburising the nitrides with release of nitrogen, adding an auxiliary metal in powder form to serve as a binder for the sintered alloy, and sintering. This alloy was an alloy consisting of one or two carbides of (W, Mo)C and (W,Mo)2C with a binder metal, but has not been put to practical use.

Molybdenum monocarbide (MoC) is considered as a useful substitute, since this carbide has the same crystal structure, a simple hexagonal type, as that of tungsten carbide as well as mechanical properties similar to those of tungsten carbide. However, the existence of the hexagonal molybdenum monocarbide as a simple substance has remained in question to this day and thus attempts to stabilize molybdenum have exclusively been carried out by forming a solid solution with tungsten carbide. This method was firstly reported by W. Dawhil in 1950, but the solid solution was not examined in detail and no commercial application was found at that time.

Recently, however, attempts to utilize the solid solution (MOxW,,)C where x+y=1 have increased with the rise of the price of tungsten.

In the prior art process for the production of a solid solution of MoC-WC, WC, Mo and C powders or W, Mo, C and Co powders are mixed, charged in a carbon vessel and reacted to a temperature of 1 660 to 20000C (W. Dawhil: "Zeitschrift f. Anorganische Chemie" 262 (1 950) 212). In the latter case, cobalt serves to aid in forming the carbide and to dissolve molybdenum and carbon in the tungsten carbide. In the absence of cobalt, it is very difficult to obtain a solid solution of (Mo,W)C.

When a (Mo,W)C powder obtained by this method is used for the production of a cemented carbide alloy with a binder metal of cobalt as a substitute for WC, however, (Mo, W)C decomposes in the alloy to deposit needle crystals of (Mo, W)2C. Deposition of even a small amount of such a subcarbide with aggregation in the alloy causes a deterioration in the strength of the alloy. For this reason, the use of MoC as a substitute for WC has not been positively attempted.

Generally, processes for the production of mixed carbides, the separate carbides are heated in the presence of each other, optionally using a diffusion aiding agent such as cobalt, to give a uniform solid solution in most cases. However, for a solid solution- containing at least 70% of MoC, a uniform solid solution cannot be obtained by counter diffusion only at high temperatures. This is due to the fact that MoC is unstable at high temperatures and is decomposed into solid solutions such as (Mo,W)C,~x and (Mo,W)3C2 and, consequently, a solid solution (Mo,W)C of the WC type cannot be formed only by cooling.As a method of stabilizing this carbide, it has been proposed to react the components once at a high temperature, to effect diffusion of Mo2C and WC, and then to hold the product at a low temperature for a long time (Japanese Patent Application (OPI) No. 146306/1976) However, a considerably long diffusion time and long recrystallisation time are required for forming (Mo, W)C from (Mo,W)C~x and (Mo,W)3C2 at low temperatures. In order to practice this method on a commercial scale the mixture needs to be heated for a long time in a furnace to obtain a complete carbide. This means that the productivity per furnace is low and a number of furnaces are thus required. When using a continuous furnace, on the other hand, a long furnace is necessary and mass production is therefore difficult industrially.

We have investigated the possibility of preparing an alloy consisting of a solid solution of (Mo W) and a (Mo-W)C powder as a hard material on a commercial scale, since the use of these materials or their cemented carbide alloys, will then increase and consequently, we have reached the invention as disclosed in US Patent No. 4,126,009.That specification discloses a process for the production of an alloy powder containing molybdenum and tungsten and having a crystal structure of the simple hexagonal WC type, which comprising mixing molybdenum and tungsten in the form of compounds thereof selected from oxides, hydroxides, chlorides, sulfates, nitrates, metallic acids, salts of metallic acids and mixtures thereof, the resulting mixture of the compounds having a particle size of less than 1 micron, reducing the mixture with one of hydrogen and ammonia to form an alloy powder or molybdenum and tungsten, and then carburizing the alloy powder.

Furthermore, we have proposed cemented carbide alloys as disclosed in Japanese Patent Applications (OP I) Nos. 145,146/1980 and 148,742/1980, which are suitable for impact resisting tools. The first specification discloses an impact resisting cemented carbide alloy containing molybdenum, characterised in that the friction coefficient is less than 70% of that between WC-Co type alloys and steels, but this alloy does not have a sufficient life, in particular in uses subject to repeated impacts because it contains a hard phase of MC type in the alloy.The second specification discloses an impact resisting cemented carbide alloy comprising a hard phase of mixed carbides of molybdenum and tungsten of the MC type and a binder phase of cobalt and nickel, represented by (Mo,W,,)C,-(Ni,Co,,) where 0.5~x < 0.95,0.5~y~1.0 and 0.90~z~0.98, but this alloy does not have a long life under severe conditions such as being subjected to high impacts for a long time.

According to the present invention there is provided a hard alloy comprising a hard phase consisting of at least one compound having a crystal structure of the simple hexagonal MC type (M: metal; C: carbon) which is a mixed carbide, carbonitride or carboxynitride of molybdenum and tungsten, and a binder phase consisting of at least one of iron, cobalt, and nickel, the hard phase being prepared by carburizing an (Mo, W) alloy obtained by reducing oxides of molybdenum and tungsten having a particle size of 1 micron or less and comprises coarse particles having a mean particle size of at least 3 microns and a uniform molybdenum to tungsten ratio and having a gross composition within the shaded portion ABCDEA of Figure 1 of the accompanying drawings.

The accompanying drawings are to illustrate the principle and merits of the present invention in more detail: Figure 1 is a graphical representation of the composition of a hard alloy according to the present invention in the relationship of W/(Mo+W) atomic ratio and binder content.

Figure 2 is a graphical representation of the relationship between the carbon content of the alloy and the change of the transverse rupture strength (TRS) in an (Mo0,7W03)C-35 wt% Co alloy.

Figure 3 is a graph showing the results of a fatigue test of an (Mo0,7W0.3)C-35 wt% Co alloy (A: MC-y alloys; 0: MC-M2C-y alloys) in which a cyclic load y is applied.

Figure 4 is a micrograph, magnified 5,000 times, of (Mo0.5W0,5)C according to the present invention, in which Mo and W are uniformly distributed.

Figure 5 is a micrograph, magnified 5,000 times, of (MoO sWO 5)C according to the prior art, in which Mo/W ratio is not uniform in each particle.

We previously made various efforts to improve (Mo W)C-iron group metal alloys and consequently, found that the uniform dispersion of granular or globular (Mo, W)2C (which will hereinafter be referred to as M2C) therein is effective for increasing the stress yield and breaking strength (US Patent No. 4,265,662), but this alloy is not suitable for uses where a high fatigue strength is required upon exposure to high impacts for a long time. This is possibly because the dispersed (Mo,W)2C acts as a harmful element for this purpose. Where a high fatigue strength is required upon exposure to high impacts for a long time, "crack propagation" is regarded as important rather than "crack initiation" and in particular, "crack propagation" tends to extend along the boundaries between the hard phase and binder metal.Thus, it is necessary to reduce the boundaries of the hard phase and binder metal and this can be achieved by increasing the particle size of the hard phase and the thickness of the binder phase.

We have carried out a heading test of bolts using dies made of the materials as shown in Table 1 by changing the particle size of the hard phase, thus obtaining the results as shown in Table 2: Table 1 Particle Compressive size of strength TRS* carbide (jut) (KG/mm2) (Kg/mm2) VHN** (Mo,W)C-25wt%Co (A) 5 327 270 760 (Mo,W)C-25wt%Co (B) 1 345 295 810 WC-22wt%Co (C) 5 310 280 830 Note: *Transverse Rupture Strength **Vickers Hardness Number Table 2 Die Life (Number of Samples Processed/Die) x 105 0 2 4 6 8 10 Alloy (A) x Alloy (B) x Alloy (C) x Note: Mark x means a broken point.

Test Conditions: workpiece: S45C Steel forging speed: 100 samples per minute As is evident from these results, the strength of an alloy and the die life are not always consistent with each other and the low hardness and low strength alloy with coarse grain size exhibits the longest life.

Thus, the tool can be used even after cracks or deformations occur. This suggests that the life of a tool does not depend on the initiation of cracks but depends on the propagation speed of the cracks leading to the overall breakage thereof.

We therefore concentrated our efforts on the preparation of a coarse grain carbide and consequently, have found that it is more difficult to obtain an (Mo, W)C with a large particle size than WC since molybdenum has the effect of retarding particle growth. However, we further found that when using particularly the solid solution (Mo-W) prepared by the process of US Patent No.

4,216,009 as mentioned above, a carbide with a particle size of 3 microns or more could readily be obtained by controlling the carburizing conditions, for example, by adjusting the carburization temperature to a temperature which is sufficiently high but lower than the decomposition point of (Mo, W)C into (Mo, W)2C, for example, to 1 4500C in the case of (Mo0,7W0.3)C. For the preparation of a carbide with a larger particle size, e.g. 6 microns or more, the carburization is preferably carried out after the solid solution (Mo-W) has been subjected to heat treatment. The heat treatment is generally carried out at a temperature of from 1100 to 1 4000C for 1 to 5 hours in a stream of nitrogen or hydrogen.In the case of (MoO 5 W0.5)C, for example, the solid solution is thermally treated at 1 3000C for 3 hours in a stream of nitrogen.

When WC, Mo2C and C are used as starting materials and subjected to carburization according to the prior art, on the other hand, it is very difficult to form a coarse carbide with a particle size of 3 microns or more, and even if coarse starting materials are used, only a carbide having a fluctuating Mo/W ratio is formed in each particle, because the carbon in the carbide acts as a diffusion barrier. The use of such a carbide with a binder metal results in a non-uniform alloy which has a low mechanical strength.

In view of the above described facts, it will clearly be understood that an alloy containing a coarse carbide with a particle size of at least 3 microns, preferably at least 5 microns and having a desired mechanical property, i.e. impact resistance must be prepared by the process which comprises reacting a solid solution of (Mo, W) with carbon, which is capable of forming a uniform and large particle size (Mo, W)C, since otherwise the preparation of such an alloy is impossible.

As a result of our studies on the sintering of an alloy consisting of two phases of (Mo, W)C and a binder metal we have found that in (Mo, W)C base alloys, the growth of carbide particles does not take place due to Ostwald ripening of the dissolving and precipitating type which can be seen in ordinary WC base alloys, but a very slow particle growth of the diffusion rate-controlling type is found. In the (Mo, W)C base alloys, the particle growth during sintering, which can be seen in the prior art WC base alloys, is scarcely expected and therefore, a carbide to be used as a raw material must be of course particles so as to prepare an alloy containing a hard phase with a large particle size.

Similarly, we have conducted various experiments and measured typical properties in order to make clear the features of the M2C precipitated alloys and M2C non-precipitated alloys.

Figure 2 is a graphical representation of the relationship between the carbon content (% by weight) and transverse rupture strength (Kg/mm2) in an (Mo0.7W0,3)C-35 weight % Co alloy. As is apparent therefrom, the transverse rupture strength rapidly lowers with the precipitation of free carbon, but does not lower even if M2C is precipitated. This is considered to be due to the fact that the M2C phase is dispersed uniformly and finely so that dispersion strengthening appears, but it is hardly related to the lowering of the transverse rupture strength.

Figure 3 is a graph showing the results of a fatigue test on an (Mo0.7W0.3)C-35 weight % Co alloy, in which a static load at a certain level is applied cyclically to a sample. It is apparent from these results that the M2C-precipitated alloy (MC-M2C-y alloys represented by mark O) is inferior to the M2C-nonprecipitated alloy (MC- alloys represented by mark A) in fatigue strength. This is possibly due to the fact that the finely dispersed M2C increases the boundaries between the hard phase and binder metal phase and acts as an element to promote crack propagation, since cracks propagate predominantly along the boundaries between the hard phase and the binder phase.In the case of wear resisting tools, in general, a high stress is intermittently applied for a long time and in addition, some factors promoting crack propagation, such as thermal impact and corrosion embrittlement are entangled, so that a high fatigue strength is required. In such a case, M2C should not be precipitated.

M2C tends to aggregate and grow abnormally large with the increase of the quantity thereof, which acts as a stress concentrating source causing a lowering in the fatigue toughness when high impact energy is applied.

As the same time, it is also proved by a field test that the quality of M2C precipitated should be as low as possible or reduced to substantially nil for the purpose of sufficiently displaying the performance of a tool where an alloy is used having a relatively large binder phase and a structure such that the mean free path of the binder phase is large.

We have made further studies on a hard alloy consisting of (Mo, W)C and an iron group metal by changing the ratio of Mo and Wand the amount of the iron group metal over a wide range and it is thus found that the two phase region {(MC+y) zone) is about 1/3 of that of WC base alloys in the case of (Mo07W03)C base alloys, and about 1/5 of that of WC base alloys in the case of (Mo09W01)C base alloys. When the binder metal is changed from cobalt to iron, some shift takes place in the two phase region, but there is little change in width. These data are collected and arranged to give the results as shown in Figure 1 wherein the boundary line of the M2C-precipitated zone and the M2Cnonprecipitated zone is drawn by line a.

Furthermore, we have conducted a number of experiments by changing the ratio of Mo and W and the quantity of iron group metals over a wide range and have thus found that when the critical width of controlling carbon industrially is 0.07% by weight, a zone wherein there is (MC+y) in an amount of at least 0.07% by weight as carbon can be represented by the relationship of q x r~4.0, p+q i.e. above line a in Figure 1, where the alloy composition consists of (MopWq)C-r weight % binder metal. In other words, it has been quantitatively determined that M2C tends to be precipitated with an increase in the ratio of Mo in the ordinary (Mo, W)C base alloys and the amount of the binder metal should be increased so as to suppress precipitation of M2C.For example, as can be seen from Figure 1, the two phase region of (MC+y) free from precipitation of M2C and free carbon amounts to at most 0.07% by weight of carbon content used as a parameter unless the amount of a binder metal is more than 13.5% by weight in the case of (MoO WO 3)C base alloys, and at least 20% by weight of a binder metal is required for a similar carbon value in the case of (Mo0.8W0.2)C base alloys. As a matter of course, line a is shifted above when the (MC+y) zone exceeds 0.07% by weight or carbon content, wherein the two phase region of (MC+y) is stable.

Referring to Figure 1 , the reason for limiting the W/(Mo+W) ratio to W 0.1 < ~0.9 Mo+W is that if the ratio is less than 0.1, the carbide is so unstable that it tends to be decomposed into M2C while if it is more than 0.9, there is little effect of molybdenum as (Mo,W)C. The reason for limiting the amount of the binder metal to 10 to 70% by weight is that if it is less than 10% by weight, the alloy itself becomes so brittle that it cannot in fact, while if it is more than 70% by weight, the sintering is so difficult that the desired shape cannot be held.

The iron group metal as the binder phase can naturally dissolve Group IVa, Va and Vla metals and it is possible to add other elements having a solubility therein such as aluminium, silicon, calcium, silver, etc.

The basic concept of the present invention can be maintained even when a part of the molybdenum and tungsten carbide is replaced by a BI type mixed carbide containing titanium, zirconium, hafnium, vanadium, niobium, tantalum, chromium, molybdenum and/or tungsten in a proportion of 30% by weight or less, preferably 0.5 to 25% by weight.

Furthermore, there is a similar relationship even for an alloy wherein a part of the C in the carbide is replaced by nitrogen and/or oxygen. Examples of the preferred embodiments in this case are as follows.

The first embodiment is the incorporation of N in (W, Mo)C to give (W,Mo) (C,N) whereby a stable starting material of the hexagonal WC type can be obtained without a long heat treatment.

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

The third embodiment is the incorporation of Cr in (W,Mo) (C,N) or (W,Mo) (C,N,O) to give (W,Mo,Cr) (C,N) or (W,Mo,Cr) (C,N,O) whereby a starting material with a low weight and low price can be obtained.

The fourth embodiment is that in the production of these starting materials a mixture of oxides, metals, carbides and/or carbon is exposed to an atmosphere having a nitrogen partial pressure of 300 Torr or more at a temperature of 7000C or higher in a part of the carburization step to form a stable starting powder.

The fifth embodiment is that, when the above described starting powder is combined with an iron group metal, two or more hard phases of the simple hexagonal WC type differing in composition are caused to be present in the finished alloy, thereby imparting a high toughness thereto.

In these five embodiments, a part of the MC type phase can also be replaced by a BI type solid solution containing one or more of Group IVa, Va and Vla metals and non-metallic elements, or the ordinary additives to cemented carbides, such as silver, silicon, bismuth, copper aluminium etc. can also be added to the iron group binder metal.

The above described embodiments will now be illustrated in greater detail: For the alloys of the present invention comprising a simple hexagonal phase containing molybdenum and tungsten, it is found in the sintered alloy with a binder metal that, when N atom % W atom q/o (Mo+W) atom % Mo+W atom % a suitable range of A is 0.005~A~0.5. If A is less than the lower limit, the effect of nitrogen does not appear, while if more than the upper limit, sintering to give excellent properties is difficult. The most preferred range of A is 0.01 ~A < 0.4.

Concerning the effect of oxygen, it is found that, when Oatom% (1- Watom% (Mo+W) atom % (Mo+W) atom % a suitable range of B is 0.005 < B~0.05. If B is less than the lower limit, there is no favourable effect of oxygen, while if more than the upper limit, sintering is difficult to give excellent properties. The most suitable range of B is 0.01 B~0.04.

On the other hand, the W/Mo ratio is preferably 10/90 to 90/10, since if less than 10/90, the alloy is unstable, while if more than 90/10, the merits of the replacement (light weight, low price) are substantially lost. The quantity of chromium used for replacing molybdenum or tungsten is 0.5 or less by atomic ratio of (W+Mo), since if more than 0.5, the alloy is brittle although the corrosion resistance increases.

As is well known in the art, it is advantageous for cutting tools to form a BI type solid solution composed of at least one Group IVa, Va and Vla metals such as titanium, zirconium, hafnium, vanadium, tantalum, chromium, molybdenum and tungsten with at least one non-metallic component such as carbon, nitrogen and oxygen in addition to the simple hexagonal phase. The quantity of the BI type solid solution preferably depends upon the cutting use.

Concerning the quantity of nitrogen in this case, we have found that, when the definition of A is changed to N atom % W atom 9/0 ) x(1 Group IVa, Va, Vla metals-atom % Group IVa, Va, Vla metals atom % a suitable range of A is also 0.005~A < 0.5 although a part of the nitrogen is occluded in the BI type solid solution. The optimum range of A is 0.01 < A= < 0.4. Concerning the quantity of oxygen, we have found that, when the definition of B is changed to O atom % W atom % Group IVa, Va, Vla metals atom % Group IVa, Va, Vla metals atom % a suitable range of B is also 0.005~B~0.05. The preferred range of B is 0.01 ~B~0.04.

As the binder metal, iron group metal is preferably used, in a proportion of 10 to 70% by weight based on the gross composition, since if less than 10% by weight, the alloy is brittle and if more than 70% by weight, the alloy is too soft.

For the preparation of the starting materials, the reaction is carried out at a high temperature in a hydrogen atmosphere for the carburization of a (Mo, W) powder with carbon, reduction and carburization of oxide powders with carbon or a combination thereof. We have found from our studies on the decomposition nitrogen pressure of (Mo, W) (C, N) that the external nitrogen pressure depending on the temperature, should be 300 Torr or more at 7000C or higher at which the carbonitrization reaction takes place. The coexistence of hydrogen is not always harmful, but it is desirable to adjust the quantity of hydrogen to at the most two times as much as that of nitrogen, preferably to the same as that of nitrogen so as not to hinder the nitriding reaction. When using an ammonia decomposition gas, it is necessary to enrich it with nitrogen.

For the preparation of starting materials containing oxygen, the coexistence of carbon monoxide and carbon dioxide is required. In this case, the quantity of hydrogen is not limited as described above, but should not exceed 50% of the gaseous atmosphere. Heating and sintering in an atmosphere of nitrogen or carbon oxide is effective for the purpose of preventing a sintered alloy from denitrification or deoxidation.

We have found that the deformation at high temperatures can be improved by changing tungsten carbide to a carbide composed of a solid solution of three elements, molybdenum, tungsten and chromium. Thus, a (Mo,W)C-Co alloy has a higher hardness at a high temperature than a WC-Co alloy and, when Cr is further dissolved in this carbide, the hardness is further raised and the high temperature hardness is also improved. Thus, the disadvantages of the prior art WC-Co alloy can be overcome. It is also to be noted that the carbide phase consists of a solid solution of (Mo, W, Cr)C. It is also found that when Cr is dissolved in a solid solution of (Mo,W)C, the carbide particles can be made finer and stabilized as a monocarbide of (Mo, W, Cr)C.On the contrary, the known method of merely adding chromium to the binder phase has the disadvantage that it is impossible to make the carbide finer and the carbide phase is not stabilized as a monocarbide of a solid solution of (Mo, W, Cr). The quantity of chromium to be added to the solid solution carbide (Mo, W)C preferably ranges from 0.3 to 10%, since if it is less than 0.3%, the carbide cannot be made finer, while if it is more than 10%, Cr3C2 separates and is precipitated in the alloy, resulting in lowering of the hardness.

In a further embodiment of the present invention, a part of the carbon in the solid solution carbide (Mo, W, Cr)C is replaced by nitrogen, oxygen and/or hydrogen. That is, it is assumed that if the carbon contained in (Mo, W, Cr)C is added as a solid and reacted with a reactivity of 100%, the crystal is stabilized, but now it is found that the incorporation not only of carbon but also of nitrogen results in the stabilization of the monocarbide as (Mo,W,Cr) (CN) and the further incorporation of oxygen and hydrogen stabilizes the monocarbide as (Mo,W,Cr) (CaNbOcHd) (a+b+c+d=1), because if there are defects in the carbide, the carbide is unstable during sintering and an M2C type mixed carbide precipitates needlewise to lower the strength.

In a still further embodiment of the present invention, one or more of manganese, rhenium, copper, silver, zinc and gold are incorporated into the binder phase to change the micronstructure of the binder phase and to make the alloy non-magnetic. At the same time, it is found that, when these elements are added, the binder phase is alloyed, whereby the corrosion resistance of the alloy is improved.

In the last embodiment of the present invention, the toughness of the alloy can be raised by using, in combination, two or more carbides having a simple hexagonal phase but differing in their ratio of MoAW. The detailed reason of increasing the toughness is not clear, but it is assumed that when (Mo, W)C is separated into two phases, the solution strain of both of the phases is lowered to give a higher toughness than in the case of a single phase. Since at least an alloy consisting of a (MoxWy)C (y > x) phase having a similar property to that of WC and a (MoxWy)C(x > y) phase having a similar property to that of MoC has two properties, i.e. the toughness of WC and the heat and deformation resistance of MoC, this embodiment is more advantageous than when using one kind of (Mo,W)C only.Most preferably, the carbide is composed of WC or a solid solution of some MoC dissolved in WC and a solid solution of WC dissolved in MoC. This corresponds to the case where the peak of plane (1,0,3) is separated in two on X-ray diffraction. Whether there are two or more simple hexagonal phases of (MoxWy)C or not can be confirmed by observation using an optical microscope after etching with an alkaline solution of a hexacyanoferrate (III) or XMA by observation.

The application or use range of the alloy of the present invention is as follows. For example, the alloy of the present invention can be used for wear resisting tools such as guide rollers, hot wire milling rollers, etc., and for cutting tools, because it has a toughness and hardness similar to or greater than those of WC-Co alloys. In particular, when the alloy of the invention as a substrate is coated with one or more wear resisting ceramic layers as such of Tic,TiN,AI203, cutting tools of better toughness as well as wear resistance can be obtained than for the prior art tools having WC-Co type alloys as the substrate. As is well known in the art, a decarburization layer called phase is formed at the boundary between the substrate and coating layer and this appears similarly in the alloy of the present invention.

In order to prevent embrittlement directly under the coating layer due to decarburization, the presence of free carbon (FC) in the surface layer within a range of 300 microns is effective without reducing the toughness.

When using the alloy of the present invention as a watch frame, it shows better properties than WC-Co type alloys, which are summarized below: (1) A good brightness can be given when the alloy is specularly finished.

(2) Grinding and polishing are possible.

(3) Corrosion resistance is good, in particular, for sweat in the case of trinkets.

(4) Mechanical strength is high.

The present invention will be further illustrated in greater detail in the following Examples. It will be self-evident to those skilled in the art that the ratio, ingredients in the following formulation and the order of operations can be modified within the scope of the present invention. Therefore, the present invention is not to be interpreted as being limited to the following Examples. All parts, percents and the like are to be taken as those by weight unless otherwise indicated.

Example 1 54 g of Mo powder and 46 g of W powder were dissolved in 20% aqueous ammonia, neutralized with 6 N hydrochloric acid to coprecipitate and then subjected three times to filtration with water washing and drying. In the resulting precipitate, W03 and MoO3 were finely mixed. The mixed oxides were fired at 8000C in the air and then reduced at 10000C in a hydrogen stream. X-ray diffraction showed that the resulting powder was of a complete solid solution of (MoO.7WO 3).

The resulting solid solution (Mo0.7W0.3), carbon powder and Co powder as a diffusion aid were mixed in such a proportion that the final composition be (MoO 7 W03)C1.0 and subjected to carburization reaction at 1 4500C for 1 hour in a nitrogen stream under a nitrogen pressure of 1 atm. It was found by X-ray diffraction that the carbide had a crystal structure of simple hexagonal WC type and measurement of the particle size using Fisher Sub Sieve Sizer showed a mean particle size of 4.5 microns.

This powder was mixed with Co powder in such a proportion that the final composition be (Mo0.7W0.3)C-30% Co, ball milled with alcohol medium, pressed in a desired shape and then sintered in a vacuum of 10-2 Torr. The thus obtained alloy had a structure consisting of two phases of MC phase and binder metal phase, and a hardness of 880 by Vickers hardness and a bending strength of 290 Kg/mm2.

Example 2 A solid solution of (Mo0.5W0.5) was prepared in an analogous manner to Example 1 except changing the Mo/W atomic ratio to 0.5:0.5. This solid solution was mixed with carbon powder and Co powder as a diffusion aid in such a proportion that the final composition be (Mo05W05) C1,0 and subjected to carburization reaction at 1 5000C for 1 hour in a nitrogen stream under a nitrogen pressure of 1 atm. It was found by X-ray diffraction that the resulting carbide powder had a crystal structure of simple hexagonal WC type and by measurement of the particle size thereof using Fisher Sub Sieve Sizer that it had a mean particle size of 5.2 microns (Cf. Figure 4).

This powder was mixed with Ni powder and Co powder in such a proportion that the final composition be (Mo0,5W0,5)C-1 5% Ni-1 5% Co, ball milled with alcohol medium, pressed in a desired shape and then sintered in a vacuum of 10-2 Torr. The thus obtained alloy had 5.49% of total carbon and 0.01% of free carbon as analytical values, a structure consisting of two phases of MC and binder metals, and a hardness of 900 by Vickers hardness and a bending strength of 300 Kg/mm2.

For comparison, WC powder with a particle size of 4 microns, Mo2C powder with a particle size of 3.5 microns, carbon powder and Co powder as a diffusion aid were mixed in such a proportion that the final composition be (Mo0.5W0.5)C1.0, and sintered at 17000C in a vacuum of 10-2 Torr. The temperature was then lowered to 1 3500C and the mixture was held at this temperature for 12 hours.

X-ray diffraction showed that the resulting carbide was substantially of (Mo, W)C having a crystal structure of simple hexagonal WC type, but there were partly peaks of (Mo,W)zC (Cf. Figure 5).

This powder was mixed with Ni powder and to powder in such a proportion that the final composition be (Mo0.5Wc.5)C1 5% Ni-1 5% Co and an alloy was prepared in the same manner as described above. The thus obtained alloy had 5.48% of total carbon and 0.02% of free carbon as 5.48% of total carbon and 0.02% of free carbon as analytical values, a structure consisting of MC phase, binder metal phase and M2C phase grown up through aggregation, and a hardness of 910 by Vickers hardness and a bending strength of 230 Kg/mm2.

When using carbides as starting materials as in this comparative example, a uniform quality mixed carbide cannot always be obtained in the formation of the mixed carbide with a large particle size, and in an alloy obtained from this carbide, there is locally a carbon deficiency zone leading to formation of an aggregated M20 phase and resulting in lowering of the strength thereof.

Referring to Figure 4 and Figure 5, the heights of peaks of W and Mo show respectively the contents of W and Mo on the lines drawn in the micrographs of carbide crystals. It will be understood from comparison of Figure 4 and Figure 5 that in the case of Figure 4 according to the present invention, the fluctuation of peaks of W and Mo on the line crossing the carbide crystal is smaller, i.e.

the Mo/W ratio is more constant, than in the case of Figure 5 according to the prior art.

Example 3 Various alloys each having a composition within the range of shaded area of Figure 1 were prepared and subjected to measurement of the hardness and transverse rupture strength, thus obtaining results shown in Table 3.

The sintering temperature was varied from 12000C to 14500C every composition: Table 3 Sample No. Carbide Binder Hardness TRS Our Invention Composition Metal (Hv) {Kg/mm2) 1 (Mo085W015)0 30Ni--5Co 870 310 2 (Mo085W015)C 40Ni 720 280 3 (Mo085W0.5)C 30Co--30Ni 500 4 (Mo0.7W0.3)C 1 ONi-1 OCo 1150 260 5 (Mo0.70W0.3)C 15Ni--30Co 600 250 6 (Mo0 W0.3)C 50Ni--15Co 445 7 (Mo0.55W0.45)C 1 5Ni 1280 245 8 (Mo0.50W0.50)C 20Ni--5Co 1020 265 9 (Mo030W07o)c 35Co 880 300 10 ( MoO 30W0,0)C 25Co--10Fe 910 270 11 (MoO.30WO.70)c 8Fe-1 2Ni 1070 245 12 (Mo0.25W0.75)C 15Co--5Ni--8Fe 900 230 13 (Mo0.20W0.80)C 15Ni--40Co 600 240 14 (MQ,20W080)0 30Fe 950 245 15 (Mo0.20W0.80)C 1 ONi-SCo 1260 240 16 (Mo0.20W0.80)C 5Ni-50o-2Fe 1280 210 Table 3 (cont.) Comparison 17 (Mo0.80W020)C 5Ni--5Co 1410 110 18 (Mo0.70W0.30)c 2Ni 4Co 1700 90 19 (MoO 50Wo 50)C 60Ni--1 5Fe 420 20 (MoOOsWO9s)c 3Ni 1820 75 Example 4 The solid solution (MoO sWO.s) prepared in Example 2 was mixed with carbon powder, Cr3C2 powder and Co powder as a diffusion aid in such a proportion that the final composition be (Mo0.45W0.45Cr0.10)C1.0, subjected to primary carburization at 1 8000C for 1 hour in a hydrogen stream and then to secondary carburization at 1 5000C for 1 hour in a hydrogen stream. It was found by X-ray diffraction that the resulting powder had a crystal structure of simple hexagonal WC type and by measurement of the particle size thereof that it had a mean particle size of 4.0 microns.

This carbide powder was mixed with Ni powder and Co powder in such a proportion that the final composition be (Mo0.45W0.45Cr0.10)C-30% Ni-I 5% Co and then sintered at 1 2200C in a vacuum of 10-2 Torr. The thus obtained alloy had a structure consisting of two phases of MC and binder metals.

When this alloy was polished with a diamond paste to give a mirror surface and subjected to a test by immersing in an artificial sweat for 24 hours, there was hardly found corrosion thereof.

The alloy obtained in this example is suitable for use as a watch frame because of its light weight as well as excellent scratch proofing property.

Example 5 Heading dies were made of the Alloy Sample Nos. 1, 9, 12 and 14 or Example 3 and WC25% Co alloy for comparison and the life tests thereof were carried out by subjecting to cold forging of bolts of S45C steel, thus obtaining results shown in Table 4: Table 4 Tool Life (Number of Bolts Processed/Die) x 10~5 Alloy Sample No, 0 2 4 6 8 1 x 9 x 12 x 14 x WC25% Co x Example 6 910 g-of the solid solution (M0.7W0.3) prepared in Example 1 was mixed with 90 g of carbon powder and 3 g of cobalt powder as a diffusion aid and then subjected to (1) carburization at 1 4500C for 1 hour in a nitrogen stream or (2) carburization at 1 3500C for 1 hour in a nitrogen stream.It was found by X-ray diffraction that both of the resulting carbides were uniformly of (Mo0,7W0.3)C and measurement of the particle size of these carbides by means of Fisher Sub Sieve Sizer showed that the carbide prepared by the process (1) had a mean particle size of 4.2 microns, while the carbide prepared by the process (2) had a mean particle size of 1.9 microns.

Each of these carbides was mixed with 30% of cobalt powder, ball milled with alcohol medium passed in a desired shape and then sintered.

Heading dies were made of the thus obtained alloys and the life tests thereof were carried out by subjecting to cold forging of bolts of S45C steel, thus obtaining results shown in Table 5: Table 5 Tool Life (Number of Bolts Processed/Die) x 105 Alloy Sample No. 0 2 4 6 8 Alloy from x Process 11) x Alloy from x Process (2) x Example 7 Piercing punches for punching a steel plate of 5 mm in thickness were made of the Alloy Sample Nos. 7 and 1 5 of Example 3 and a WC-1 2% Co alloy for comparison and used therefor 100,000 times. The amounts of wear of the piercing punches at that time are shown in Table 6: Table 6 Alloy Sample No.Amount of wear (mm) 7 0.07 15 0.08 WC-12%Co 0.21 Example 8 68.5% of an (MoO.7W0.3)C powder with a particle size of 3 microns, 30% of Ni powder, 1% of Mn powder and 0.5% of Re powder were mixed while adjusting the quantity of carbon to 97% of the theoretical quantity 6.10%, and the mixed powder was sintered at 1 2500C for 1 hour in a vacuum of 10-2 Torr.The resulting alloy was non-magnetic and had the following properties: Density: 9.9 g/cm3 Hardness (HRA): 84.5 Transverse Rupture Strength: 290 Kg/mm2 Example 9 75% of an (Mo0.5W0.5)C powder with a mean particle size of 4 microns, 10% of Ni powder, 13% of Co powder, 1% of Re powder, 0.8% of Mg powder and 0.2% of B powder were mixed while adjusting the quantity of carbon to 98% of the theoretical quantity of carbon 5.93%, and the mixed powder was sintered at 1 3500C for 1 hour in a vacuum of 10-2 Torr.The resulting alloy had the following properties: Density: 10.1 g/cm3 Hardness (HRA): 86.5 Transverse Rupture Strength: 265 Kg/mm2 For examination of the corrosion resistance, the above described alloy and a WC20% Co alloy for comparison were subjected to tests using various corrosion solutions to give results as shown in Table 7: Table 7 Amount on Corrosion (mg/cm2/hr) 1* //** III*** Alloy of Our Invention 2.5 0.2 0 WC20% Co 2.8 2.1 0 Note:I Hot 10% H2SO4 Solution II 35% HCI Solution at Room Temperature 11110% NaOH Solution at Room Temperature Example 10 30% of an (Mo0.7W0.3)C powder with a mean particle size of 5 microns, 35% of an (MoO.3WO 7)C powder with a mean particle size of 0.5 micron, 25% of Ni powder and 10% of Co powder were mixed while adjusting the quantity of carbon to 97.5% of the theoretical quantity of carbon, i.e.5.15% and the mixed powder was sintered at 1 3200C for hour in a vacuum of 10-2 Torr. The resulting alloy had the following properties: Density: 11.2 g/cm3 Hardness (HRA): 82.5 Transverse Rupture Strength: 280 Kg/mm2 Heading dies for nut former were made of the above described alloy and a WC-25% Co alloy for comparison and the life tests thereof were carried out by cold forging nuts of Si SC steel, thus obtaining results shown in Table 8: Table 8 Tool Life (Number of Nuts Processed/Die) x 105 Alloy Sample No. 0 2 4 6 8 10 12 Alloy of our x (12.0 Invention WC25% Co x(4.2)

Claims (14)

Claims

1. A hard alloy comprising a hard phase consisting of at least one compound having a crystal structure of the simple hexagonal MC type (M: metal; C: carbon) which is a mixed carbide, carbonitride or carboxynitride of molybdenum and tungsten, and a binder phase consisting of at least one of iron, cobalt and nickel, the hard phase being prepared by carburizing an (Mo, W) alloy obtained by reducing oxides of molybdenum and tungsten having a particle size of 1 micron or less and comprises coarse particles having a mean particle size of at least 3 microns and a uniform molybdenum to tungsten ratio and having a gross composition within the shaded portion ABCDEA of Figure 1 of the accompanying drawings.

2. A hard alloy as claimed in claim 1 wherein the hard phase having a mean particle size of at least 3 microns is obtained by controlling the carburizing conditions.

3. A hard alloy as claimed in claim 2 wherein the carburizing conditions are controlled by adjusting the carburization temperature to a temperature which is as high as possible but lower than the decomposition temperature of (Mo, W)C to (Mo,W) > C.

4. A hard alloy as claimed in claim 1 wherein the hard phase having a mean particle size of at least 3 microns is obtained by subjecting the (Mo, W) alloy to a heat treatment.

5. A hard alloy as claimed in claim 4 wherein the heat treatment is carried out at a temperature of from 1100 to 1 4000C in a stream of nitrogen or hydrogen.

6. A hard alloy as claimed in claim 1 wherein a part of the MC type compound is replaced by at least one hard BI-type compound containing Ti,Zr, Hf,V,Nb,Ta,Cr,Mo or W.

7. A hard alloy as claimed in claim 6 wherein the quantity of the BI type hard compound replaced is up to 30% by weight.

8. A hard alloy as claimed in claim 1 wherein at least one of the mixed carbides is a solid solution of (Mo,W,Cr)C.

9. A hard alloy as claimed in claim 8 wherein the quantity of Cr is from 0.3 to 10% by weight.

10. A hard alloy as claimed in claim 1 wherein a part of the carbon in the carbides forming the hard phases in is replaced by nitrogen and/or oxygen.

11. A hard alloy as claimed in claim 10 wherein the quantities of nitrogen and oxygen are defined, in connection with the alloy composition, by the relationship of: W atom % N atom % 0.005 < -( x ( )x( Group IVa,Va,Vla metal atom % Group IVa,Va,Vla metal atom % and W atom q/o 0 atom % 0.005(1-- atom % )x( ) < 0.05 Group IVa,Va,Vla metal atom % Group IVa,Va,Vla metal atom %

1 2. A hard alloy as claimed in claim 1 which includes at least one of Be, Mg,Ca, B,Si,P,Mn,Fe or Re.

13. A hard alloy as claimed in claim 1 wherein at least one of Mn, Re,Cu,Ag,Zn, and Au is incorporated in the binder phase to make the alloy non-magnetic.

14. A hard alloy as claimed in claim 1 wherein the hard phase comprises two or three simple hexagonal phases differing in their ratio of Mo/W.

1 5. A hard alloy as claimed in claim 1 substantially as hereinbefore described.

1 6. A hard alloy as claimed in claim 1 substantially as hereinbefore described with reference to any one of the Examples.

1 7. A wear resisting tool which is formed from an alloy as claimed in any one of the preceding claims.