JPS5949298B2 - Dispersed precipitation type heat-resistant and wear-resistant sintered alloy - Google Patents

Description

[Detailed Description of the Invention] This invention has extremely high strength and toughness, and
This invention relates to a dispersed precipitation type heat-resistant and wear-resistant sintered alloy that has significantly improved wear resistance and adhesion resistance and is suitable for use as cutting tools that can be applied in particular from low-speed cutting to high-speed cutting, hot tools, etc. It is something.

Previously, the same applicant has proposed that it has strength and toughness, as well as excellent heat and abrasion resistance and erosion resistance, and is particularly effective when used as a hot tool or a cutting tool. Dispersion precipitation type heat-resistant and wear-resistant sintered alloy 0, that is, rNi: 50-70%, which exhibits excellent performance such as light and heavy cutting that is possible from low-speed cutting to high-speed cutting.
The basic composition is Ti: 2-10% and At: 0.5-10%, Fe: 0.1-10%, Co: 1-2
0chi, Cr: 1-20%, B: 0.01-1.0%
, Zr: 0.01 to 1.0%, Nb: 5% or less,
v: 5% or less, Si: 3% or less and Mn: 5% or less, in a binder phase in which one or more types are separated, Group 4a,
Carbide consisting of one or more carbides of group 5a and group 6a transition metal carbides or double carbides: 10-
A dispersion precipitation type heat-resistant and wear-resistant alloy characterized in that 90% of r is dispersed and r is precipitated in the binder phase.

” (hereinafter referred to as “prior invention alloy”) was filed in patent application No. 47-8582.
The application was filed as No. 3 (Special Publication No. 53-6922).

The present invention provides a dispersed precipitation type heat-resistant and wear-resistant sintered alloy which is a further improvement over the above-mentioned prior invention alloy. One or more types of nitrides, carbonitrides, double nitrides, and double carbonitrides of transition metals in Groups 4a, 5a, and 6a of the Table of Contents (hereinafter, these are collectively referred to as transition metal nitrides, carbonitrides, double nitrides, and double carbonitrides). When substituted with (referred to as carbon/nitride), the dispersed phase particles in the alloy structure become extremely fine, further improving the hardness and toughness at room temperature.

■ In the bonded phase, there is an r-phase, that is, (N i3A Z (T
i)) Since a larger amount of the phase is precipitated, the high temperature strength is further improved.

(b) Furthermore, a part of the dispersed phase forming components in the above-mentioned prior invention alloy is replaced with carbon/nitride of the above-mentioned transition metal, and desirable aldinium (hereinafter referred to as A7N)
By containing AlN, the excellent wear resistance and adhesion resistance of AlN will be synergistically imparted to the alloy.
The alloy now has even better wear and welding resistance, and a portion of the AAN is solidly dissolved in the binder phase, making the γ' phase precipitation much finer and more uniform. As a result, the hardness and toughness of the alloy at room and high temperatures are further significantly improved.

This was done based on the research results shown in sections (a) and (b) above.

Therefore, the present invention provides, as a dispersed phase-forming component, one or more carbides and double carbides of transition metals in groups 4a, 5a+, and 6a of the periodic table (hereinafter referred to as component A): 10 to 95% , Fe: 1 to 10%, Co: 1 to 70%, and Cr: 1 to 10% as bonding phase forming components, respectively.
20%, Mo: 1-50, W: 1-20, B:
0.01~1.0 ratio, Zr: 0.01~1.0%
, Nb: 5% or less, V: 5% or less, Si: 3% or less,
It has a composition consisting of one or more of the group consisting of Mn: 5% or less, Ti: 2 to 10%, A4: 0.5 to 10%, Ni and inevitable impurities: 20 to 70%. In addition, in a dispersed precipitation type sintered alloy having a binder phase in which the r phase, that is, the (Ni3A4(Ti)) phase is precipitated: 5 to 90% (or more by weight), a part of the dispersed phase forming components is periodically Formation of a dispersed phase consisting of one or more types (hereinafter referred to as component B) of nitrides, carbonitrides, double nitrides, and double carbonitrides of transition metals in Groups 4a, 5a, and 6a of the Table of Contents In the ingredients, A component: B component = 1:0.1~0
．． By substituting the weight ratio of 8 within a satisfying range and further containing 0.1 to 5 weight ratio of AtN as necessary, strength, toughness, heat resistance, abrasion resistance, and welding resistance can be improved. This is a dispersion precipitation type sintered alloy with improved heat and wear resistance.

Next, the reason why the substitution ratio of transition metal carbon/nitride and the content of AtN are limited as described above in the alloy of the present invention will be explained.

(a) Substitution ratio of carbon/nitride of transition metals Carbides of transition metals of groups 4a, 5a + and 6a of the periodic table,
The substitution ratio of transition metal carbon/nitride (hereinafter referred to as FB component) to one or more of the double carbides (hereinafter referred to as A component) is less than 10% by weight, that is, B component/A component = 0.1 If it is less than that, the desired characteristic improvement effect cannot be obtained,
On the other hand, if the substitution ratio exceeds 80% by weight, that is, if the B component/A component = 0.8, the sinterability will be significantly reduced.
Since the impact resistance decreases, the substitution ratio was set such that the weight ratio of component A:component B=i:0.1 to 0.8.

(b) A7N content If the content is less than 0.1%, the desired characteristic effects cannot be obtained, while if the content exceeds 5%, N1AA, Co
Since it is undesirable to form intermetallic compounds such as A7 and FeA7 and cause embrittlement, the content should be set at 0.1 to 5.
%.

Next, the dispersion precipitation type heat-resistant and wear-resistant sintered alloy of this invention (
The alloy of the present invention (hereinafter referred to as the alloy of the present invention) will be explained using examples and in comparison with the alloy of the prior invention.

Example 1 TaC powder with an average particle size of 1.5 μm as raw material powder, Ta
C powder and TaN powder, both of which have an average particle size of 2
0 μm Co powder, Ni powder, Ni-At alloy (A,
/, Content: 31.7 Ti) powder, Ti powder, Mo powder,
and Cr powder, and the final composition was 15% Ti.
c -10%TaN-10%Co-52%Ni-3%A
7-2% Ti-3% Mo-5% Cr, wet mixed, press molded, and then 10 tuiH9
Sintered by holding at a temperature of 1350°C in vacuum for 13 hours, and after sintering, holding at a temperature of 1150°C in vacuum for 4 hours, followed by oil quenching.

Alloy 1 of the present invention was produced by solution treatment, followed by incineration treatment at a temperature of 760° C. for 3 hours.

For comparison purposes, TaC powder was used instead of TiN powder as the raw material powder, and the final component composition was 15%.
TiC-10%TaC-10%Co-52%Ni-
Prior Invention Alloy 1 was produced under the same conditions as applied to the production of Invention Alloy 1, except that the raw material powders were blended so as to have 3% Al, 2% Ti, 3% Mo, and 5% Cr.

The hardness of the resulting invention alloy 1 and prior invention alloy 1 during sintering, quenching, and tempering, and the transverse rupture strength after the quenching and tempering heat treatment were measured, and the measurement results were It is shown in Table 1.

Moreover, FIG. 1 shows hardness change curves with respect to temperature for the present invention alloy 1 and the prior invention alloy 1 after the above heat treatment, and the commercially available JIS-8KH4 alloy (high speed steel).

As shown in Table 1, the alloy 1 of the present invention has better hardness and transverse rupture strength at room temperature than the alloy 1 of the prior invention. Of course, it is clear that it has extremely high high temperature hardness compared to the prior invention alloy 1, and as shown in the figure, the 5KH4 alloy has a softening point of about 600°C, and the prior invention alloy 1 has a softening point of about 600°C. It is clear that Alloy 1 of the present invention exhibits a softening point of about 800°C, which is higher than any of the alloys, while each exhibits a softening point of 700°C.

Furthermore, the above-mentioned present invention alloy 1, prior invention alloy 1, and 5
Cut out the KH4 alloy cutting tool, work material: JIS@S
NCM-8 (Hardness H: 220), Feed J: 0.2w/
rev, depth of cut: 1. Omm, cutting speed: 20,50
．． A continuous cutting test was conducted under the conditions of 80 m/min and cutting time: 5 m1n.

In addition to measuring the amount of flank wear, the welding state and the finished surface state were observed.

The results are shown in Table 2.

As seen in Table 2, it can be seen that Invention Alloy 1 exhibits superior cutting performance compared to Prior Invention Alloy 1 and SKH4 alloy.

Example 2 10% Tic powder with an average particle size of 1.5 μm, 5% WC powder with an average particle size of 1.5 μm, and 1.5 μm as dispersed phase forming components
n CrC powder: 5%, 1.5 μm TiN powder: 1
0% and the average particle size as a binder phase forming component is 2.
．． 0 μm Co powder: 1.5%, Ni powder: 40%, N
i-At alloy (Ni: 68.3%, A, la: 3
1.7%) Powder: 10chi, Ti powder: 3chi, B powder: 0
．． 1%, Nb powder: 1.9% (by weight) is prepared, and from this mixed powder, under the same conditions as in Example 1, the final component is substantially the same as the previous mixed powder composition. Invention alloy 2 having the composition was produced.

In addition, for the purpose of comparison, the composition of the dispersed phase forming components of the blended powder in the production of the above-mentioned invention alloy 2 was changed without using Tin powder, TiC powder: 20%, We powder: 5%.
%, Cr3C2 powder: 5q6, but the prior invention alloy 2 was produced under the same conditions.

When the hardness and transverse rupture strength of both of the resulting alloys were measured after heat treatment, the hardness of the prior invention alloy 2 was HRC: 6.
3. Inventive alloy 2 showed an excellent hardness HRC of 65 and 7/11j in terms of hardness HRC 200.

Example 3 WC powder with an average particle size of 1.0 μm and (W, Ti)CN powder with an average particle size of 1.5 μm as dispersed phase forming components = 20%
and the average particle size as a binder phase forming component is 2.0, respectively.
μm Co powder wood: 35%, Ni powder: 10%, NiA1
Alloy (A7: containing 31.7%) powder: 2%, Ti powder =
Prepare a blended powder consisting of 2%, ■powder = 1% (more than % by weight), wet mix, press mold, and then 0.5 m
Sintering was carried out by holding at a temperature of 1400°C for 1 hour in a N atmosphere of mHP, followed by sintering at a temperature of 1100°C in vacuum for 4 hours, followed by solution treatment of oil quenching, followed by sintering at a temperature of 70°C.
Alloy 3 of the present invention having substantially the same final component composition as the above blended powder composition was produced by subjecting it to a tempering treatment held at 0° C. for 4 hours.

In addition, for the purpose of comparison, the blending of the dispersed phase forming components in the blended powder in the production of the above-mentioned Invention Alloy 3 is shown below.

The same 1.0μm WC without containing nitride powder
Powder: 301 1.5 μm (W, Ti)C powder: Prior invention alloy 3 was produced under the same conditions except that only carbide powder consisting of 20 pieces was mixed.

When the hardness and transverse rupture strength of both of the resulting alloys were measured after heat treatment, the prior invention alloy 3 showed a hardness of HRC near 2 and a transverse rupture crab of 200 positive/-, whereas the present invention Invention Alloy 3 exhibited a higher hardness HRC of near 5 and 9/- of 220.

In addition, the above-mentioned present invention alloy 3, prior invention alloy 3, and commercially available SKH4 alloy have a workpiece material of JIS/SNCM.
-8 (Hardness HB: 220), Feed: 0, 2 vm/r
av, depth of cut: 1. A continuous cutting test was conducted under the following conditions: Om, cutting speed = 3060, and 100''/win, cutting time: !Mlin, and the cutting performance was observed in the same manner as in Example 1.

The results are shown in Table 3.

As shown in Table 3, in Example 3, Example 1
It is clear that Invention Alloy 3 exhibits superior cutting performance compared to Prior Invention Alloy 3 and SKH4 alloy.

Example 4 Tie powder with an average particle size of 1.5 μm as a dispersed phase forming component: 45%, Tac powder with an average particle size of 2 μm: 10%,
1.0 μm Mo powder: 5%, 1.5 μm NbN
Powder: 15%, Ni powder with an average particle size of 2.0 μm as a binder phase forming component: 10 cm, Ni-At alloy (Containing At: 31.7%) powder = 3 cm, Ti powder: 2 cm
%, Mo powder: 10% (by weight) was prepared, wet-mixed, press-molded, held at a temperature of 1450°C for 1 hour in a 10-degree vacuum, sintered, and then By holding the temperature at 1,160°C for 3 hours in vacuum, performing oil quenching solution treatment, and then performing tempering treatment at 720°C for 4 hours, the composition of the blended powder was substantially the same as that of the previous stage. Invention alloy 4 with the final composition was produced.

In addition, for the purpose of comparison, the blending of the dispersed phase forming components in the blended powder in the production of the above-mentioned Invention Alloy 4 was changed to 1.5 μm.
m Tie powder = 45%, same 1.2 μm Tac powder:
10%, 1.0ttm Mo powder: 5%, 1.5μ
Prior Invention Alloy 4 was produced under the same conditions except that only carbide powder consisting of 15 m of Nbc powder was used.

When the hardness and transverse rupture strength of both of the resulting alloys were measured after heat treatment, it was found that Invention Alloy 4 had a hardness of HRA:
Hardness HRA: 91.7, 160 positive/-, while prior invention alloy 4 had hardness HRA: 91.0, 150 positive/-
Showed ume/mit.

Further, for the above-mentioned present invention alloy 4, prior invention alloy 4, and commercially available Tie-based cermet, the workpiece material: JIS
-SNCM -8 (Hardness HB: 220), Feed F)
: 0.3mm/rev, depth of cut: 1.5+m,
Cutting speed: 80,150. A continuous cutting test was conducted under the conditions of 220 m/min and 10 m1n of cutting time, and the cutting performance was measured and observed in the same manner as in Examples 1 and 3.

The results are shown in Table 4. From the above results, the present invention alloy 4
It is clear that the alloy exhibits superior mechanical properties and cutting performance not only to commercially available Tie-based cermets but also to the prior invention Alloy 4.

Example 5 30% of Tie powder with an average particle size of 1.5 μm as a dispersed phase forming component, and 1.0% of Tie powder with an average particle size of 1.5 μm. u77Z Moc powder: 10%
, 1.5 μm TicN powder: 30%, 1.5 μm
VN powder: 5%, Co powder as binder phase forming component: 2%, Ni powder: 10%, Ni-At alloy (At: 3%).
Prepare a blended powder consisting of powder (containing 1.7%): 5%, Ti powder: 2.5%, W powder: 5%, and Si powder (20.5% by weight), and from this blended powder. Inventive alloy 5 was produced under the same conditions as in Example 4, having the same final component composition as the blended powder composition of the previous stage.

In addition, for the purpose of comparison, the composition of the dispersed phase forming components of the blended powder in the production of the above-mentioned invention alloy 5 was changed to
．． 5μ door Tie powder = 60%, same i, 0μm Mo,
e Powder = 10 cm, same 1.51 tm Vc powder: 5%
Prior invention alloy 5 was manufactured under the same conditions except that it was composed only of carbide powder consisting of.

The hardness and transverse rupture strength after heat treatment of both alloys obtained as a result are as follows: Inventive alloy 5 has a hardness of HRA292.0.
, the transverse rupture strength was 150 positive/ynA, and the prior invention alloy 5 had a hardness HRA of 91.1 and a transverse rupture strength of 150 positive/ynA.
It showed 130kr/#71.

Further, when the change in hardness with respect to temperature was measured for Invention Alloy 5 and the commercially available We-based cemented carbide PIO), the results shown in FIG. 2 were obtained.

Furthermore, the above-mentioned present invention alloy 5, prior invention alloy 5, and W
When used as an e-based cemented carbide cutting tool and conducted a continuous cutting test under the same cutting conditions as in Example 4,
The cutting performance shown in Table No. 5 was shown.

As is clear from the results shown in FIG. 2 and Table 5, the present invention alloy 5 has superior mechanical properties, high-temperature hardness, and cutting ability compared to the prior invention alloy 5 and the We-based cemented carbide. It has performance.

゛Example 6 Tie powder with an average particle size of 1.5 μm as a dispersed phase forming component: 25, 1. Qμm Nbc powder = 5%, same 1.
5 μm Zrc powder: 5%, 1.5 μm ZrN powder: 5%, 1.0 μm AtN powder: 3%, and Co powder each having an average particle size of 2.0 m as a binder phase forming component. : 11.5%, Ni powder: 33%, Ni-A
1 alloy (containing At: 31.7%) powder: 5%, Ti powder: 5%, Mn powder: 2T, Zr powder: 0.5T, (more than weight %) is prepared, Inventive alloy 6 having substantially the same final component composition as the previous blended powder composition was produced from this blended powder under the same conditions as in Example 3.

In addition, for the purpose of confirming the property improvement effect of the AtN component, the composition of the dispersed phase forming component of the blended powder in the production of the above-mentioned invention alloy 6 was changed to an average particle size of 1.5 μm without blending the AtN powder. Tie powder: 25%, same 1
．． Alloy 7 of the present invention was manufactured under the same conditions except that it was composed of Nbc powder of 5 μm in diameter, 5% Zrc powder in 1.5 μm diameter, and 8 grams of ZrN powder in 1.5 μm diameter.

Furthermore, for the purpose of comparison, the blending of the dispersed phase forming components was similarly carried out without blending the nitride powder, and the average particle size was 1.5A.
m Tie powder: 30%, same 1. Prior invention alloy 6 was produced under the same conditions as in the production of the invention alloy 6 above, except that it was composed of carbide powder consisting of Q μm Nbc powder: 5% and 1.5 μm Zrc powder: 13%. .

After heat treatment of the resulting invention alloys 6 and 7 and the prior invention alloy 6, the hardness (Rockwell A scale) and transverse rupture strength at room temperature, as well as the high temperature hardness (Vickers) at a temperature of 800°C were measured. The results shown in Table 6 were obtained.

As shown in Table 6, the alloy 6 of the present invention containing the AJaN component has superior mechanical properties and high temperature hardness than the alloy 7 of the present invention not containing the A7N component, and the alloy of the present invention containing nitrides. It is clear that 6°7 has superior properties compared to the prior invention alloy 6 which does not contain nitrides.

Example 7 (Ti, Mo)C powder having an average particle size of 1.8 μm as a dispersed phase forming component = 10 Ti, 2.0/! ZffL
VeC powder: 10 cm, TiN powder: 1.5 μm: 15
q6 and the average particle size as a binder phase forming component are each 2.
．． 7pm Ni powder = 49%, Ti powder: 1%, Ni-
A blended powder consisting of A1 alloy (containing 31.7% At) powder = 5mm, W powder = 5mm, Cr: powder: 5% (or more by weight) was prepared, and Example 1 was prepared using this blended powder. Inventive alloy 8 was produced under the same conditions as in Example 1, having a final component composition substantially the same as the blended powder composition.

In addition, for the purpose of comparison, the composition of the dispersed phase forming components of the blended powder in the production of the above-mentioned Invention Alloy 8 was changed to the average particle size:
(Ti9MO)・C powder with 1.8 μm: 25 Ti,
Same 2. Prior invention alloy 7 was manufactured under the same conditions except that only the carbide powder consisting of 10% Oμrn VcC powder was used.

゛When the hardness and transverse rupture strength of the resulting invention alloy 8 and prior invention alloy 7 after heat treatment were measured,
The values shown in Table 7 are shown.

Table 7 also shows that these alloys were tested under the same conditions as in Example 1 (however, the cutting speed was 50 m/yni).
The results of a continuous cutting test conducted in the case of n only) are shown.

As shown in Table 7, it is clear that Invention Alloy 8 has superior mechanical properties and cutting performance compared to Prior Invention Alloy 7.

Example 8 T having an average particle size of 1.5 μm as a dispersed phase forming component
ie powder: 15 t, 5% WeC powder of 0.8 μm, (Ti, V)N powder of 1.0 μm = 15 t, and Ni with an average particle size of 2.7117 n as a binder phase forming component.
Powder: 48%, Ti powder: 2%, Ni-At alloy (At
: Contains 31.7%) Powder: 5chi, ■Powder: 5%, Fe
Powder: A blended powder consisting of 6% (or more by weight) was prepared, and from this blended powder, under the same conditions as in Example 1,
Invention Alloy 9 was produced having a final component composition substantially the same as the previous blended powder composition.

In addition, for the purpose of comparison, the composition of the dispersed phase forming components of the blended powder in the production of the above-mentioned invention alloy 9 was changed to the average particle size:
2 parts 30% Tie powder with 1.5 μm and 0.8 μm
Prior Invention Alloy 8 was produced under the same conditions except that it was composed only of carbide powder consisting of 5% We powder of m.

The hardness and transverse rupture strength of the resulting alloy 9 of the present invention and the alloy 8 of the prior invention after heat treatment were measured and showed the values shown in Table 8.

Table 8 also shows that both alloys were tested under the same conditions as in Example 1 (however, the cutting speed was 50/min).
The results of continuous cutting tests conducted in the following cases are shown.

As shown in Table 8, it is clear that Invention Alloy 9 has superior mechanical properties and cutting performance compared to Prior Invention Alloy 8.

As mentioned above, the dispersion precipitation type heat-resistant and wear-resistant incineration alloy of the present invention has a higher hardness at room temperature and high temperature than a dispersed precipitation type heat-resistant and wear-resistant incineration alloy that does not contain nitrides due to the inclusion of nitrides. The alloy of the present invention has significantly improved toughness and high-temperature strength, and therefore has high strength, toughness, and excellent wear and welding resistance, and is particularly suitable for cutting tools and hot-working tools. It exhibits extremely excellent performance when applied to manufacturing such as.

[Brief explanation of drawings]

FIGS. 1 and 2 are curve diagrams showing changes in hardness with respect to temperature for the alloy of the present invention, the alloy of the prior invention, and the conventionally known alloy.

Claims (1)

[Claims] 1. 4a, 5. of the periodic table as dispersed phase forming components. a. and one or more carbides and double carbides of group 6a transition metals (hereinafter referred to as component A) = 10 to 9
5% each as a bonding phase forming component, Fe: 1 to 1
0%, co=1 to 70 inches, Cr: 1 to 20 inches, Mo: 1
~50%, W: 1-20 inches, B: 0.01-1.0
%, Zr: 0.01 to 1.0%, Nb: 5% or less, V: 5% or less, Si: 3% or less, Mn: 5% or less. , Ti: 2~
A binder phase having a composition of 10%, At2, 0.5% to 10%, Ni and unavoidable impurities: 20% to 70%, and in which the r phase, that is, the (NiA7(Ti)) phase was precipitated:
In the dispersion precipitation type sintered alloy having a concentration of 5 to 90% (or more by weight), a part of the dispersed phase forming components is replaced with nitrides, carbonitrides, and carbonitrides of transition metals of groups 4at, 5a+, and 6a of the periodic table. A dispersed phase forming component consisting of one or more types of double nitrides and double carbonitrides (hereinafter referred to as component B), in a weight ratio of component A: component B = 1:0.1 to 0.8. 1. A dispersion-precipitation type heat- and wear-resistant sintered alloy, characterized in that its strength, heat- and wear-resistance, and welding-resistance are improved by having a composition in which the following elements are substituted within a satisfying range. 2 As a dispersed phase forming component, 4a s5a of the periodic table
, and one or more carbides and double carbides of Group 6a transition metals (hereinafter referred to as component A): 10-
95%, each containing Fe: 1 to 1 as a bonding phase forming component
0%, co=1~70chi, Cr:1~20%, Mo1~
50%, W: 1-20 inches, B: 0.01-1°0 inches, z
One or more of the following groups: r: 0.01 to 1.0, Nb: 5 or less, V: 5 or less, Si: 3% or less, Mn: 5% or less, Ti: 2-10%, At:
0.5-10%, Ni and inevitable impurities = 20-70
It has a composition consisting of
13A Z (T i ) phase precipitated binder phase: 5
In a dispersed precipitation type sintered alloy having a content of ~90% (or more by weight), a part of the first dispersed phase forming component is
A dispersed phase forming component consisting of one or more types (hereinafter referred to as component B) of nitrides, carbonitrides, double nitrides, and double carbonitrides of transition metals of the a + 5 a s and 6a groups. , by substituting the weight ratio of component A:component B=1:0.1 to 0.8 within a satisfying range, and by creating a composition containing 0.1 to 5% by weight of aluminum nitride, strength, A dispersed precipitation type heat and wear resistant sintered alloy characterized by improved heat and wear resistance and welding resistance.