CA1115994A - Hard alloy containing molybdenum and tungsten - Google Patents

Abstract

TITLE OF THE INVENTION
A hard alloy containing molybdenum and tungsten ABSTRACT Of THE DISCLOSURE
This invention relates to a hard alloy comprising a hard phase consisting of at least one compound having a crystal structure of simple hexagonal MC type (M: metal;
C: carbon) selected from the group consisting of mixed carbides, carbonitrides and carboxynitrides of molybdenum and tungsten as a predominant component, and a binder phase consisting of at least one element selected from the group consisting of iron, cobalt, nickel and chromium, in which a hard phase consisting of a compound of M2C type having a crystal structure of hexagonal type is evenly dispersed.

Description

BACKGROUND OF THE INVENTION
1. FIELD OF THE INVENTION
This invention relates to a hard alloy containing molybdenum and a process for the production of the same and more particularly, it is concerned with a hard alloy comprising, as a predominant component, a hard phase con-sisting of a compound having a crystalline-structure of simple hexagonal type and a process for the production of the same.

2. DESCRIPTION OF THE PRIOR ART
Up to the present time, as a starting material for cemented carhides, there has been used tungsten car-bide (WC) powder as a predominant component with a suitable binder metal, typically an iron group metal, to which carbides or carbonitrides of high mel~ing point metals such as titanium tTi), tantalum ~Ta), niobium (Nb), molyb-denum ~Mo), hafnium (Hf), vanadium (V) and chromium (Cr) are added depending upon the requirements of a desired alloy. However, it is also true that tungsten is a relatively expensive metal and that it is found in only a few parts of the world. Accordingly, it is considered to be a so-called "strategic" material, and its availability can be subject to political considerations. Therefore, increase of the demand for cemented carbides consisting of tungsten carbide mainly meets inevitably with a problem of natural resources and if the tungsten carbide can be exchanged for another high melting point metal carbide, this exchange has a great influence upon the industry.
Molybdenum monocarbide (MoC) is considered as a useFul substitute, since this carbide only has the same L5~

crystal structure of simple hexagonal type as tungsten carbide as well as the similar mechanical properties to tungsten carbide. However, the existence of the hexa-gonal molybdenum monocarbide as a simple substance has remained in question to this date and thus an attempt to stabilize molybdenum monocarbide has exclusively been carried out by forming a solid solution with tungsten carbide. This method was firstly reported by W. Dawihl in 1950, bu~ this solid solu~ion was not examined in detail and the commercial worth was not found in those days.
Of late, however~ the study to utilize the solid solution (MoxWy)C wherein x ~ y = 1 has become active with the rise of the price of tungsten. It is very interesting why a study on this solid solution and an attempt to use the same has not been carried out so actively up to the present time.
Molybdenum carbide is stabili~ed as a monocar-bide having a crystal structure of simple hexagonal type when a solid solution is formed with tungsten carbide.
If this stable carbide of (Mo, W)C can readily be pre-pared, replacement o-f tungsten by molybdenum would be possible. For the embodiment of this purpose, there has been proposed a process for the stable production of (Mo, W)C (Japanese Patent Application OPI No. 146306/1976 - U.S. Patent No. 4,049,380 - ). When the ~Mo, W)C powder obtained by this process is used as a starting material of a (Mo, W)C-CO alloy as a substitute for WC, however, MoC is not stable in the alloy and Mo2C tends to precip-itate often. Furthermore~ this process has not been put to prac~ical use, since it requires a heat treatment for a long time.
Furthermore, it has been proposed to produce a molybdenum-tungsten carbonitride having a crystalline structure of tungsten carbide by heating molybdenum and tungsten in combined form and carbon in a proportion sufficient to form the monocarbide in a nitrogen-containing atmosphere (Japanese Patent Application (OPI) No. 104617/
1978.) However9 this method aims at stabilizing the alloy by incorporating nitrogen 50 that (Mo 9 W) 2C is not pre-cipitated.

SUMMARY OF THE INV~NTION
It is an object of the present invention to pro-vide a hard alloy containing molybdenum.
It is another object of the present invention to provide a hard alloy corresponding to a cemented carbide alloy consisting mainly of tungsten carbide (WC) a part of which is replaced by molybdenum carbide (MoC~.
It is a further object o the present in~ention to provide an alloy having a hard phase consisting of a hexagonal monocarbide of (Mo, W)C in which (Mo, ~)2C is dispersed.
It is a still further object o the present inven-tion to provide a hard alloy having 2 hard phase consist-ing of a cornpound of tungsten and molybdenum with carbon, nitrogen and oxygen, having a crystal structure of simple hexagonal type.
These objects can be attained by a hard alloy comprising a ha~d phase consisting of a compound of ~Mo9W)C
of simple hexagonal type and a binder phase consisting ~ 5~

of at least one element selected from the group consisting of iron, cobalt, nickel and chromiumg in which a compound represented by ~Mo 7 W) 2C having a crystal structure of hexagonal type is uniformly dispersed as a hard phase.

BRIEF DESCRIPTION OF THE DRAWING
The accompanying drawings are to illustrate the principle and merits of the present in~ention in more detail.
Fig. 1 is a micrograph, magnified 200 times, of a prior art cemented carbide alloy containing molybdenum, showing the appearance of a carbide of (Mo, W)2C type precipitated needlewise.
Fig. 2 is a micrograph, magnified 200 times, of a (Mo, W)C-(Mo, W)2C-Co alloy according to the present invention, in which a carbide of (Mo, W)2C type is uni-formly dispersed.
Fig. 3 is an X-ray diffraction pattern of an alloy of the present invention.
Fig. 4 is a graph comparing the high temperature hardness of a WC-Co alloy according to the prior art and a (Mo, W, Cr)C-Co alloy according to the present invention, in which A shows WC-10 % Co, B shows (Mo, W, Cr)-(9 % Co ~ 5 % Ni), C shows WC-15 % Co and D shows (Mo, W, Cr) 15 % Co.
Fig. 5 is a graphical representation of the com-pressive stress and tensile stress of a WC-Co type alloy as a function of the strain, in which A = WC-5%Co, B = WC-10%Co, C = WC-25%Co, D = WC-30%Co, E = WC-7%Co, F = WC-12%Co and G = WC-15%Co. The percent of Co is by weight.

Fig. 6 is a graph comparing the compressive stress as a function of the strain of a (Mo, W~C-Co alloy accord-ing to the present invention and a WC-Co alloy of the prior art, in which H = WC-11%Co, I = WC-16~Co, J =
(Mo0 7W0 3)C-11%Co, K = (Mo0 5W0 5)C-19%Co and L = WC-24%Co.
The quantity of the binder metals Co and Ni is by volume percent.

DETAILED DESCRIPTION OF THE INVENTION
In accordance with the present invention~ there is the provision of a hard alloy comprising a hard phase consisting of at least one compound of simple hexagonal MC type ~M: metal; C: carbon) selected from the group consisting of mixed carbides, carbonitrides and carbooxy-nitrides of molybdenum and tungsten as a predominant com-ponent, and a b;nder phase consisting of at least one ele-ment selected from the group consis*ing of iron, cobalt, nickel and chromium, in which a hard phase consisting of a compound of M2C type having a crystal structure of hexa-gonal type is evenly dispersed. The quantity of the compound of M2C type precipitated in preferably 3G % by volume or less, in particular, 5 to 25 ~ by volume, since if more than 30 % by volume, M2C grows to large particles such that the objects or effects as a dispersed strength-ened alloy cannot be achieved. As an essential condition for dispersing evenly the hard phase of M2C type, it is necessary that the carbon content of the hard phase is in an atomic proportion of 0.98 to 0.80 to the theoretical carbon content of the hard phase of MC type.
Briefly stated, and in accordance with the pre-sently preferred embodiment of the invention, there is provided a hard alloy which comprises one or more carbide phases consisting of 80 % by weight or more of a carbide of MC type, solid solution containing molybdenum and tung-sten and having a crystal structure of simple hexagonal type and 20 % by weight or less of a mixed carbide of M2C
type containing, as a main component, Mo2C and having a granular or globular shape with a size of 10 microns or less, the carbide of M2C type being dispersed in the alloy, and 3 to 50 % by weight of a binder phase consisting of an iron group metal.
The inventors have made various studies mainly on the relation between the carbon content and the tough-ness of a cemented carbide alloy from (Mo~Wy)C as a start-ing material and consequently, have found the following facts:
When the carbon content of the alloy is less than the theoretical carbon content, a molybdenum-tungsten mixed carbide of M2C type precipitates as needle crystals as in the prior art (Fig. 1). When the alloy contains a micro quantity of an impurity element5 on the other hand, the mixed carbide is not of needle crystals~ but it pre-cipitates in a finely granular -form ~Fig. 2). It is found as a result of measurement of the stTengths of the alloy containing needle-shaped molybdenum-tungsten mixed carbide of M2C type precipitated and the alloy containing granular mixed carbide of M2C type that the latter is superior to the ~ormer in toughness. It is well known that the strength of such an alloy depends to a great extent on the di-Efer-ence of precipitated forms as described above. In the former case, a stress is concentrated on needle crystals ~ ~r of molybdenum-tungsten mixed carbides (M2C~ M3C2) resulting in starting points of breakage and in lowering of the s~rength of the alloy~ while the latter is a so-called dispersed type alloy, in which granular molybdenum-tungsten carbides of M2C type are evenly or widely dispersed so that the stress concentration on the mixed carbides is prevented and an external force added to the alloy is rather absorbed, thus increasing the strength of the alloy.
The reason why the granular crystals of molybdenum-tungsten mixed carbide of M2C precipitate is not clear in detail, but can be considered to be as follows:
In an ordinary (Mo, W)C alloy, precipitation of the needle crystals of molybdenum-tungsten mixed çarbide of M2C type is due to that the precipitation temperatures of tungsten and molybdenum are different in the cooling step of the liquid phase in which tungsten, molybdenum and carbon are dissolved in the binder phase. That is to say, WC precipitates at a relatively high temperature and only Mo and C remain in the binder phase to the end. At a temperature o lower than 1180 C, MoC is decomposed into Mo2C and C and thus Mo2C remain as an agglomerate.
The free carbon and Mo2C precipitated can well be dispersed by a rapid cooling treatment so as to prevent them from agglomeration. However, this method can with a suf-Eiclent cooling effect be adapted to a small-sized alloy having a small thermal capaclty, but it is difficult to treat a large-sized cemented carbide alloy having a large thermal capacity by this method The inventors have found as a result of X-ray diffraction analysis of the binder phase of an alloy in which needle Mo2C is precipitated that the lattice con-stant of the binder phase is not changed frcm ~hat of the pure metal and the binder phase is not alloyed, nor embrittled. Thus, it is assumed that if a needle-shaped Mo2C precipita~ed can be dispersed in a granular or globular form, an alloy having a sufficient strength can be prepared. If there are micro amolmts of impurity elements in the binder phase, Mo, W and C are combined on nuclei of such elements to form or precipitate a number of nuclei of M2C type molybdenum-tungsten mixed carbide before the liquid phase vanishes or solidifies and Mo and C are not in the binder phase. Thus~ there is no precipi-tation of needle crystals of M2G type even at a tempera-ture of 1180 C or less at which the liquid phase ~anishes~
In general, the precipitate of a large number of M2C
nuclei is in a globular or rod-like form and, in order to disperse and precipitate more -finely the molybdelaum-tungsten mixed carbide of M2C type in the alloy, it is effective to inhibit the precipitation and growth of the molybdenum-tungsten mixed carbide of M2C type by subjec:t-ing to rapid cooling from the sintering temperature to the solidification temperature of liquid phase.
For the purpose of dispersing evenly the ~2C
phase in the alloy, there are a method comprising prepar-ing previously a carbide (Mo, W)2C, adding to starting powders to be mixed and controlling the sintering con-ditions to precipitate uni-formly ~Mo, W)2C phase, a method comprising, during the step of producing a carbide, synthesizing not only a complete solid solution of ~Mo, W)C but also a carbide in the surface layer of which fine (Mo, W}2C is dispersed, adding an iron group metal such as Co, Ni or Fe to the carbide and sintering the mixture with precipitation of (Mo, W)2C and a method com-prising adding Mo and W to a carbide of (Mo, W)C and thus precipitating Mo and W dissolved in the binder phase as (Mo, W)2C during the sintering step.
Furthermore, the inven~ors have made s~udies on the conditions for dispersing (Mo9 W~2C in the alloy and consequently, have found that a micro amount of one or more impurity elements is added to the alloy and the M2C
phase is precipitated round the impurity nuclei in the steps of sintering and cooling, thereby dispersing the M2C phase uniformly in a globular form. In particular~
during formation of the carbide, an impurity element can be added and dispersed uniformly. Impurities such as iron are effective for promoting the carburization reao tion and ~e3C formed at this time serves as nuclei to disperse (Mo, W)2C. When the impurity is not added, a needle-like M2C phase tends to precipitate as a primary crystal. In order to prevent this precipitation, it is necessary to control the quantity of Mo and W dissolvecl during sintering. To this end, the quantity of W dis-solved in the binder metal is increased more than that of Mo thereby precipitating uniform (Mo, W)2C.
Examples of the element added as an impurity ele-ment to the binder metal are one or more of berylliwn, magnesium, calcium, boron, silicon, phosphorus, manganese, iron and rhenium. These elements are added individually or in combination to the binder metal in a proportion of at most 3 % by weight, since if more than 3 % by weight 9 the molybdenum-tungsten mixed carbide of M2C phase is embrittled and the strength is not so increased. Addition of titanium, zirconium, hafnium, tantalum and niobium as an element to inhibit the precipitation and growth of the molybdenum-tungsten mixed carbide of M2C type is also effective for the dispersed precipitation of the mixed carbide.
The carbide or mixed carbide of M2C mentioned in this specification includes not only (Mo, W)2C and ~Mo 9 W) 3C2 but also other lower carbides containing other metals.
The size of the granular precipitate 9 molybdenum-tungsten mixed carbide of M2C type is preferably 0.1 to 10 microns, more effectively 1.0 to 2 microns, since if the precipitated particles are too coarse, the strength and hardness of the alloy are lowered, while i-f too small, the mixed carbide is deposited on the boundary of (Mo, W~C
or binder phase thereof, so that the boundary strength is is lowered and thus the alloy strength is deteriorated.
The quantity of carbon when a carbide of M2C type is dis-persed in the alloy is preferably 80 to 98 % of the theoretical quantity when all the carbides are regarded as of MC type. This corresponds to the presence of 2 to 30 ~ by volume.
A granular or globular molybdenum-tungsten mixed carbide of M2C phase has a large influence upon the pro-perty of the alloy depending on the quantity of the mixed carbide. X-ray diffraction using CuKa, under conditions of 40 KV, 80 mA~ FS 4000 c/s and TC 0.2 sec shows that the alloy has properties at least similar to those of WC-Co -1~-alloys when the ratio of the X-ray peak of M~C type to the peak of MC type appearing near 39.4 and 48.4 D in the X-ray diffraction angle (2 a) i5 in the Tange of 0.01 to 0.5, in particular, 0.05 to 0.20. In this specifica~ion~
for the convenience of illustration, ~he molybdenu~-tungsten mixed carbide of M2C type is sometimes represented by M2C or Mo2C, but, even though W, Co, Ni, N and/or O
are dissolved in M2C or Mo2C and the ratio of the metallic components and non-metallic components is fluctuated near 2 : 1, the effects o~ merits of the present invention are not lost.
As a binder metal there is preferably used an iron group metal in a proportion of 3 to 50 % by weight based on the alloy composition~ since if less than 3 %
by weight, the alloy is bri~tle and if more than 50 % by weight, the high temperature property is deteriorated.
The iron group metal as a binder phase can naturally dis-solve Group IVa, ~a and VIa metals and it is possible to add even other elements having solubility therein such as aluminum, silicon, calcium, silver9 etc. with holding the merits of the present invention.
The basic concept of the present invention can be held even when a part of molybdenum and tungsten carbide is replaced by a Bl ~ype mixed carbide containing titanium, zirconium, hafnium, vanadium, niobium, tantalum, chromium "nolybdenum and/or tungsten in a proportioll of 30 % by weight or less, preferably 0.5 to 25 ~ by weight.
Furthermore 7 there is the similar relationship even in the case of an alloy wherein a part of C in the carbide is re-placed by nitrogen and/or oxygen. Examples of the preferred embodiment in this case are as follows.
The first embodiment is incorporation of N in (W, Mo)C to give ~W, Mo)(C, N) whereby a stable star~ing material of hexagonal WC type can be obtained without a heat treatment for a long time.
The second embodiment is incorporation of O in ~W,Mo)(C,N) to give (W9 Mo)(C, N, 0) which is more stable.
The third embodiment is incorporation of Cr in ~W, Mo)(C, N) or (W9 Mo)~C, N, 0~ to give (W, Mo, Cr3(C, N) or (W9 Mo, Cr)(C, N, 0) whereby a starting ma~erial with a low weight and low price can be obtained.
The fourth embodiment is that in the production of these starting material powders, 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 700 C or higher in a part of the carburization step to -form a stable starting powcler.
The fifth embodiment is that, when the above des-cribed starting powder is combined with an iron group metal, two or more kinds of hard phases of simple hexa-gonal 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 Bl type solid solution containing one or more of Group IVa, Va and VIa metals and non-metallic elementsl or the ordinary additives to cemented carbides, such as silver, silicon, bismuth, copper, aluminum, etc. can also be added to the iron group binder metal with holding the merits of the present invention.

~5~

The above described embodiments will now be illus-trated in greater detail:
In the important system of the present invention wherein there are a simple hexagonal phase containing molybdenum and tungsten and an M2C phase, it is found in the sintered alloy with a binder metal that, when A = ( N atom % ) x ~1 W atom % ), the suitable ~ Mo+W) atom % ~Mo +W) a~om ~
range of A is 0.005 < A < 0.5. If A is less than the lower limit, the ef-fect of nitrogen does not appear, while if more than the uppeT limit, sintering is difficult to give excellent properties. The most suitable range of A is 0.01 < A _ 0.4.
Concerning the effect of oxygen) it is found that, h B~=~( atom ~ ) x (1 - W atom % _), the (Mo+W) atom % (Mo~W) atom %
suitable range of B is 0.005 < B < 0.05. I 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 0O01 < B < 0.04.
On the other hand, a W/Mo ratio ls preferably 5/95 to 90/10, since if less than 5/95, the alloy is unstable, while îf more than 90/10, the merits of the replacement (light weigh~, 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 i-f more than 0.5, the alloy is brittle although the corro-sion resistance is increased.
As well known in the art, it is advantageous for cutting tools to form a Bl type solid solution composed of at least one of Group IVa, Va and VIa metals such as ~L~D~9 4 titanium, zirconium, hafnium, vanadium, tantalum, chromium, molybdenum and tungsten with at least one of non-metallic components such as carbon, nitrogen and oxygen in addition to the simple hexagonal phase. The quantity of the Bl type solid solution is preferably changed depending upon the cutting use.
Concerning the quantity of nitrogen in this case, it is found as a result of our ~arious experiments that, when the definition of A is changed to N atom %
( - -) x (1 -- Group IVa,Va,VIa metals atom %
W atom % ), the suitable range of Group IVa, Va, VIa metals atom ~
A is also 0.005 < A < 0.5 although a part of the nitrogen is occluded in the Bl type solid solution. The optimum range of A is 0.01 < A < 0.4. Concerning the quantity of oxygen, it is found as a result of our various experi-ments that, when the definition of B is changed to O atom %
) x Group IVa,Va,VIa metals atom ~
(1 - W atom % _ ), the suitable range Group IVa,Và,VIa metals atom %
of B i5 also 0.005 < B < 0.05. The optimum range of B
is 0.01 c B < 0.04.
- As the binder metal, there is preferably u-sed an iron group me*al in a proportion of 3 to S0 % by weight based on the gross composition, since i less than 3 % by weight~-the alloy is brittle and if more than 50 ~ -by weight, the alloy is too soft.
For the preparation of starting materials, the - reaction is carried o~t at a high temperature- in a hydrogen atmosphere in the case o carburization of a (Mo, W) powder with carbon, reduction and carburization of oxide powders with carbon or combina~ion thereof. At this time, it is ~ ~ ~ 5 ~ ~ ~

found as a result of our studies on the decomposition nitrogen pressure of (Mo~ W)(C, N) that the external nitro-gen pressure, depending on ~he temperature, should be 300 Torr or moTe at 700 C or higher at which the carbonitriza-tion reaction takes place. The coexistence of hydrogen is not always harmful, but lt is desirable to adjust the quantity of hydrogen to at most two times as much as that of nitrogen, in particular, at most the same as that of nitrogen not so as to hinder the nitriding reaction. In the case of using an ammonia decomposition gas, it is necessary to enrich with nitrogen.
For the preparation of starting materials contain-ing oxygen, the coexistence of carbon monoxide and carbon dioxide is required in an atmosphere. In this case 9 the quantity of hydrogen is not limited as described abo~e, but should not exceed 50 % of the atmosphere. Heating and sintering in an atmosphere of nitrogen or carbon oxide is effective for the purpose of preventing an alloy sintered from denitrification or deoxidation.
In ~he abo~e described five embodiments, the dis-persing treatment of an M2C type phase can be omitted and in this case, considerably excellent effects can also be given.
In cemented carbide alloys consisting predominant-ly of WC, excellent properties as alloy uses such as drills, hubs, taps, etc. can be obtained by reducing the particle size of the carbide when containing a binder metal in a proportion of up to 15 % by weight, but, when the alloy contains a binder metal in a higher proportion, this procedure has no effect. In alloys for low speed cutting, for example, drills, in particular$ the edge portion is deformed by frictioII heat.
The inventors have further made studies to develop an alloy having a higher wear resistance and toughness and consequently, have found that the deformation a-t a high temperature can remarkably be improved by changing tungsten carbide to a carbide composed of a solid solution of three elements, molybdenum, tungsten and chromium. That is to say, a (Mo, W)C-Co alloy has a higher hardness at a high ~emperature than a WC - Co alloy and~ when Cr is further dissolved in this carbide, the hardness is further ~aised and the high temperature hardness is also improved.
Thus, the disadvantages of the prior art WC-Co alloy can be overcome by on~ effort ~Cf. Fig. 4). It is 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 adding merely chromium to the binder phase has the disadvantages that it is im-possible to make finer the carbide and the carbide phase is not stabilized as a monocarbide of a_solid solution of (Mo, W, Cr). The quantity o~ chromium to be added to the solid solution carbide (Mo, W)C ranges preferably 0.3 to 10 %, since if less than 0.3 %, the carbide cannot be made finer, while ;f mvre than 10 %, Cr3C2 is separated and precipitated in the alloy, resulting in lowering of the hardness.
In a further embodiment of the present invention, a part o the carbon in the solid svlution carbide (Mo, W, Cr)C is replaced by nitrogen, oxygen and/or hydrogen.
That is, it is assumed that if the carhon contained in (Mo, W, Cr)C is added as solid and reacted with a reactivity of lOU %, the crystal is stabilized, but now it is found that incorporation of not only carbon but also nitrogen results in stabilization of the monocarbide as (Mo, W, Cr) (CN) and further incorporation of oxygen and hydrogen stabilizes more 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 M~C type mixed carbide precipitates needle~wise to thus lower the strength.
When the quantity of chromium contained in (Mo, W, Cr)C is`limited to 0.3 to lD % by weight to thus obtain a finer carbide and one or more of iron group metals such as iron~ nickel and cobalt are added as a binder phase in a proportion of 15 to 30 % by weight, the so obtained alloy can be used as a cemented carbide alloy for low speed cutting, for example, drills 9 taps and hubs with an excellent performance. When the binder metal is within a range of 3 to 15 % by weight, the alloy can also be used effectively as a corrosion resisting alloy.
Useful examples of the corrosion resisting alloy are corro-sion resisting seal rings, watch frames, ends of slide calipers~ mechanical seals, etc.
As a material for a cemented carbide alloy there is chosen an alloy having a relatively large cobalt con-tent, which deformation strength is high. As shown in Fig. 5; breakage is hard to occur by deformation with the increase of the quantity of coba]t. If the quantity of ~ 3~

cobalt is increased, however, the alloy shows a decreased yielding stress and tends tQ be deformed. This tendency of deformation is a disadvantage in the case of using the alloy as a forging tool such as headers, although it is hardly cracked.
In accordance with the present invention, there is provided an alloy having a high deformation resistance as well as a sufficient elastic strength without lowering the hardness and it is expected that the properties are more improved than those of the prior art WC-Co type alloys.
That is to say, as a result of our detailed studies on WC-CQ type alloys, it is found tha~, in the WC-Co type alloys, an alloy in WC-~ ~one between the free carbon precipitating zone and &-phase (n-Co3W2C phase) precipitating zone is excellent in mechanical properties and thus alloy in WC-~
zone have mainly been used. Thîs y-phase is a phase such that tungsten is dissolved in cobalt and~ as well kno~n, the alloy property is changed with the change of the quantity of this solid solution. The deformation strength depends on the quantity of tungsten dissolved in the binder phase cobalt. If none is dissolved in the cobalt, the deforma-tion resistance of the alloy is considered to be increased further, but this i5 unreasonable unless free carbon is precipitated.
The inventors have made efforts to find an alloy in which the binder phase is held as pure as possible without the coexistence of free carbon. That is to say, one aspect of the present invention consists in that the deformation strength of the binder phase can be held high without sub-stantial dissolving of tungsten and molybdenum in the binder -1~ -s~

phase even if the quantity of carbon is changed and, in addition, an M2C type compound occurring due to the lack of carbon is evenly dispersed to prevent stress concentra-tion. In general, alloys comprising carbides of molybdenum and tungsten have not been put to practical use because of precipitation of a needle carbide (Mo, W)2C which causes a marked decrease of the alloy strength. The inventors, however, have succeeded in increasing the deformation resistance of the alloy without deterioration of the strength thereof by dispersing well ~Mo, W)2C.
Fig. 6 is a graph comparing the compressive stress as a function of the strain of a (Mo, W)C-Co alloy accord-ing to the present invention and a WC-Co alloy of the prior art. It is fownd that the prior art WC-Co type alloy shows a strain of about 2 to 4 ~ at compression9 whilst the alloy of the present invention shows a strain of 4 to 5 ~. For example, a WC-24 % by volume Co alloy shows a yielding stress of 400 Kg/mm2 and a deformation of about 4 ~, while, on the contrary, the alloy of the present invention exhibits a higher yielding stress, i.e.
500 Kg/mm2 and a deformation amounting to about 5 %.
In the alloy of the present invention, in which the composition ratio of molybdenum and tungsten is repre-sented by (MoxWy)C, the composition of (MoxWy)C in the alloy is not always limited to one, but two or more com-binations can be used to change the alloy property. In this case, a carbide of M2C type, i.e. (Mo, W)2C should be uniformly dispersed to give a desired effect. The advantages of the present invention can be held even when the alloy contains carbides, nitrides and carbonitrides of Group I~a, Va and VIa metals. Furthermore, replacement of C in (Mo, W)C or ~Mo, W)zC by N, 0 and/or H can be carried out with keeping the advantages of the present invention. These alloys having~ in particular, an excel-lent shock resisting toughness can favourably be used as can-maXing tools, dies, mining tools, rolls, etc. in addition to headers well-known as a shock resisting tool.
In a still further embodiment of the present invention, one or more of manganese, rhenium, copper, silver, zinc and gold are incorporated in the binder phase to change the microstructure of the binder phase and to make non-magnetic. At the same time, it is found that, when these elements are added, the binder phase i5 alloyed, whereby the corrosion resistance of the alloy is improved. The hardness and wear resistance of the alloy are deteriorated if the quantity of the binder phase exceeds 30 % by weight and the wear resistance of the alloy is not lowered unless the quantity of these elements exceeds 5 % by weight.
In the case of WC type alloys, it is desirable, as is well known, to design the alloy in a low carbon alloy so as to improve markedly the non-magnetization and corrosion resistance thereof and if the quantity of the alloyed carbon is less than WC-y phase zone, tungsten is dissolved in the binder phase in a large amount to lower readily the magnetism. In the case of (Mo7 W~C type alloys, on the other hand, the magnetism is hard to be lowered even if it is designed in a low carbon alloy. When using nickel as the binder metal, however, there is obtained a remarkable effect in combination with the above described ~ ~ ~ 5~ ~ ~

additives. In a low carbon alloy, an M2C phase is pre-cipitated and preferentially corrosed because it is relatively base electrochemically as compared with a (Mo, W)C phase. It is found~ however3 that when an M2C
phase is evenly dispeTsed in a proportion of 30 % by volume or less in the alloy, the alloy base is not corrosed and the corrosion resistance as the whole alloy body is rather improved because the M~C phase is in a fine globular form. That is, a low carbon al.loy in which an M2C phase is uniformly dispersed in a proportion of at most 30 %
by volume is desired which is made corrosion resistant and non-magnetic by using nickel as a binder phase and adding at least one of manganese, rhenium, copper, silver, zinc and gold thereto.
In the production of the alloy of the present invention, the alloy i5 unavoidably contaminated with small amounts of impurities such as iron~ cobalt, etc., but as far as the sum of these impurities does not exceed 1 %, the advantages of the present invention can well ~0 be kept.
In a still fuTther embodiment of the present invention, the quantity of iron in the alloy is preferably controlled by the relation of:
Amount of Iron (wt %) 0.001 < --- - < 0.1 Amount of Binder Metals twt %~
In this case, the hard phase consists of (Mo, W)C, and the binder phase consists of Co and Ni, to which Fe is added as additive element. Thus, a carbide of (Mo, W)2C
type is precipitated in a granular form, not in a needle-like form. The quantity of Fe to be added as the additive element is preferably 0.1 to 10 % by weight, since if ~ s~

less than 0.1 ~, the effect of Fe is little, and if more than 10 %, the precipitate of M2C type is too coarse to hold the alloy strength. For the addition of Fe, it can be added to the alloy or the reaction mixture during production of the carbide. The carburization reactivity of the carbi.de can be controlled by changing the quantity of Fe. If the quantity of Fe is less than 0.1 %, the carburization does not proceed sufficiently and, when using the thus resulting carbide for the production o-f an alloy, the carbide of M2C type is hardly dispersed in a granular form in the alloy, while i-f more than 10 %, the carbide is alloyed and grinding thereof is veTy diffi-cult resulting in lowering of the yield of the carbide useful for the production o a hard alloy. A ~Mo, W)C
alloy in which an M~C type carbide is precipitated and dispersed according to ~he feature of the present inven-tion has a high alloy strength as a so-called dispersion type alloy. When the particle size distribu~ion and dis-persed state of the M2C ~ype carbide are not uniform or differ in the interior and exterior portion of the alloy, however, the alloy strength CaTlnot be held so high. In the case of large-sized alloys or alloys having a large content of a binder phase, for example~ Morgan Rolls, super-high pressure pistons, anvils and the like, there is often the following problem. That is to say, where carbon is d.iffused from outside the alloy resulting in a marked difference of carbon contents between the exterior and interior portions of the alloy, or where the cooling speed differs in the surface portion and interior portion of the alloy, the precipitation conditions of molybdenum ~ ~ 5~ ~ ~

and tungsten from the liquid phase are not the same and the shapes and dispersed states of the M2C type carbide are different in the exterior and interior portions of the alloy. In the exterior or surface portion of the alloy, the M2C tends to be coarsened and agglomerated, thus resulting in lowering of the strength.
According to the embodiment of the present inven-tion, this problem can be solved. When using cobalt and nickel as the binder phase and adding iron in a propor-tion of 0 1 to 10 %, the carbide of M2C type is stably precipitated and dispersed independently on the quantity of the binder phase and the shape of the alloy to thus keep a high alloy strength. If the binder phase is of cobalt only, M2C tends to be agglomerated and if it is of nickel only, the hardness and compressive strength of the alloy are lowered. When the quantity of iron in the binder phase is less than 0.1 ~, there is not such a large ef-fect thereof, while if more than 10 %, the corro-sion resistance and strength of the alloy are deteriorated.
In this embodiment also, it is desired that a binder phase comprising an iron group metal as a predominant component is in a proportion of 3 to S0 % by weight of the gross composition, since if less than 3 % by weight, Ihe alloy ls too brittle, while if more than 50 % by weight, the high temperature property is deteriorated. It is also natural that the iron group metal as the binder phase dissolves Group IVa, Va and VIa metals and, moreover, the merits or effects of the present invention will not be lost even by the addition of elements having a solubility therein such as aluminum, silicon, calcium, silver~ etc.

~5~3~

The basic concept of the present invention can be held even when a part of the molybdenum and tungsten carbide i5 replaced by a Bl type mixed carbide containing titanium, zirconium, hafnium, vanadium, niobium, tantalum, chromium, molybdenum and/or tungsten. Furtherrnore, the properties of our alloy are not so changed even if a part of Mo and W in (Mo, W)(CNO) or ~Mo, W)C is replaced by other elements as far as it holds the simple hexagonal structure As is apparent from the aspect oE this embodiment, a micro amount of iron is essential as a stabilizer in a (Mo, W)C alloy or (Mo, W)tCNO~ alloy. As a method of dispersing iron, it îs desirable to add iron during form-ation of the carbide, and to effect the carburization reaction at a temperature of 1500 C or higher in a stabilizing atmosphere of nitrogen or carbon oxide.
A process for the production of (Mo, W)C has hitherto been known which comprises adding a large amount of a diffusion aiding agent such as iron or cobalt to Mo2C and WC and subjecting to reaction at 2000 C or higher (Japanese Patent Application (OPI) No. 146306/
1976~. In this process, iron is added for the purpose of promoting the solid solution forming reaction of WC
and Mo2C.
In the embodiment of the present invention, a small amount of iron can be added when a complete Mo-W
solid solution, (Mo, W) alloy powder is carburized, or when a (Mo, W) oxide is directly carburized. The iron added is.used for the stabilizing reaction at a tempera-ture of 1500 C or higher and has no bad influence upon the carbide.

s~

In this embodiment, a part of the carbon in the carbide can also be replaced by nitrogen and/or oxygen with holdin~ substantially the similar effects.
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 the ratio of Mo/W. -The detailed reason o~ increasing the toughness is not clear, but it is assumed that when ~Moj W~C is separated into two phases, the solution strain of both the phases is lowered to give a higher touglmess than in the case of a single phase.
Since at least an alloy consisting o-f a ~MoxWy)C (y > x) phase having the similar property to that of WC nd a ~lox~Yy)C (x > y3 phase having the similar property to that of MoC has two properties, i.e. toughness of WC and heat and deformation resistance of MoC, this embodiment is advantageous more than when using one kind of (Mo, W)C
only. ~lost 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 a case where the peak of plane ~1, 0, 3~ is separated in two in 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 alk~line solution of a hexacyanoferrate ~III) or by X~ observ3tion.
The ~plicat;on or use range of the alloy of the present invelltion is as follows. For example, the alloy of tl~e presellt in~ention can be used for we~r resisting tools sllcll as guide rollers, hot wire milling rollers, etc., ~

and for cutting tools, because of having a toughness and hardness similar to or more than those of WC-Co alloys.
In particular, when the alloy of the invention as a sub-strate is coated with one or more wear resisting ceramic layers such for example as of TiC, TiN, A1203, cutting tools more excellent in toughness as well as wear resis-tance can be obtained than the prior art tools having WC-Co type alloys as a substrate. As well known in the art~
at this time, a decarburization layer called n-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 the 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 deteri-orating the toughness.
When using the alloy of the present invention as a watch case, it shows more excellent properties as a watch case than WC-Co type alloys~ which are summarized below:
(1) Beautiful brightness can be given when the alloy is specularly finished.
(2) Grinding and polishing workings are possible.
~3) Corrosion resistance is excellent, in parti-cular, for sweat in the case of trinkets.
(4) Mechanical strength is considerably high.
According to the present invention, the deforma-tion resistance of an alloy can be increased without deteriorating the strength thereoE by dispersing well or evenly (Mo, W)2C. A carbide of M2C type itself has a low hardness (Vickers hardness oF Mo2C: 1500 Kg/mm2), but3 when this M2C is dispersed uniformly in an alloy, the alloy can hold a high toughness without lowering as a whole the hardness thereof because the soft M2C can moderate an impulsive force added to the alloy. Because of the high wear resistance and high toughness with the low price, the alloy of the present invention is suitable for spikes for shoes or ice spikes. When a carbide of M2C is suitably or uniformly dispersed in an alloy, the a:Lloy shows a good sliding property on a concrete surface and can absorb a shock from the roughness of a concrete surface.
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 ratios, 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 ~MoO 7Wo 3)C containing 0.2 % of Fe as an addi-tive was used as a starting material. Starting materials were taken by weighing so that the gross composition be (MoO 7Wo 3~Cz-15 ~ Co and ~ in this formula (carbon content in alloy/theoretical carbon content) be 100, 98 and 96 atomic %, mixed by wet process in an organic solvent, dried, compacted and sintered at 1450 C in vacuum. For comparison, the similar procedure was repeated except using a Fe-free starting material. The properties of the ~f~'~L5~

resulting alloys are shown in Table 1:
Table 1 T R S Hardness T C F C. z Fe-containing (Kgjmm2) (~!RA) ( (~
Sample (A) 190 86.8 7.61 0.02 100 (B) 290 87.1 7.43 0.00 98 (C) 260 87.2 7.28 D.00 96 Fe-free Sample ~) 175 86.8 7.59 0.00 100 (E) 140 87.0 7.43 0.00 98 ~F) 120 87.1 7.28 0.00 96 Note: T.R.S. = TransveTse Rupture Strength T.C. = Total Carbon F.C. = ~ree Carbon As can be seen from these results~ the alloys of the pre-sent invention (B) and ~C) are more excellent in toughness than the alloys of the prior art (A~ 9 ~D), ~E) and (F).
In the alloys of the present invention, the alloy strength is not lowered even if the carbon content is less than the theoretical carbon content, while in the prior art alloys, a molybdenum-tungsten mixed carbide of M2C type due to lack of carbon is precipitated as a needle crystal, resulting in lowering of the toughness of the alloy.
Fig. 1 is a micrograph of the prior art alloy {E) and Fig. 2 is a micrograph of our alloy (B).
Example 2 A powdered solld solution of ~MoO gWo 1) with a particle size of 6 microns was mixed with 0.2 ~ of Fe powder and variable amounts of carbon to give a z value as shown in Table 2, subjected to carburization at 1600 C
for 1 hour in nitrogen gas and pulverized. The carbide was heated for 30 minutes in CO gas and stabilized. The resulting carbide was a carbide in which MC and M2C phases f~

were coexistent as shown in Table 2. The carbide was mixed with 10 % of Co and 10 ~ of Ni and sintered at 1300 C. The properties of the so obtained alloys are shown in Table 2:
Table 2 ~Mo W )C Carbide Properties of Alloys 0.9 0.1 z Charpy Impaet z T.C. % 2% N2%M2C vol-% T.R.S. Value (K~-mm) 1.0 10~2~ 0.3 0.2 0 180 0.4 0.95 9.74 0.2 0.05 7 260 0.7 0.9 9.22 0.2 0.05 14 2~0 0 O.S 8.23 0.2 0.03 30 220 0.7 Q.7 7.20 0.1 0.02 45 170 0.3 0.~ 6.17 ~.1 0.01 55 140 0.2 0.5 5.14 0.1 0.01 75 100 0.1 As can be seen from this table 9 the practical alloy strength cannot be obtained in cases where the valu0 of z in (Mo, l~')C is 1.0 or less than 0.7.
.
Ex~mple 3 gO to 90 ~ of a carbide having the theoretical combined carbon content ~MoO 7Wo 3)C wi~h a particle size of 5 microns, 0 to 10 ~ of (MoO 7W0 3)2C with a particle size o 3 microns, 9 to 10 % of Co and 0.1 to 0.5-% of Fe, ~e, Si, B and Be were mixed and alloys were prepared in an ;lnalo~ous manner to Example 1. In the texture of tlle allo) oltained in this way, there was evenly dispersed a carl~ of h~2C type ~Mo/ W)2C as shown in Fig . 2, while in th~ prior art alloy of (Mo, W)C- ~Mo~ W)2C-Co, needle cr~stals ~ere precipitated as shown in Fig. 1. For co~ rison, the properties of these alloys are shown in r;~ble a:

Table 3 Run No-~oO 7W0 3)C(MoO 7Wo 3)2C CoDispersing Agent T.R.S.2 (~) (%) ~%~(%) (Kg/mm ) 1 85 5 9.9 Fe 0.1 260 2 82 8 9.5 Re 0,5 190

3 8~ 2 ~.8 Si 0.2 240

4 88 l 9.9 B 0.1 270 87 3 9.7 Be 0.3 200 7 82 8 10 - ~20 8 89 1 lO - 160 Note: Run Nos. 1-5: Present Invention;
Run Nos. 6-8: Prior Art As evident from-this table~ the alloys of the present invention, in which (Mo, W) 2C i5 dispersed by the addi-tion of an impurity, exhibit a high toughness.
Example 4 t 86 ~ of ~MoO 7Wo 3)C powder with a particle size of 5 microns, 5 % of ~MoO 7Wo 3)2C powder with a particle - size of 2 m~crons, 9 % of Co powder and 0.2 ~ of Fe powder were weighed, mixed by wet process in an organic solvent, dried, compacted and sintered at 1450 C in vacuum, thus obtaining an alloy having a trans~erse rupture strength of 260 Kg/mm2 and a hardness (Hv) of 1400 Kg/mm~. Then, this alloy was subjected to a carburizing treatment to precipitate free carbon within a range of 300 microns from the surface layer and coated with a layer of Tic r double layer of TiC and TiN or double layer of TiC and Al2O3.
For compaIison, a commercially sold WC type nlloy was similarly coated. The so obtained inserts Snml)lc Nos. 1-6 were subjected to a cutting test under the following ~30-conditions (Form No. SNU 432):
Workpiece SCM 3 (HB = 280) Cutting Speed v = 170 m/min Feed f = 0.86 mm/rev Depth of Cut d = 1.5 mm Cutting Time 30 min Test results are shown in Table 4:
Table 4 Sample No. Coating Layer VB (mm~ KT (mm) 10Our Invention No. l TiC 0.16 0.09 No. 2 TiC, TiN 0.14 0~03 No~ 3 TiC, A12O3 0.12 0.02 I~C-t~pe Alloy No. 4 TiC 0.21 0.10 No. 5 TiC, TiN 0.17 0.04 No. S TiC, Al~O3 0.15 0.03 .~ e~-ident from these results, the alloy of the pr~sent in~-~ntion is similar to or superior to the prior art WC-t~pe alloy as to VB ~Elank Wear) and KT (Depth o Crater).
There was found no decarburization layer (n-phase~ .in t3le interface between the substrate of our alloy ~nd the TiC
l~.ver and F.C. ~free carbon) was found within a rall~e of ~on microns directly under the coating layer.
Example 5 Mo2C powder with a particle size of 2 m;~l~ons, I~`C l-o-r-ler with a particle size of 2 microns ;In~ I)o~
~o~ r with Co powder as a diffusion aid wero m;~od so n~
to ~ e a final gross composition of ~Mo0 8Wn 2) ~' t~ 5No 05)1 o and then reacted at 1800 C fclr ~ n ~I;nll~o~

ll ~ 3 ~

in a nitrogen-hydrogen stream having a nitrogen partial pressure of 0.5 atm. X-ray diffraction showed formation of a simple hexagonal crystal o~ WC type.
This powder was mixed wiEh Co powder to give a final alloy composition of (Mo~ 8W0 2)(C, N)-10 % Co, compacted to form a desired shape and then sintered. The sintering was carried out by hea~ing in a v~cuum of 10 2 Torr up to 1000 C and in Co atmosphere under a reduced pressure of 10 Torr from 1000 C to 1400 C. On the other hand, for comparison, an alloy was prepared in a similar manner but using no nitrogen in the step of pro-, ducing ~he carbide and no carbon monoxide in the step of sintering. The results are shown in Table 5:
Table S
Composition of Hard Phase A-Value B-Value Texture __ _ Our Invention ~0.8W0.2)~Co.g3No.o4oo ol~0 98 0.04 0.01 l~C type phase + Co phase Prior Art ~0.SW0.2)(co 976No.0o4)o.98 0.004 - WC type phase + needle M2C type phase + free carbon ~ Co phase The alloy of the present invention is light and excellent in shock resistance as well as high temperature hardness according to test results. Therefore, the alloy of our invention is suitable for various tools, in particular, wear resisting tools.
F.xample 6 A previously prepared solid solution powder of ~Mo0 7Wo 15Cr0 15) with a particle size of 2 microns was mixed with carbon and 0.2 ~ of Fe as a diffusion aiding agent, carburized at 1800 C in hydrogen and then reacted at 1300 C for 30 minutes in a mixed gas of nitrogen and carbon monoxide. The hard phase thus obtained, having a gross composition of (~oO 7~0 15crO.lS)(co~ goNo . 060.01)' was mixed with 9.5 % of a binder metal consistin~ of Co/Ni (1/1) containing a micro amount of Fe and sintered X-ray diffraction showed that the resulting alloy was com-posed of a hexagonal monocarbide of ~Mo, IY, Cr)C and a (Mo~ W, Cr)2C phase with the binder phase. In view of the structure, a granulaT carbide of M2C type was ev~nly dispersed in the alloy.
This alloy has a better corTosion resistance, in particular, for sweat as compared with the prior art WC-Co type alloys and, in addition, it is suitable for use as trinkets such as watch case because of i~s light weight and as non-magnetic alloys.
Example 7 (M0-7wo-3)(co.9No 08) prepar~d by mixing (Mo 7 W) 03 with carbon and 0.05 % of Fe and reacting at 1700 C in nitrogen at 1.1 atm, 30 % of ~Tio 7W0 3) (C0 85No 15)~ 5 % of Co and 5 ~ of Ni was mixed and ball milled bywet process. The thus mixed powder was compacted, then heated in vacuum up to 800 C, in H2 of 30 Torr up to 1200 C and in C0 of 30 Torr at 1200 to 1400 C ~Ind held at 1400 C for 1 hour to finish the sintering~
Analysis of the alloy showed: A = 0.11 and B =
0.02 and examination of the structure showed that ~hcre were a (Mo, W)(CN) phase and a (Mo, W)zC phase dis~-crsed well in a globular form.
For comparison, a nitrogen-free alloy was l~rel~ared by sintering in vacuum only.
These samples were subjected to a cutting test under the following conditions:
Workpiece Ordinary Steel (Hardness: RC 20-29) Cutting Speed 150 m~min Feed 0.381 mm/rev Depth of Cut 0013 mm The results are shown in Table 6:
Table 6 Cutting Time Flank Wear Crater Wear Edge Deforma-(min) (mm) (mm) tion ~mm~
Alloy of Our Invention 15 O.lS 0.010 O~OQ6 Alloy of Prior Art 15 0.20 0.017 0.015 As can be seen from these results, our alloy is more excellent in wear resistance as well as edge deformation resistance than the prior art alloy. When the alloy of the present invention was coated with one or more of carbides, nitrides, oxides and borides in monolayer or multilayer to form a so-called coated insert, the excel-lent edge deformation resistance of the alloy could well be held.
Example 8 MoO3 powder and W03 powder were taken by weighing to give a calculated quantity of Mo/W ratio of 8/2; and mixed with carbon in a proportion sufficient to remove the oxygen in the oxides and 0.2 % of Fe as a catalyst for fixing nitrogen during the reaction. The mixture was reacted at 1500 C for 1 hour in a gaseous stream of (NH3 ~ 10 vol ~ CO) to complete the reducing reaction.
X-ray analysis of the resulting compound showed the formation of a hexagonal type compound of (Mo, I~)~CNO~.

-3~-~ ~ 5 ~

This carbide was mixed w;th 10 % of Co and Ni and an alloy was prepared therefrom in an analo~ous manner to Example 5. Analysis of the resulting alloy showed A = 0.2 and B = 0.03 and examination of the structure thereof showed that a granular carbide of M2C type was evenly dis-persed in a proportion of 2 % by volume. This alloy was particularly excellent in shock resistance.
Example 9 527 g o WC powder with a particle size of 1 micron, 430 g of Mo2C po~der with a particle size of 2 microns and 13 g of Cr3C2 were mixed, to which 5 g of Co powder as a diffusion aid and 27 g of carbon black for filling up the lack of carbon were further added, and ball milled by dry proc~ss for about 30 hours. The thus mixed powder was reacted at 2000 C in a H2 stream to form a primary carbide. The primary carbide was well ground and then subjected to secondary carburization. In the secondary carburi~ation, the carbide was further reacted under each of the ~ollo~in~ conditions, whereby a part of the carbon in the carbide (~lo, 1~, Cr)C was replaced by oxygen, ; nitrogen and hydro~en and the carbide was more s~abilized:
I) in NH3 stream9 1400 C x 1 hour II) in CO ~tream, 1600 C x 2 hours III) in H2 ~tream, 1500 C x 1 hour IV) in ~acuum, 1500 C x 1 hour Variolls carbide~ re obt~ined as shown ;n Iable 7:
T~lle 7 ~lethod T.C. F.C. ~2 N2 ~l2 Co (~) ~ (%) (%~ ~%) B.3~ o.n8 0.030.25 0.01 0.5 ~0 I~ 8.3~ n.ns o~l 0.05 0.001 0.5 III) 8~45 n.n~ n.on~ o.ol 0.02 0.5 I~') 8~5~ 0~0 n.nOl 0~001 0.001 0.5 The carbides prepared by heating in gaseous atmospheres were all of a monocarbide, while the carbide obtained by heating in vacuum contained free carbon in a large amount and Mo2C precipitated according to the results of X-ray analysis.
When a (Mo9 W9 Cr)C-16 % (Co + Ni) alloy was then prepared using Carbide II) of the above described carbides, there was found an M2C phase dispersed evenly in the alloy in a proportion o 10 ~ by volume.
The properties of our alloy and the prior art WC-Co alloy were compared using an end mill under the following conditions:
Workpiece: SCM 3, HRC 8 13, Length 385 mm, End Mill: 8 mm ~, Two Cutting Edges, Right Cutting Edge Right Twist 25 , Solid Cutting System: Cutting of Groove in 5 mm Depth on Above ~orkpiece, Comparison of Lifes by Measurement of Time or Cutting Length until VB = 0.3 mm or until chipping Machine: NoO 4 Plain Milling Machine Cutting Conditions: V = 26.5 m/min f = 0.0285 mm/edge Water-insoluble Cutting Oil The results are shown in Table 8:
Table 8 Cutting Length Number of Grooves Cut Alloy of Present Invention 36.3 m 92.6 Fine Particle WC Alloy of 24 5 62.5 Prior Art Cemented Carbide Alloy K10 1 54 4 o Prior Art High Speed Steel 2.55 67 -~6-'.1~ 5~

As evident from these results, the alloy of the present invention is superior to other known alloys concerning the wear resistance and chipping resistance, since in our alloy, the high temperature hardness is high and, thus, the toughness can be raised wi~hout lowering the wear resistance even if the quantity of the binder phase is increased.
Example 10 The solid solution carbide (Mo, W, Cr)C obtained by the procedure set forth in Example 9 was mixed with 10 % of Ni powder and ball milled sufficiently for 100 hours by wet process in an organic solvent. The thus mixed powder was compacted under a pressure of 1 tons/cm2 and alloyed at a temperature of 1400 C. In the resulting alloy, there was found a (Mo, W~2C phase with 1 micron or less dispersed evenly in a proportion of 5 % by volume.
The resulting alloy was subjected to rubbing using a diamond paste to give a specular surface. The physical properties of this alloy are shown in Table 9, from which it is apparent that the alloy of the present invention is superior to the prior art WC-Co type alloys in density, toughness and corrosion resistance to sweat:
Table 9 Density Hardness T.R.S. Corrosion _ _ _ _ (Kg/mm2) Resistance *
Alloy of Present 10 8 91.0 180 Good Invention WC-Co Alloy of 14 5 90.S 160 Not Good Note: Corrosion Resistance was measured by immers-ing in an artificial sweat for 48 hours.
When a watch frame was made of this alloy and subjected to a performance test, the alloy of the present invention s~

was more useful because of its light weight and excellent resistance to scratching and to sweat.
Example 11 80 % of a carbide o:E molybdenum and tungsten with a molar ratio of 7 : 3 9 (MoO 7Wo 3)C was mixed with 10 %
of Co and 10 % of Ni, after which the quantity of carbon was controlled so that the carbon content in the alloy be 98 at % based on the carbon content 7.11 % in the stoichio-metric composition of the alloy.
For a comparative test, Sample Nos. (I) to ~III) were prepared as shown in the following:
Sample No. (I): 0.1 % of Fe was added as an impurity to the above described alloy so as to disperse (Mo, W)~S
evenly according to the present invention.
Sample No. ~ No impurity was added to the above described alloy as in the prior art.
Sample No. (III): Prior art WC-18 vol~Co alloy for impact tools.
In the present invention, (Mo, W)2C was uniformly dispersed in a gTanular form with proportion of about 8 %
by volume as shown in Fig. 2, whilst in the prior art (Mo, W)C-Co alloy, needle crystals were precipitated as shown in Fig. 1.
Headers were made of these alloys and the life tests thereof were carried out by subjecting to plastic working of the head of a screw consisting of SCr 4 stee~
rod, thus obtaining results as shown in Table 10:

Table 10 (Number of Screws/Die~
0 2 4 6 8 10 x 1~2 Sample l~o. (I) --~
Sample No. (II) x Sample No. (III) - x Note~ Mark x means a broken point.
As can be seen from the t0st Tesultg the alloy of the present invention shows the maximum life and a sufficient performance e~en if cracks or deformations occu~.
Example 12 A (MoO 7Wo 3~Cz powder where z = 0.9 was synthe-sized, mixed with 15 % of Ni powder and 2 % of Mn powder, ball milled adequately by wet process, compacted and sintered at 1350 C ;n vacuum. Examination of the texture showed the pTesence of a granula~ (Mo, W)2C with a size of 2 microns dispersed in a proportion of about 10 ~ by volume.
The properties of the so obtained alloy are as follows:
Density: 9.~8 g/cm3 Hardness: HRA ~ 88-5~ Hc = 0-4~a = o TTansve~se Rupture Strength: 170 Kg/mm2 This cemented carbide alloy i.s non-magnetic.
Example 13 The quantity of carbon is an alloy consisting of 85 % of (MoO 7Wo 3)C, 16 % of Ni, 0.6 % of Mn and 3 % of Re was controlled so that the alloyed carbon content be 95 at ~ based on the theoretical carbon content ~7.59 %) and 0.1 % of Fe was added to the alloy. The mixture was sintered at 1450 C for 1 hour in vacuum, thus obtaining an alloy having the following properties:
Density: 10.25 g/cm3 Hardness: HRA 89.4 Transverse Rupture Strength: 165 Kg/mm2 4~ = 0 Hc =
About 1 % by volume of an M2C phase was found in the structure of the present alloy. The corrosion resistance of the alloy of the present invention and that of the prior art WC-7 % Co alloy are tabulated below:
Table 11 (Unit: mg/cm2/hr~
Hot 10 % H~SO4 Hot 35 % HNO3 Alkali Solution Solution .
WC-7 ~ Co Alloy 15 9 Our Alloy 0.4 4 0 Example 14 Using (Mo0 7W0 3)C prepared by adding 0.~ % of Fe during production of the carbide~ there were obtained for trial Composition (A) (Mo0 7Wo 3~C-15 ~ Co, Composition ~B) (Mo0 7~ 3)C-7.5 % Co-7.5 % Ni and Composition (C) (Mo0 7Wo 3)C-15 % Ni. During the same time, starting materials were taken by weighing so that the alloyed carbon content z in (Mo0 7Wo 3)Cz be 98 at %, ball milled by wet process in an organic solvent, dried, compacted and then sintered at 1350 C in vacuum to thus obtain : alloys having the properties shown in Table 12.
For comparison, the alloy properties of the prior art compositions, i.e. Composition (D) (Mo, W)C-15 % Co and Composition (E) (Mo, W)C-7.5 % Co-7.5 % Ni are also shown in Table 12. The sum of oxygen and nitrogen in our alloys (A), (B) and (C) was 0.15.

Table 12 T.R.S. Hardness l'.C. F.C. Value z Amount of Fe (~/non2) (HRA) _ ~a$ %) in Binder Phase Our Invention Composition ~A) 240 87.0 7.43Q.00 98 1.1 Compositioll tR) 290 86.9 Co~l~osition ~C) 230 86.5 " 1l .. ..

~ior Art Composition (D~ 175 86.8 7.5~ " 100 Compositioll ~E) 140 86.6 7.43 " 98 .~s is evident from this table, ~he alloys of the present in~ention tA), tB) and (C) have a higher toughness than the prior ar~ alloys tD~ and IE).
According to the present inYention, the alloy is stabili-ed and, con~equently, the alloy strength is not lo~ered b~ the presence of Fe, N and O in the alloy e~en i~ tlle ~llo~ed c~lrhon content is less than the theo~etical carbon content. In the alloy of the present invention, a ~r~nul~r ~lo, I~ C ~s evenly precipitated, while in the prior nrt ~llo~ lo, IY~C due to lack of carbon ~as pre-~n ci~it~lted ~ needle crystals, resulting in lowering of the tou~hness.
: E~mple 15 ~ al~till~ m~ltçrials were mixed so that the hardh;~ of ~ desir~ lloy be composed of ~MoO 7W~ 3)C of

5 m;erolls ;Ind ~lo(~ 3)~C of 3 microns and the binder rh;l~e b~ co~ o~l of Co ~nd Ni with 0.1 to 1.0 % of Fe ~nd ~n ;~ re~lred therefrom as shown in Exa~ple ~ the ~tt~l~otul~e of the alloy obtained by this proce-d~re~ ~ o;lrl~id~ o~ C type was uniformly dispersed. The rclertie~ of o~lr ~lloy ~Ire shown in Table 13:

5~3 Table 13 tMoO 7W0 3)C tMoO 7W0 ~)2L Co Ni ~nount of T.R.S.
(~) (%) ~ ~%) ~%) Fe i B nder (Kg/mm ) ( % ) 1 85 5 10 - 0.1 Z60 2 " " " - 1.0 275 3 " " 6.6 3.4 0.5 3~0 ~ " " 7.5 2.5 0.5 290 In the alloy of the present invention, the toughness can be increased by adding Fe as an additive element to the binder phase consisting of Co and Ni to disperse a carbide of M2C type well in the alloy.
Example 16 (I) WC powder with a particle size of 6 microns was mixed with Mo2C powder with a particle size of 2 microns and carbon to give a final composition o~ carbide ~MoO 5W~ 5)C, ball milled by wet process for 30 hours and the mixed powder was reacted at 2000 C for 1 hour in an H2 stream. The carbide had a carbon content of T.C. 7.81 ~ and F.C. 0.03 %
; and the combined carbon content was near the theoretical value of ~MoO 5Wo 5~C. X-ray diffraction showed that the peak o Mo2C disappeared and there were the peaks of (Mo, W)C only. However, examination of the cross section of the powder taught that the powder was of a core struc-ture. According to the examination of the peaks at the high angle side, there were two phases separated.
tII) WC powder with a particle size of 1 micron was mixed with Mo2C powder with a particle size of 2 microns and carbon to give a final composition of carbide (MoO 5Wo 5)C, to which 0.5 ~ of Co was then added as a diffusion aid. The mixed powder was heated at 2000 C
for l hour in an H2 stream, the temperature being lowered ~ ~ ~ S~3~ ~

to 1400 C, and the product was held at the same tempera-ture for 10 hours. Analysis of the product showed T.C.
7.71 %, F.C. 0.05 % and Co 0.S %, the carbon content being near the theoretical carbon content. X-ray diffraction showed no separation into two phases 9 but showed a completely single phase.
Alloys were prepared from the carbides prepared by the above described procedures (I) and ~ . Starting materials were taken by weighing to give a composition of ~Mo, W)C-15 % Co, ball milled by wet process in an organic solvent, dried~ compacted and sintered at 1400 C in vacuum, thus obtaining alloys having the following pro-perties:
Table 14 Density Hardness Transverse Rupt~re (HRA) Strength ~Kg/mm Our Inven~ion ~I) 11.6 88.6 280 Prior Art (II)11.5 88.0 180 In he alloy of the present invention, there are two kinds of MC type phases and a granular M2~ phase dispersed evenly, whereby a high toughness can be given with holding the properties of the prior art WC-Co type cemented car-bide alloy. On the contrary, the prior art alloy is a uniform solid solution, but lacks in toughness.

Claims (15)

The embodiments of the invention in which an exclusive property or privilege is claimed are defined as follows:

1. A hard alloy comprising a hard phase consisting of at least one compound having a crystal structure of simple hexagonal MC type (M: metal; C: carbon) selected from the group consisting of mixed carbides, carbonitrides and carboxynitrides of molybdenum and tungsten as a predominant component, and a binder phase consisting of iron, cobalt, nickel and chromium, in which a hard phase consisting of a compound of M2C type, wherein M and C are defined as above, having a crystal structure of the hexagonal type is uniformly dispersed in a proportion of at most 30% by volume based on all the hard phases, said compound of M2C type being in a granular or globular form with a size of at most 10 microns, wherein the carbon content in the hard phases of the alloy is in an atomic proportion of 0.98 to 0.8 with respect to the theoretical carbon content of the MC type compound.

2. The hard alloy as claimed in claim 1, wherein a part of the compound of MC type is replaced by a B1 type hard compound.

3. The hard alloy as claimed in claim 2, wherein the B1 type hard compound contains at least one of titanium, zirconium, hafnium, vanadium, niobium, tantalum, chromium, molybdenum and tungsten.

4. The hard alloy as claimed in claim 2, wherein the quantity of the B1 type hard compound replaced is at most 30%
by weight.

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

6. The hard alloy as claimed in claim 5, wherein the quantity of Cr is 0.3 to 10% by weight,

7. The hard alloy as claimed in claim 1, wherein a Part of the carbon in the carbides forming the hard phases is replaced by at least one of nitrogen and oxygen,

8. The hard alloy as claimed in claim 7, wherein the quantities of nitrogen and oxygen are defined, in connection with the alloy composition, by the relations of:
;
.

9. The hard alloy as claimed in claim 1, wherein the binder phase is incorporated in a proportion of 3 to 50% by weight of the alloy composition.

10. The hard alloy as claimed in claim 1, wherein the quantity of iron in the alloy composition is defined by the relation of:

11. The hard alloy as claimed in claim 1 wherein the dispersion of the hard phase consisting of a compound of M2C type is carried out by adding an impurity element to the binder phase.

12. The hard alloy as claimed in claim 11, wherein the impurity element is at least one of beryllium, magnesium, calcium, boron, silicon, phosphorus, manganese, iron and rhenium.

13. The hard alloy as claimed in claim 11, wherein the impurity element is added in a proportion of 0 to 3% by weight.

14. The hard alloy as claimed in claim 1, wherein at least one of manganese, rhenium, copper, silver zinc and gold is incorporated in the binder phase in an amount to make the alloy non-magnetic.

15. The hard alloy as claimed in claim 1, wherein the hard phase consisting of a compound of MC type comprises two or more simple hexagonal phases differing in a ratio of Mo/W.