JPH0629475B2 - cemented carbide - Google Patents

Description

The present invention relates to the field of cemented carbide, particularly in the field of tools, having cobalt as a binder material. More particularly, the present invention relates to the formation of hard, heat-resistant oxides, nitrides, etc. on the surface.
The present invention relates to a cemented carbide metal cutting insert having a boride or carbide coating.

In the past, various hard heat-resistant coatings have been applied to the surface of cemented carbide cutting inserts to improve the wear resistance of the cutting edges and increase their cutting life. For example, see U.S. Patent Nos. 4,035,541 (assigned to the present applicant), 3,564,683, 3,616,506, and 5,616,506.
No. 3,882,581, No. 3,914,473, No. 3,736,107, No. 3,96
See, for example, US Pat. Nos. 7,035, 3,955,038, 3,836,392, and Reissue No. 29,420. However, these heat-resistant coatings can reduce the toughness of the cemented carbide inserts to varying degrees. The extent of this degradation depends, at least in part, on the structure and composition of the coating and the method used to apply it. Thus,
Although heat resistant coatings have improved the wear resistance of metal cutting inserts, they have not reduced the susceptibility of the cutting edge to failure by chipping or breakage, especially in intermittent cutting operations.

Previous efforts to improve toughness or cutting edge strength in coated cutting inserts have centered around the creation of a cobalt-enriched layer extending inward from the substrate/coating interface. It has been discovered that cobalt enrichment of the surface layer of certain C-type porosity substrates can be achieved during a vacuum sintering cycle. These cobalt-enriched zones were characterized by A-type porosity, with the majority of the substrate having C-type porosity. The cobalt-enriched zones were usually depleted of solid solution carbides to various depths and degrees. Cobalt enrichment is desirable because:
It is well known that increasing the cobalt content increases the toughness or impact resistance of cemented carbides. However, the level of enrichment produced is difficult to control in type C porosity substrates. Typically, a cobalt and carbon coating is formed on the surface of the substrate. This cobalt and carbon coating is removed before coating the refractory material on the substrate to allow for an intimate bond between the coating and the substrate. Sometimes, the level of cobalt enrichment in the subsurface layer of the substrate is so high that it adversely affects the wear of the flank surface. As a result, sometimes the cobalt-enriched layer on the flank surface of the substrate is ground away, leaving the cobalt-enriched layer only on the rake surface and the possibility of type C porosity material on the flank surface.
Compared to substrates with type B porosity, substrates with type C porosity are chemically less homogeneous, resulting in less controlled formation of ake (η) phase (a hard, brittle phase that affects toughness) at the coating/substrate interface, reducing coating adhesion and increasing non-uniform coating growth.

By definition, the porosity observed in cemented carbides is determined by ASTM (American
ican Society For Testing Materials) No. ANSI/AST
They fall into one of the three categories recommended by B-276-79:

Type A for pore sizes smaller than 10 microns in diameter.

B for pore sizes between 10 and 40 microns in diameter
Type.

Type C for irregular pores caused by the presence of carbon inclusions. These inclusions are pulled out of the sample during metallographic preparation, leaving behind the irregular pores.

In addition to the above classifications, the observed porosity can be assigned a number from 1 to 6 to indicate the degree or frequency of the observed porosity. The method for making these classifications is described in "The porosity classification" by Dr. D. Schwarzkopf, published by MacMillan, New York (1960).
Kopf and Dr. R. Kieffer, Cemented Carbides, pp. 116-120.

Cemented carbides can also be classified according to the carbon and tungsten content in their binder.

Tungsten carbide-cobalt alloys with excess carbon are C
Tungsten carbide-cobalt alloys are characterized by type C porosity, which, as noted above, is effectively free carbon inclusions. Tungsten carbide-cobalt alloys with low carbon content and where the cobalt is saturated with tungsten are characterized by the presence of eta phase, _M12C , or _M6C carbides (M represents cobalt and tungsten). Between the extremes of type C porosity and eta phase, there is a range of intermediate binder alloy compositions containing various levels of dissolved tungsten and carbon, yet lacking the presence of free carbon or eta phase. The tungsten level present in the tungsten carbide-cobalt alloy can also be characterized by the magnetic saturation of the binder alloy.
That is, the magnetic saturation of a cobalt alloy is a function of its composition. Carbon-saturated cobalt has a magnetic saturation of 158 ga
It has been reported to have a magnetic saturation of 100 gauss-cm ³ /gr cobalt, while a magnetic saturation of less than 126 gauss-cm ³ /gr cobalt indicates the presence of the eta phase.

It is an object of the present invention to provide a cemented carbide tool having a binder-thickened layer near its surface, with substantially all of the porosity throughout the tool being Type A or Type B, and further to provide a tool body having a carbon level ranging from the eta layer and a binder-thickened layer near its circumferential surface. It is also an object of the present invention to combine the above cemented carbide tool with a heat-resistant coating to provide a coated cutting insert having a combination of high wear resistance and high toughness.

Other and further objects of the present invention will become more fully apparent upon consideration of the following description of the invention.

In accordance with the present invention, it has been found that a binder-thickened layer can be formed near the peripheral surface of a cemented carbide tool by:

Milling and blending a first carbide powder, a binder powder, and a chemical powder selected from the group consisting of metals, alloys, hydrides, nitrides, and carbonitrides of transition elements, the carbides of which have a more negative free energy of formation than the first carbide near the binder melting point; then sintering or subsequently heat treating a compact of the blended materials to at least partially convert the chemical to its carbide.

In accordance with the present invention, a cemented carbide tool is provided with a binder-enriched layer near the periphery of a cemented carbide tool having substantially only type A or B porosity throughout the body. Enrichment can also be achieved in cemented carbide tools having carbon levels ranging from eta-phase to type C porosity.

It has also been found that the cemented carbide tool according to the invention has a partially binder-depleted layer beneath the binder-rich layer.

The first carbide is preferably 70-90.5% by weight tungsten carbide. The binder alloy is preferably 5-10% by weight cobalt, nickel, iron or alloys thereof, most preferably cobalt.

The chemical agent is selected from hydrides, nitrides and carbonitrides of Group IVA and VA elements and is preferably added in a small effective amount, most preferably 0.5 to 2 weight percent of the powder charge. The most preferred chemical agent is titanium nitride or titanium carbonitride.

The cemented carbide tool of the present invention has also been found to have a first layer near its peripheral surface that is at least partially depleted of carbide solid solution.

The cemented carbide tool of the present invention has also been found to have a second layer below the solid solution depleted layer that is enriched in carbide solid solution.

The cemented carbide tool of the present invention preferably has a rake face and a flank face, a hard, high-density, heat-resistant coating intimately bonded to the faces and having a cutting edge at the junction of the faces, and the bond-thickened layer may be ground off the flank faces prior to application of the coating.

The heat resistant coating is preferably a metal oxide, carbide,
It consists of one or more layers of nitride, boride or carbonitride.

The nature of the present invention will become more apparent from the following detailed description taken in conjunction with the accompanying drawings.

The foregoing objects of the present invention are achieved by heat treating a cemented carbide former containing an element whose carbide has a more negative free energy of formation than tungsten carbide at temperatures approaching or above the melting point of the binder.

For cutting insert applications, the element or chemical agent may be selected from the group IVA and VA transition metals, their alloys, nitrides, carbonitrides, and hydrides. It has been found that the layer of material adjacent to the periphery of a tungsten cemented carbide tool can be consistently binder-enriched and usually at least partially depleted of carbide solid solution during sintering or reheating at temperatures above the melting point of the binder alloy by adding nitrides, hydrides, and/or carbonitrides of the group IVA and VA transition elements to the powder charge.

During sintering, these IVA and VA group additives react with carbon to form carbides or carbonitrides which exist partially or totally in solid solution with tungsten carbide and other carbides present.

The level of nitrogen present in the final sintered carbide is generally reduced from that added as nitrides or carbonitrides because these additions are unstable at high temperatures above and below the melting point of the binder alloy.
If the sintering atmosphere contains nitrogen at a concentration lower than its equilibrium vapor pressure, this will cause at least partial volatilization of nitrogen from the sample. If the chemical agent is added as a metal, alloy, or hydride, it will also be converted to a cubic carbide, typically in solid solution with tungsten carbide and other carbides present. The hydrogen in the added hydride is volatilized during sintering.

Tantalum, titanium, niobium, hafnium metals, hydrides, and carbonitrides can be used alone or in combination to promote cobalt enrichment through sintering or subsequent heat treatment of tungsten carbide-cobalt based alloys having a wide range of carbon, collectively approximately 15% by weight.
Additions up to 1000 ppm have been found to be useful. Zirconium and vanadium metals, nitrides, carbonitrides, and hydrides are also believed to be suitable for this purpose. In alloys with A and B porosity and low carbon alloys containing eta phase, cobalt enrichment occurs without capping of the surrounding cobalt or carbon, thus eliminating the need to remove excess cobalt and carbon from the hardmetal surface before applying a refractory coating.

Additions of approximately 0.5 to 2 weight percent titanium, particularly in the form of titanium nitride or titanium carbonitride, are preferred for tungsten carbide-cobalt based alloys. Because titanium nitride is not completely stable during vacuum sintering, causing at least partial volatilization of nitrogen, it is preferred to add half a mole of carbon per mole of starting nitrogen to maintain the carbon level required for tungsten-poor cobalt binder alloys. Cobalt enrichment by heat treatment of tungsten carbide-cobalt based alloys has been found to occur more readily when the alloy contains a tungsten-poor cobalt binder.

The tungsten-poor cobalt binder preferably has a cobalt magnetic saturation of 145 to 157 gauss-cm ³ /gr.

The addition of titanium nitride to a tungsten carbide-cobalt based powder mixture, along with the necessary carbon, facilitates the formation of a magnetically saturated cobalt binder alloy of 145-157 gauss-cm ³ /gr, which is normally difficult to achieve.
Although cobalt binder alloys having 125 gauss-cm ³ /gr cobalt magnetic saturation are preferred, alloys containing tungsten saturated cobalt binder alloys (up to 125 gauss-cm ³ /gr cobalt) can also be enriched.

It has been found that a cobalt-enriched layer thicker than 6 microns results in a significant improvement in cutting edge strength of the refractory coated cemented carbide cutting insert.

Deep cobalt enrichment was achieved at just over 125 microns,
A cobalt-enriched layer having a thickness of 12 to 50 microns is suitable for the coated cutting insert. The cobalt content of the cobalt-enriched layer in the heat-resistant coated cutting insert is 150 to 300% of the average cobalt content as measured on the surface by energy dispersive X-ray analysis.
It is also preferred that

Binder thickening is the process of sintering all tungsten carbide-binder-cubic carbides (i.e., tantalum, niobium, titanium, vanadium,
This is believed to occur in alloys containing 3% or more by weight of binder. These alloys should be thickened using the method described above.

However, for cutting insert applications, it is preferred that the binder content be 5-10% by weight cobalt and 5-20% by weight total cubic carbide content.
Cobalt is the preferred binder, but nickel, iron and alloys thereof with each other and with cobalt may be substituted for cobalt. Other binder alloys containing nickel or cobalt or iron may also be suitable.

The sintering and heat treating temperatures used to achieve binder densification are typical liquid phase sintering temperatures. For cobalt-based alloys, these temperatures are between 1285°C and 1540°C. The sintering cycle is performed for at least 15 cycles at a given temperature.
The results can be further optimized by using controlled cooling rates from the heat treatment temperature to a temperature below the melting point of the binder alloy. These cooling rates should be 25-85°C/h, preferably 40-70°C/h.
/h. Most preferably, the heat treatment cycle for cutting insert substrates with cobalt binder is 1370
~1500 °C for 30-150 min, followed by 40-70 °C/
The temperature is then cooled to 1200°C at a cooling rate of 10 -3 torr. The pressure level during the heat treatment can vary from 10 ^-3 torr to the high pressures typically used in hot isostatic pressing. The preferred pressure level is
0.1 to 0.15 torr. If nitride or carbonitride additions are utilized, they are preferably below the equilibrium vapor pressure of nitrogen in the sintering atmosphere, thereby allowing volatilization of nitrogen from the substrate.

Although initial densification occurs during sintering, subsequent grinding steps in the metal cutting insert manufacturing process may remove the densified zone, in which case subsequent heat treatment according to the above parameters may be used to develop a new densified zone below the circumferential surface.

The bond-thickened substrate to be used in the coated cutting insert can have a bond-thickened layer on both the rake and flank faces. However, depending on the type of insert, the bond-thickened layer on the flank faces may be removed in some cases, although this is not necessary in all cases to obtain optimum performance.

The bond-enriched substrate can be coated using heat-resistant coating techniques well known to those skilled in the art. While the applied heat-resistant coating can have one or more layers of materials selected from carbides, nitrides, borides, and carbonitrides of elements of Groups IVA and VA, and oxides and oxynitrides of aluminum, it has been found that a good combination of cutting edge strength and flank wear resistance can be achieved by combining a substrate having a bond-enriched layer according to the present invention with a coating consisting of an aluminum oxide coating on an inner layer of titanium carbide, or an inner layer of titanium carbide bonded to an intermediate layer of titanium carbonitride bonded to an outer layer of titanium nitride, or a coating consisting of titanium nitride bonded to an inner layer of titanium carbide. Most preferred are cemented carbide tools having a bond-enriched layer according to the present invention in combination with a titanium carbide/aluminum oxide coating. In this case, the coating should have a total thickness of 5 to 8 microns.

If the metal binder content is less than 5% by weight, it will not be tough enough to give the best properties, and if the content is greater than 10% by weight, it will be too soft.

Metal carbide contents below 5 weight percent do not provide good resistance to surface cratering or thermal stability. Metal carbide contents above 20 weight percent are very expensive. Also, in compositions containing titanium and tantalum, titanium provides resistance to deformation, while tantalum provides toughness. For this reason, the tantalum content is preferably about twice that of titanium. The remainder of the composition is tungsten carbide.

1, there is shown diagrammatically an embodiment of a coated metal-cutting insert 2 according to the present invention. The insert 2 comprises a substrate or cemented carbide body 12, which is coated with a second layer, a bond-enriched layer 14.
14 and the majority of the interior, excluding the surface layer, of substrate 12 having a chemical composition substantially equal to that of the original powder blend.
A first layer, a binder-reduced layer 16, on the bulk 18.
and

A binder-thickened layer 14 is present on the rake face 4 of the cemented carbide body 12 and is ground away from the flank face 6. Located inward of the binder-thickened layer 14 is a binder-reduced zone 16. This binder-reduced zone 16 has been found to develop along with the binder-thickened layer when the cemented carbide body is fabricated in accordance with the present invention.

The binder-reduced zone 16 is partially depleted of binder while being enriched with solid solution carbides.

The binder-enriched layer 14 is partially depleted of carbide solid solution. Inward of the binder-depleted zone 16 is the bulk 18 of the substrate material.
There is.

A cutting edge 8 is formed at the junction of the rake surface 4 and the flank surface 6. While the illustrated cutting edge 8 is honed, honing of the cutting edge 8 is not necessary for all applications of the present invention. As can be seen in FIG. 1, a binder-thickened layer 14 extends into this cutting edge area and preferably adjacent most, if not all, of the honed cutting edge 8. A binder-reduced zone 16 extends to the flank surface 6 immediately below the cutting edge 8 and is preferably adjacent to the substrate.
It is softer than 12, which is advantageous for receiving coating and reducing the possibility of cracks. Heat-resistant coating 10
is intimately bonded to the circumferential surface of the cemented carbide body 12 .

Thus, the difference between the technology of the present invention and the prior art is that the prior art does not include a starting composition as in the present invention and does not sinter a composition to achieve a binder-thickened layer beginning at the peripheral surface of the substrate and extending inward therefrom, whereas the technology of the present invention has a binder-thickened layer beginning at the peripheral surface of the substrate and extending inward therefrom.

These and other features of the present invention will become more apparent from a consideration of the following examples.

Example 1 A mixture containing 7000 grams of powder was mixed with paraffin, surfactant, solvent, and cobalt-bonded tungsten carbide cycloids in the amounts and proportions below.
The mixture was milled and blended for 16 hours.

Cuboid insert preforms with dimensions of 15.1 mm x 15.1 mm x 5.8-6.1 mm and a weight of 11.6 gr were pill pressed using a force of 8200 kg.
It was vacuum sintered at 1496°C for 30 minutes and then cooled under ambient furnace conditions. After sintering, the insert had a weight of 11.25g and dimensions of 13.26mm x 13.26mm x 4.95mm.

These inserts were then machined to the SNG433 grinding dimensions (this identification number is an American Standard) as follows:
(The insert is based on an insert identification system developed by the ds Association and commonly adopted in the cutting tool industry. The international designation is SNGN12 04 12.) 1. The top and bottom (rake faces) of the insert were ground to a thickness of 4.75 mm.

2. Inserts were heated at 1427°C for 6 hours under a vacuum of 100 microns.
The alloy was heat treated for 10 minutes, then cooled to 1204°C at a cooling rate of 56 C/h, and then cooled under ambient furnace conditions.

3. The periphery (flank) was ground to produce a 12.7 mm square and the cutting edge was honed to a radius of curvature of 0.064 mm.

A titanium carbide/titanium carbonitride/titanium nitride coating was then applied to the ground insert using the following chemical vapor deposition (CVD) techniques in the following order of application:

Inserts without TiN and its associated carbon additives were also made from the same powder formulation as the above inserts. Microstructural data obtained from the coated inserts are as follows:

Example 2 Utilizing the formulation of Example 1, pill pressed green inserts were fabricated with and without TiN and associated carbon additives according to Example 1. These inserts were
The inserts were sintered at 1496°C for 30 minutes under a 25 micron vacuum and then cooled under ambient furnace conditions. The inserts were then honed (0.064 mm radius) and then coated with a TiC/TiCN/TiN CVD coating according to the technique shown in Table 1. It should be noted that in this example, a cobalt-enriched layer was present on both the flank and rake faces.

These coated inserts have been evaluated and
The following results were obtained:

Example 3 A mixture of the following materials was charged to a cylindrical mill along with a surfactant, a fuming binder, a solvent and 114 kg of cycloids.

The powder charge was balanced to produce 6.25% total carbon by weight in the charge. The mixture was blended and milled through 90.261 revolutions to obtain an average particle size of 0.90 microns. The mixture was then wet screened, dried, and ground in a hammer mill. Compacts were pressed and sintered at 1454°C for 30 minutes, then cooled under ambient furnace conditions.

This process produced a sintered preform containing cobalt with an overall magnetic saturation of 117-121 gauss- ^cm3 /gr (i.e., the measurement included the bulk material and binder-thickened material). Evaluation of the microstructure of the sintered preform revealed the presence of eta phase throughout the preform, and a porosity of A-2.
to B-3, and the cobalt enriched zone thickness is approximately
26.9 microns and a solid solution depleted zone thickness of approximately 31.4 microns.

Example 4: The following materials were added to a tungsten carbide-cobalt alloy-lined mill jar, measuring 190 mm inside diameter and 194 mm long. Additionally, 17.3 kg of 3.2 mm tungsten carbide-cobalt cycloids were added to the jar. These materials were milled and blended by rotating the mill jar around its cylindrical axis at 85 revolutions per minute for 72 hours (i.e., 367,200 revolutions).

This mixture was balanced to produce a binder alloy of 2 w/o W and 98 w/o Co. After milling and blending, the slurry was wet screened to remove oversized particles and contaminants, dried at 93°C under nitrogen, and then hammer milled in a Fritz Patrick J-2 Fitzmill to break down clumps.

This powder is used to press compacts, which are then
It was sintered at 1454°C for 30 minutes and then cooled under ambient conditions.

The top and bottom (i.e., rake) surfaces of the inserts were then ground to final thickness. This was followed by heat treatment at 1427°C under 100 microns of vacuum.
After holding for 10 minutes, the inserts were cooled at a rate of 56°C/hour to 1204°C, then furnace cooled under ambient conditions. The peripheral (or flank) surfaces were then ground to a 12.70 mm square and the insert cutting edges were honed to a 0.064 mm radius. These treatments resulted in an insert substrate with a cobalt-enriched zone and a solid solution-depleted zone on the rake surface only, which were ground away from the flank surfaces.

The inserts are then loaded into a coating reactor,
The inserts were coated with a thin layer of titanium carbide using the following chemical vapor deposition technique: The hot zone containing the inserts was first heated from room temperature to 900°C. During this heating period, hydrogen gas was allowed to flow through the reactor at a rate of 11.55 liters per minute. The pressure within the reactor was maintained at just under 1 atmosphere. The hot zone was then heated from 900°C to 9
The reactor was heated to 82°C. During this second heating stage, the pressure inside the reactor was maintained at 180 torr, and a mixture of titanium tetrachloride and hydrogen and pure hydrogen gas were admitted to the reactor at rates of 15 liters per minute and 33 liters per minute, respectively. The mixture of titanium tetrachloride and hydrogen gas was obtained by passing hydrogen gas through an evaporator that maintained the titanium tetrachloride at a temperature of 47°C. Once the temperature reached 982°C, methane gas was again admitted at a rate of 2.5 liters per minute.
The titanium tetrachloride was allowed to enter the reactor at a flow rate of 1.53 liters per minute under a pressure slightly less than 1 atmosphere while argon was passed through the reactor at a flow rate of 1.53 liters per minute under a pressure slightly less than 1 atmosphere. The pressure in the reactor was reduced to 140 torr. Under these conditions, titanium tetrachloride reacted with methane in the presence of hydrogen to form titanium carbide on the surface of the heated insert. These conditions were maintained for 75 minutes, after which the flows of titanium tetrachloride, hydrogen, and methane were terminated. The reactor was then allowed to cool while argon was passed through the reactor at a flow rate of 1.53 liters per minute under a pressure slightly less than 1 atmosphere.

Examination of the microstructure in the final insert revealed a cobalt-enriched zone extending inwardly to 22.9 microns and a cubic carbide solid solution-depleted zone extending inwardly from the substrate rake face to 19.7 microns. The porosity in the cobalt-enriched zone and the remainder of the substrate was determined to be between A-1 and A-2.

Example 5 The materials in this example were mixed with the following material charges:
The blend was milled using a stage milling process.

This was balanced to produce a binder alloy of 2 w/o W and 98 w/o Co.

Test inserts were then fabricated and coated with TiC according to and with the test blanks described in Example 4.
It was coated.

Microstructural evaluation of the coated inserts showed that the cobalt-enriched zone and the porosity in the bulk of the material were A-1. The cobalt-enriched zone and the solid solution-depleted zone were approximately 100 mm from the rake face, respectively.
It extended inward to depths of 32.1 microns and 36 microns.

Example 6 The following materials were loaded into a 190 mm inner diameter mill jar:

This mixture was balanced to produce a binder alloy of 2 w/o W and 98 w/o Co.

Additionally, cycloids were added to the mill. The mixture was then milled for four days. The mixture was dried in a Sigma blender at 121°C under partial vacuum and then passed through a 40 mesh screen in a Fitzmill.

SNG433 inserts were then fabricated using the techniques described in Example 4. However, the inserts in this example were CVD coated with a TiC/TiN coating. The coating procedure used was as follows:

1. TiC Coating: In the coating reactor, the sample was maintained at approximately 1026-1036°C under a 125 torr vacuum. Hydrogen carrier gas was introduced at a flow rate of 44.73 liters per minute.
The _iCl4 was flowed into the evaporator.

The evaporator was maintained at 33-35°C under vacuum. _TiCl4 vapor was entrained in a hydrogen carrier gas and transported into the coating reactor. Free hydrogen and free methane were released at 19.8°C, respectively.
The TiC coating was carried out at flow rates of 8 and 3.98 liters per minute into the coating reactor. These conditions were maintained for 100 minutes to produce a dense TiC coating that was tightly bonded to the substrate.

2. TiN Coating - The methane flow into the reactor was discontinued and nitrogen was allowed to flow into the reactor at a rate of 2.98 L/min. These conditions produced a dense TiN coating tightly bonded to the TiC coating, which was maintained for 30 minutes.

Evaluation of the coated inserts yielded the following results:

Example 7 A mixture of materials was made using the following two-stage milling cycle.

In step 1, 17.3 kg of the following material was poured into a 4.8 m 2 mill with an inner diameter of 181 mm and a length of 194 mm and a WC-Cc lining.
The mill was rotated around its cylindrical axis at 85 revolutions per minute (224,800 rpm) for 48 hours.

In stage II, 6314 g (90.2 wt%) WC and 1500
ml of Soltrol 130 was added, bringing the total charge to an additional 16
The mixture was rotated for 81,600 rpm.
The o, Co were balanced to produce a 95 w/o binder alloy. After grinding, the slurry was wet screened through a 400 mesh, dried at 93°C under nitrogen for 24 hours, and then passed through a 40 mesh screen in a Fitzmill.

The test specimen was 15.11mm x 15.11mm x 5.28mm with a total force of 16,400kg.
It was pressed uniaxially to dimensions of 8.6 gr/cc (specific gravity).

The unfired test specimens were baked at 1468°C under a 1 micron vacuum.
The inserts were sintered for 150 minutes. The inserts were then cooled under ambient furnace conditions. Graphite flakes were used as a release agent between the test inserts and the graphite sintering tray.

The sintered inserts were honed to a cutting edge radius of 0.064 mm. The inserts were then TiC/TiCN/TiN coated according to the following procedure.

1. The insert was placed in the reactor and air was evacuated from the reactor by flowing hydrogen through the reactor.

2. The insert was heated to approximately 1038° C. while maintaining a flow of hydrogen through the reactor. The pressure in the coating reactor was maintained at just above 1 atmosphere.

3. TiC coating: A mixture of _H2 and _TiCl4 is approximately 92%.
/min for 25 minutes and methane
A flow rate of 3.1 L/min was introduced for the same time. _{The TiCl4} vaporizer was maintained at approximately 6 psi and 30°C.

4. TiCN Coating--The flow of the _H2 + _TiCl4 mixture was maintained for 13 minutes, the flow of methane was reduced to half and nitrogen was introduced into the reactor at a flow rate of 7.13 liters/minute.

5. TiN Coating - For 12 minutes, the methane flow was discontinued and the nitrogen flow rate was doubled. After completion of TiN coating, the _H2 + _TiCl4 mixture and _N2 flows were discontinued, the reactor heating element was closed, and the reactor was heated to approximately 250°C.
The reactor was purged with free hydrogen until cooled to 250° C. At 250° C., the reactor was purged with nitrogen.

The insert substrates were determined to have A-1 to A-2 porosity in their non-enriched interior or bulk material. The cobalt-enriched and solid solution-depleted zones extended 25 microns and 23 microns from the surface, respectively. The non-enriched interiors had an average hardness of 91.7 Rockwell A. The coercivity, Hc, of the substrates was found to be 186 oersteds.

Example 8 A batch containing the following materials was ball milled.

30.0 w/o WC (1.97 micron average particle size) 750 kg 51.4 w/o WC (4.43 micron average particle size) 1286 kg 6.0 w/o Co 150 kg 5.0 w/o WC-TiC carbide solid solution 124.5 kg 6.1 w/o TaWC carbide solid solution 152 kg 1.5 w/o W 37.5 kg This mixture was charged to a total carbon content of 6.00 w/o. These materials were used to prepare 3409 kg of cycloids and
Crushed with 798 naphtha for 51.080 rpm. 0.8
A final particle size of 2 microns was produced.

5000gr of powder was divided from the blended milled batch and to this the following ingredients were added:

1.9 w/o TiN (pre-milled to approximately 1.4-1.7 microns)
96.9g 0.2w/o C (Ravin 410) 9.4g 1500 ml perchloroethylene These materials were then mixed with 50% by volume of cycloid (17.3kJ
g) for 16 hours in a 190 mm internal diameter, tungsten carbide lined mill jar. After milling was completed, the lot was wet screened through a 400 Lesch screen and sieved under partial vacuum in a Sigma Blender.
It was dried at 121°C and then passed through a 40 mesh sieve in a Fitz mill.

SNG433 semi-finished product was pressed using a force of 3600 kg to produce 8.24g
The semi-finished products were molded to a density of 10-25 microns/cc and a height of 5.84-6.10 mm. The semi-finished products were heated at 1454°C on a NbC powder separator under vacuum.
The sample was sintered at RT for 30 minutes and then allowed to cool in the furnace. The sintered sample had sintered dimensions of 4.93 mm x 13.31 mm square, a density of 13.4 gr/cc, and an overall magnetic saturation value of 146-150 gauss- ^cm3 /gr Co. Microstructural evaluation of the sample showed type A porosity throughout and a cobalt-enriched layer approximately 21 microns thick.

The top and bottom of the insert were then ground to an overall thickness of 4.75 mm. The insert was then ground 100 microns under vacuum.
Heat treated at 427°C for 60 minutes, then heated to 1204°C at a rate of 56°C/h.
The mixture was cooled to 100°C and then furnace cooled.

The flanks of each insert were ground to 12.70 mm square and the cutting edges were honed to a 0.064 mm radius of curvature.

The inserts were then titanium carbide/
It was CVD coated with aluminum oxide.

The insert is placed in the coating reactor and heated to approximately 1026°C.
The reactor was heated to 1030°C and held under a vacuum of 88-125 torr. Hydrogen gas was introduced at a flow rate of 44.73 liters per minute into the reactor under vacuum of 35-30 torr.
The TiCl4 vapor was passed through an evaporator containing _TiCl4 at 8°C. The _TiCl4 vapor was entrained in hydrogen and directed into the coating reactor. Simultaneously, hydrogen and methane were evaporated at 19.88 and 2.98°C.
The hydrogen and chlorine were then allowed to flow into the reactor at a rate of 19.8 kJ/min. These conditions of vacuum, temperature, and flow rate were maintained for 180 minutes to produce an adherent TiC coating on the insert. The hydrogen flow to the vaporizer and the methane flow into the reactor were then terminated. Hydrogen and chlorine were then allowed to flow into a generator containing aluminum particles at a temperature of 380-400°C and a pressure of 0.5 psi. The hydrogen and chlorine were each allowed to flow at a rate of 19.8 kJ/min.
The hydrogen and AlCl3 were flowing into the reactor at rates of 0.8/min and 0.8-1.0/min. The chlorine reacted with the aluminum to produce _AlCl3 vapor, which was then directed into the reactor. While the hydrogen and _AlCl3 were flowing into the reactor, _Co2 was also flowing at a rate of 0.5/min.
These flows were flowing into the reactor at a rate of 1
The temperature was maintained for 80 minutes, during which time the insert was held at 1026-1028° _C under a vacuum of approximately 88 torr. This procedure produced a dense coating of _Al2Cl3 tightly bonded to the TiC inner coating.

Evaluation of the coated inserts yielded the following results:

Example 9 A further 5000 gr of material was divided from the initial batch of material produced in Example 9. This material was mixed with 95.4 gr (1.9 w/o)
of pre-ground TiCN and 1.98 gr (0.02 wt%) of Ravi
n 410 carbon black was added as in Example 8,
It was mixed for 16 hours, sieved, dried and put through a Fitz mill.

The specimens were pill pressed and vacuum sintered at 1496°C for 30 minutes, then furnace cooled at the ambient furnace cooling rate.

Evaluation of the sintered samples yielded the following results:

EXAMPLE 10 A further 5000 gr of material was divided from the initial batch made in Example 9. Pre-ground TiCN in the amount of 95.4 gr (1.9 w/o) was added as in Example 9, mixed for 16 hours, sieved, dried and put through a Fitts mill. Test specimens were then pressed and, together with the test specimens of Example 9, 1496
It was sintered at ℃.

Evaluation of the sintered samples yielded the following results:

Example 11 The following mixture was charged using the following two-stage mix-mill cycle:

Stage I: The following materials were added to a 181 mm inner diameter, 194 mm long mill jar with a WC-Co lining along with 17.3 kg of 4.8 mm WC-Co cycloids. The mill jar was rotated around its cylindrical axis at 85 revolutions per minute for 48 hours (244.8
00 rpm).

Stage II The following materials are then added to the mill jar, followed by 16
Time rotated (81,600 rpm).

This mixture was balanced to produce a binder alloy of 10 w/o W and 90 w/o Ni. After discharging the mixture slurry from the mill jar, the slurry was wet sieved through a 400 mesh sieve (Tyler) and sieved under a nitrogen atmosphere.
It was dried at 3°C and then passed through a 40 mesh sieve in a Fitz mill.

The test specimens were pill pressed and sintered at 1450°C for 30 minutes under a 6.9 x ¹⁰ dynes/cm ^nitrogen atmosphere, followed by furnace cooling at the ambient furnace cooling rate. After sintering, the specimens were sintered at a 1 x 10 ^dynes /cm
The material was hot isostatically pressed at 1370°C for 60 minutes in a ³⁰⁰⁰ cc helium atmosphere. Optical metallographic evaluation of this sample showed that the material had A-3 porosity throughout and the solid solution depleted zone was approximately 25.8 microns thick.

The samples were then reconditioned and examined by energy dispersive X-ray line scan analysis (EDX) at various distances from the rake face. Figure 3 is a graph showing the change in relative concentrations of nickel, tungsten, titanium, and tantalum as a function of the distance from the rake face of the sample. As clearly shown in this figure, near the rake face there is a layer at least partially depleted of titanium and tantalum, which form carbides in solid solution with tungsten carbide.
This solid solution depleted zone extends inward approximately 70 microns. The discrepancy between this value and the values stated above is believed to be due to the fact that the sample was reconditioned between each evaluation, so that a different cross section of the sample was examined for each evaluation.

Corresponding to the reduction in titanium and tantalum is a nickel-enriched layer (see Figure 3). The nickel concentration in the enriched layer decreases as the distance from the rake face decreases from 30 to 10 microns, indicating that the nickel in the zone is partially evaporated during vacuum sintering.

The titanium and tantalum concentration peaks at 110 microns are believed to be due to the scanning of any large particles with high concentrations of these elements.

The two parallel horizontal lines represent a typical distribution of the bulk analysis values of the sample obtained for the mixture chemical composition.

Example 12 The following mixture was charged using the following two-stage mix-mill cycle:

Step I The following materials were milled according to Step I of Example 11:

Stage II: The following materials are then added to the mill jar and
It was rotated for six hours.

5980 gr (83.6 w/o) WC 140 gr Sunoco 3420 This mixture was balanced to produce a 10 w/o W - 90 w/o Ni binder alloy.

After discharging the mixture, it was sieved, dried and filtered according to Example 12.

The pressed test specimen was heated to 1466℃ under a 35μm vacuum atmosphere.
The sintered samples were sintered at 100°C for 30 minutes.
It had a porosity of 3 and a solid solution depleted zone 13.1 microns thick.

Example 13 A mixture was charged using the following two-stage mix-grind cycle.

Stage I The following materials were added to a 190mm internal diameter, 194mm long mill jar with a WC-Co lining along with 17.3kg of a 4.8mm WC-Co cycloid.

The mill jar was rotated on its axis at 85 revolutions per minute for 48 hours (244,800 revolutions).

Stage II: The following materials are then added to the mill jar and
Rotated for 6 hours (81,600 rpm) 6328g (87.9% by weight)
WC 140gr Sunoco3420 This mixture was balanced to produce a binder alloy of 10 w/o W and 98 w/o Co.

After discharging the slurry from the mill jar, the slurry was
It was wet screened through a misch screen, dried at 93°C under nitrogen, and then passed through a 40 mesh screen in a Fittz mill.

The insert blanks were pressed and then sintered at 1468°C for 30 minutes under a 35 micron vacuum to evaporate most of the hydrogen in the samples. During sintering, the samples were supported on NbC powder separators.

The sintered sample had A-2 porosity in the enriched zone and A-4 porosity in the unenriched bulk. The sample had an average Rockwell A hardness of 90, a 9.8 micron thick solid solution-depleted zone, and a coercive strength, Hc, of 150 oersteds.

Example 14: A batch of material having a composition equivalent to that of Example 8 was compounded, milled, and pressed into insert blanks. The blanks were then sintered, ground, heat treated, and ground (flanks only) substantially according to the procedure used in Example 8. However, a cooling rate of 69°C per hour was used in the final heat treatment.

The inserts were analyzed by EDX line scan analysis at different distances from the rake face.

The results of this analysis are shown in the graph of Figure 2. As shown, a cobalt-enriched layer extends inward from the rake surface to a depth of approximately 25 microns, followed by a layer of partially cobalt-depleted material extending inward from the rake surface to a depth of approximately 90 microns. Although not shown in the graph of Figure 2, partial solid solution depletion was found in the cobalt-enriched layer and solid solution enrichment was found in the partially cobalt-depleted layer.

The two horizontal lines represent a typical distribution of bulk material analysis values obtained for the formulation chemistry.

The above description and detailed examples are provided to illustrate some of the possible alloys, products, methods and applications that fall within the scope of the invention as defined by the claims.

[Brief explanation of the drawings]

Figure 1 is a schematic cross-sectional view of one embodiment of a coated metal cutting insert according to the present invention. Figure 2 is a graph showing typical cobalt enrichment levels produced in a cemented carbide hardmetal according to the present invention as a function of depth below the rake surface. Figure 3 is a graph showing the change in relative concentrations of binder and solid solution carbide as a function of depth below the rake surface for Example 11. 2...insert 4...rake 6...flank surface 8...cutting edge 10...coating 12...hardmetal 14...binder-enriched layer 16...binder-reduced layer 18...bulk

──────────────────────────────────────────────────── Continued from the front page (72) Inventor: George Paul Grubb 553 Hampton Street, Greensburg, Pennsylvania, USA 15601 (56) Reference: Japanese Patent Application Publication No. 1982-73104 (JP, A)

Claims (7)

[Claims] [Claim 1] A cemented carbide formed by sintering a substantially homogeneous mixture of components, said alloy consisting of 70-90.5 weight percent tungsten carbide, 5-10 weight percent metal binder selected from the group consisting of cobalt, nickel, iron and alloys thereof, and 5-20 weight percent metal carbide selected from the group consisting of transition metal carbides of titanium, tantalum, niobium and hafnium, said alloy having a degree of porosity throughout its length with pore sizes ranging from less than 10 microns in diameter to between 10 and 40 microns in diameter, said metal carbide fused with the tungsten carbide to form a carbide solid solution, said cemented carbide having a first layer near the peripheral surface of said alloy enriched in metal binder and correspondingly depleted in carbide solid solution, and a second layer below said first layer depleted in metal binder and correspondingly enriched in carbide solid solution. Claim 2: The metal binder is a cobalt binder alloy;
2. The cemented carbide of claim 1, wherein the cobalt binder alloy has a magnetic saturation of less than 158 gauss-cm ³ /gr CO. 3. The cemented carbide of claim 2, wherein the cobalt binder alloy has a magnetic saturation value of 145 to 157 gauss-cm ³ /gr Co. 4. The cobalt binder alloy having a resistivity of 126 gauss-cm ³ /g.
3. The cemented carbide of claim 2, having a magnetic saturation value less than rCo. 5. The cemented carbide of claim 1, wherein the first layer of thickened metal binder extends inward from the peripheral surface of the alloy to a depth of at least substantially 6 microns. 6. The cemented carbide of claim 5, wherein the first layer of thickened metal binder extends inward from the peripheral surface of the alloy to a depth of 12 to 50 microns. 7. The cemented carbide of claim 1, wherein the thickened layer of metal binder has a cobalt content at the peripheral surface that is 1.5 to 3 times the average cobalt content of the cemented carbide.