Title:
Sintered valve guide and manufacturing method thereof
Kind Code:
A1


Abstract:
Disclosed is a sintered valve guide guide formed of a sintered alloy consisting essentially of 3.5 to 5% copper, 0.3 to 0.6% tin, 0.04 to 0.15% phosphorus, 1.5 to 2.5% carbon and the balance iron, by mass, and as occasion needs, further containing 0.46 to 1.41% metal oxide, and MnS and/or magnesium silicate. The metallographic structure has: a matrix containing a pearlite phase, a Fe—P—C compound phase and a Cu—Sn alloy phase; pores; and a graphite of 1.2 to 1.7% by mass of the sintered alloy. In the cross section, the ratio of the pearlite phase to the matrix is 90 area % or more, the ratio of the Fe—P—C compound phase is 0.1 to 3 area % of the cross section, the ratio of the Cu—Sn alloy phase to the cross section is 1 to 3% by area, and the ratio of a portion of the Fe—P—C compound phase having a thickness of 15 microns or more is 10 area % or less of the whole Fe—P—C compound phase.



Inventors:
Chikahata, Katsunao (Chiba, JP)
Hayashi, Koichiro (Chiba, JP)
Fujitsuka, Hiroki (Chiba, JP)
Tsuboi, Toru (Chiba, JP)
Application Number:
11/182638
Publication Date:
02/16/2006
Filing Date:
07/14/2005
Primary Class:
Other Classes:
419/11
International Classes:
C22C33/02
View Patent Images:



Primary Examiner:
ZHU, WEIPING
Attorney, Agent or Firm:
Kilpatrick Townsend & Stockton LLP - East Coast (ATLANTA, GA, US)
Claims:
What is claimed is:

1. A sintered valve guide formed of a sintered alloy consisting essentially of 3.5 to 5% copper, 0.3 to 0.6% tin, 0.04 to 0.15% phosphorus, 1.5 to 2.5% carbon and the balance iron, by mass, wherein the sintered alloy has a metallographic structure comprising: a matrix having a pearlite phase, a Fe—P—C compound phase and a Cu—Sn alloy phase; pores; and a graphite phase being dispersed at a ratio of 1.2 to 1.7% by mass of the sintered alloy, and wherein, in a cross section of the metallographic structure of the sintered alloy, the ratio of the pearlite phase to the matrix is 90% by area or more, the ratio of the Fe—P—C compound phase to the cross section of the metallographic structure is 0.1 to 3% by area, the ratio of the Cu—Sn alloy phase to the cross section of the metallographic structure is 1 to 3% by area, and the ratio of a portion of the Fe—P—C compound phase having a thickness of 15 microns or more to whole of the Fe—P—C compound phase is 10% by area or less.

2. A sintered valve guide formed of a sintered alloy consisting essentially of 3.5 to 5% copper, 0.3 to 0.6% tin, 0.04 to 0.15% phosphorus, 1.5 to 2.5% carbon, 0.46 to 1.41% metal oxide and the balance iron, by mass, wherein the sintered alloy has a metallographic structure comprising: a matrix having a pearlite phase, a Fe—P—C compound phase, a Cu—Sn alloy phase and a metal oxide phase; pores; and a graphite phase being dispersed at a ratio of 1.2 to 1.7% by mass of the sintered alloy, and wherein, in a cross section of the metallographic structure of the sintered alloy, the ratio of the pearlite phase to the matrix is 90% by area or more, the ratio of the Fe—P—C compound phase to the cross section of the metallographic structure is 0.1 to 3% by area, the ratio of the Cu—Sn alloy phase to the cross section of the metallographic structure is 1 to 3% by area, and a portion of the Fe—P—C compound phase having a thickness of 15 microns or more to whole of the Fe—P—C compound phase is 10% by area or less.

3. A sintered valve guide formed of a sintered alloy consisting essentially of 3.5 to 5% copper, 0.3 to 0.6% tin, 0.04 to 0.15% phosphorus, 1.5 to 2.5% carbon, 1% or less of at least one solid lubricant selected from the group consisting of manganese sulfide and magnesium silicate minerals, and the balance iron, by mass, wherein the sintered alloy has a metallographic structure comprising: a matrix having a pearlite phase, a Fe—P—C compound phase and a Cu—Sn alloy phase; pores; a graphite phase; and said at least one solid lubricant being dispersed in the pores or intergranularly dispersed in the metallographic structure, and wherein, in a cross section of the metallographic structure of the sintered alloy, the ratio of the pearlite phase to the cross sectioin of the metallographic structure is 80% by area or more, the ratio of the Fe—P—C compound phase to cross section of the metallographic structure is 0.1 to 3% by area, the ratio of the Cu—Sn alloy phase to the cross section of the metallographic structure is 1 to 3% by area, the graphite phase is dispersed in the pores at a ratio of 0.8 to 3.2% by area of the cross section of the metallographic structure, and the ratio of a portion of the Fe—P—C compound phase having a thickness of 15 microns or more to whole of the Fe—P—C compound phase is 10% by area or less.

4. A sintered valve guide formed of a sintered alloy consisting essentially of 3.5 to 5% copper, 0.3 to 0.6% tin, 0.04 to 0.15% phosphorus, 1.5 to 2.5% carbon, 0.46 to 1.41% metal oxide, 1.6% or less of at least one solid lubricant selected from the group consisting of manganese sulfide and magnesium silicate minerals, and the balance iron, by mass, wherein the sintered alloy has a metallographic structure comprising: a matrix having a pearlite phase, a Fe—P—C compound phase, a Cu—Sn alloy phase and a metal oxide phase; pores; a graphite phase; and said at least one solid lubricant being dispersed in the pores or intergranularly dispersed in the metallographic structure, and wherein, in a cross section of the metallographic structure of the sintered alloy, the ratio of the pearlite phase to the cross sectioin of the metallographic structure is 90% by area or more, the ratio of the Fe—P—C compound phase to cross section of the metallographic structure is 0.1 to 3% by area, the ratio of the Cu—Sn alloy phase to the cross section of the metallographic structure is 1 to 3% by area, the graphite phase is dispersed in the pores at a ratio of 0.8 to 3.2% by area of the cross section of the metallographic structure, and the ratio of a portion of the Fe—P—C compound phase having a thickness of 15 microns or more to whole of the Fe—P—C compound phase is 10% by area or less.

5. The sintered valve guide as set forth in claim 2, wherein the metal oxide includes at least one oxide of a metal selected from the group consisting of aluminum, silicon, magnesium, iron, calcium and tin.

6. The sintered valve guide as set forth in claim 1, wherein, in the cross section of the metallographic structure, the ratio of another portion of the Fe—P—C compound phase having a thickness of 5 microns or more and less than 15 microns to whole of the Fe—P—C compound phase is 10 to 40% by area, and the balance Fe—P—C compound phase has a thickness of less than 5 microns.

7. A method of manufacturing a sintered valve guide, comprising: preparing a powder mixture comprising, by mass: 0.27 to 0.7% of a Fe—P alloy powder; 3.93 to 5.44% of a Cu—Sn alloy powder; 1.7 to 2.7% of a graphite powder and the balance iron powder, wherein the Fe—P alloy powder consists essentially of 15 to 21% phosphorus, an inevitable amount of impurities and the balance iron, the Cu—Sn alloy powder consists essentially of 8 to 11% tin, an inevitable amount of impurities and the balance copper; forming the powder mixture into a tubular compact by pressing the powder mixture in a tubular cavity; and sintering the tubular compact in a non-oxidizing atmosphere at a sintering temperature of 950 to 1,050 degrees C.

8. The manufacturing method as set forth in claim 7, wherein the iron powder includes an ore reduced iron powder containing 0.5 to 1.5% by mass of metal oxide.

9. The manufacturing method as set forth in claim 7, wherein the iron powder is a mixture of an ore reduced iron powder and an atomized iron powder, wherein the content of the atomized iron powder in the mixture is 10 to 30% by mass.

10. The manufacturing method as set forth in claim 7, wherein the iron powder has a maximum particle size of 104 to 200 microns.

11. The manufacturing method as set forth in claim 7, wherein the Fe—P alloy powder has a maximum particle size of 61 to 104 microns, and the Cu—Sn alloy powder has a maximum particle size of 35 to 61 microns.

12. The manufacturing method as set forth in claim 7, wherein the sintering time is 15 to 90 minutes.

13. The manufacturing method as set forth in claim 7, wherein the preparing of the powder mixture further comprises: mixing at least one powder of a manganese sulfide and magnesium silicate minerals into the powder mixture to adjust the ratio of said at least one powder in the powder mixture to 1.6% by mass or less.

14. The manufacturing method as set forth in claim 7, wherein the tubular cavity is tapered at a ratio of 1/5000 to 1/1000 by inclining at least one of an inner bore surface of a die and a circumferential surface of a punch which define the tubular cavity.

15. The manufacturing method as set forth in claim 7, further comprising: dipping in an oil the sintered compact obtained by the sintering.

16. The manufacturing method as set forth in claim 7, further comprising: cooling the sintered compact obtained by the sintering at a cooling rate of 8 degrees C./min.

Description:

BACKGROUND OF THE INVENTION

1. Field of the Invention

The present invention relates to a valve guide for use in internal-combustion engines, and in particular, a sintered valve guide made of a sintered alloy superior in wear resistance and machinability, and a method of manufacturing the same.

2. Related Art

Valve guides made of cast iron have been conventionally used in internal-combustion engines, and sintered alloy valve guides are also used recently, from the viewpoints of wear resistance and mass production. The applicant of the present application has proposed valve guide materials superior in wear resistance, for example, in the patent applications (Japanese Patent Application Publication (JP-B) No. 55-34858 and Japanese Patent 2680927).

The valve guide material disclosed in the document of JP-B No. 55-34858 has a composition consisting of C: 1.5 to 4%, Cu: 1 to 5%, Sn: 0.1 to 2%, P: 0.1 to 0.3%, and Fe: the balance, by weight ratio, and has a structure that iron-phosphorus-carbon compound phase made of an eutectic Fe—P—C compound, Cu—Sn phase and free graphite are dispersed in the matrix which is mixture of pearite and ferrite. It has wear resistance better than that of conventional cast iron valve guides and a machinability better that that of the conventional iron-based sintered alloys, and thus has been used in many car manufacturers although it's processing is relatively difficult in comparison with the cast iron products.

On the other hand, the valve guide material disclosed in the document of Japanese Patent 2680927, which is a valve guide material improved from that disclosed in JP-B No. 55-34858, has machinability improved by dispersing a magnesium silicate mineral into the area of grain boundary in the metallographic structure of JP-B No. 55-34858, without impairing the wear resistance.

The valve guide material disclosed in Japanese Patent 2680927 exhibits an excellent wear resistance which is equivalent to that of the valve guide material disclosed in JP-B No. 55-34858, but its machinability is lower than that of cast iron products although improved to some extent. Therefore, further improvement in the machinability is desirable. The applicant of the present application have conducted studies aimed at improving the machinability even at sacrifice of wear resistance to some extent, and the valve guide material disclosed in Japanese Patent Application Laid-Open (JP-A) No. 2002-69597 has been developed.

The valve guide material disclosed in JP-A No. 2002-69597 has a composition consisting of C: 1.5 to 4%, Cu: 1 to 5%, Sn: 0.1 to 2%, P: 0.01 to 0.1%, and Fe: the balance, by weight ratio, and has a structure containing free graphite dispersed in the matrix whose main body is pearlite.

However, in a trend toward improving efficiency of the manufacturing process, there exist an increasing demand for improvement in the processability of valve guide materials and thus an increasing need for a valve guide material more superior in machinability.

BRIEF SUMMARY OF THE INVENTION

With the above problems in mind, it is therefore an object of the present invention to provide a novel sintered valve guide higher in durability which is manufactured efficiently with use of a sintered alloy having well-balanced wear resistance and machinability as a valve guide material, and a manufacturing method thereof.

In order to achieve the above-mentioned object, a sintered valve guide, according to one aspect of the present invention, is formed of a sintered alloy consisting essentially of 3.5 to 5% copper, 0.3 to 0.6% tin, 0.04 to 0.15% phosphorus, 1.5 to 2.5% carbon and the balance iron, by mass, wherein the sintered alloy has a metallographic structure comprising: a matrix having a pearlite phase, a Fe—P—C compound phase and a Cu—Sn alloy phase; pores; and a graphite phase being dispersed at a ratio of 1.2 to 1.7% by mass of the sintered alloy, and wherein, in a cross section of the metallographic structure of the sintered alloy, the ratio of the pearlite phase to the matrix is 90% by area or more, the ratio of the Fe—P—C compound phase to the cross section of the metallographic structure is 0.1 to 3% by area, the ratio of the Cu—Sn alloy phase to the cross section of the metallographic structure is 1 to 3% by area, and the ratio of a portion of the Fe—P—C compound phase having a thickness of 15 microns or more to whole of the Fe—P—C compound phase is 10% by area or less.

A sintered valve guide, according to another aspect, is formed of a sintered alloy consisting essentially of 3.5 to 5% copper, 0.3 to 0.6% tin, 0.04 to 0.15% phosphorus, 1.5 to 2.5% carbon, 0.46 to 1.41% metal oxide and the balance iron, by mass, wherein the sintered alloy has a metallographic structure comprising: a matrix having a pearlite phase, a Fe—P—C compound phase, a Cu—Sn alloy phase and a metal oxide phase; pores; and a graphite phase being dispersed at a ratio of 1.2 to 1.7% by mass of the sintered alloy, and wherein, in a cross section of the metallographic structure of the sintered alloy, the ratio of the pearlite phase to the matrix is 90% by area or more, the ratio of the Fe—P—C compound phase to the cross section of the metallographic structure is 0.1 to 3% by area, the ratio of the Cu—Sn alloy phase to the cross section of the metallographic structure is 1 to 3% by area, and a portion of the Fe—P—C compound phase having a thickness of 15 microns or more to whole of the Fe—P—C compound phase is 10% by area or less.

A sintered valve guide, according to another aspect, is formed of a sintered alloy consisting essentially of 3.5 to 5% copper, 0.3 to 0.6% tin, 0.04 to 0.15% phosphorus, 1.5 to 2.5% carbon, 1.6% or less of at least one solid lubricant selected from the group consisting of manganese sulfide and magnesium silicate minerals, and the balance iron, by mass, wherein the sintered alloy has a metallographic structure comprising: a matrix having a pearlite phase, a Fe—P—C compound phase and a Cu—Sn alloy phase; pores; a graphite phase; and said at least one solid lubricant being dispersed in the pores or intergranularly dispersed in the metallographic structure, and wherein, in a cross section of the metallographic structure of the sintered alloy, the ratio of the pearlite phase to the cross sectioin of the metallographic structure is 90% by area or more, the ratio of the Fe—P—C compound phase to cross section of the metallographic structure is 0.1 to 3% by area, the ratio of the Cu—Sn alloy phase to the cross section of the metallographic structure is 1 to 3% by area, the graphite phase is dispersed in the pores at a ratio of 0.8 to 3.2% by area of the cross section of the metallographic structure, and the ratio of a portion of the Fe—P—C compound phase having a thickness of 15 microns or more to whole of the Fe—P—C compound phase is 10% by area or less.

A sintered valve guide, according to another aspect, is formed of a sintered alloy consisting essentially of 3.5 to 5% copper, 0.3 to 0.6% tin, 0.04 to 0.15% phosphorus, 1.5 to 2.5% carbon, 0.46 to 1.41% metal oxide, 1.6% or less of at least one solid lubricant selected from the group consisting of manganese sulfide and magnesium silicate minerals, and the balance iron, by mass, wherein the sintered alloy has a metallographic structure comprising: a matrix having a pearlite phase, a Fe—P—C compound phase, a Cu—Sn alloy phase and a metal oxide phase; pores; a graphite phase; and said at least one solid lubricant being dispersed in the pores or intergranularly dispersed in the metallographic structure, and wherein, in a cross section of the metallographic structure of the sintered alloy, the ratio of the pearlite phase to the cross sectioin of the metallographic structure is 90% by area or more, the ratio of the Fe—P—C compound phase to cross section of the metallographic structure is 0.1 to 3% by area, the ratio of the Cu—Sn alloy phase to the cross section of the metallographic structure is 1 to 3% by area, the graphite phase is dispersed in the pores at a ratio of 0.8 to 3.2% by area of the cross section of the metallographic structure, and the ratio of a portion of the Fe—P—C compound phase having a thickness of 15 microns or more to whole of the Fe—P—C compound phase is 10% by area or less.

A method of manufacturing a sintered valve guide, according to one aspect of the invention, comprises: preparing a powder mixture comprising, by mass: 0.27 to 0.7% of a Fe—P alloy powder; 3.93 to 5.44% of a Cu—Sn alloy powder; 1.7 to 2.7% of a graphite powder and the balance iron powder, wherein the Fe—P alloy powder consists essentially of 15 to 21% phosphorus, an inevitable amount of impurities and the balance iron, the Cu—Sn alloy powder consists essentially of 8 to 11% tin, an inevitable amount of impurities and the balance copper; forming the powder mixture into a tubular compact by pressing the powder mixture in a tubular cavity; and sintering the tubular compact in a non-oxidizing atmosphere at a sintering temperature of 950 to 1,050 degrees C.

In accordance with the above construction, a sintered valve guide that is higher in durability is efficiently manufactured with use of a valve guide material having well-balanced abrasion resistance and machinability.

BREIF DESCRIPTION OF THE DRAWINGS

FIG. 1 is a schematic view illustrating the cross section of the metallographic structure of a sintered valve guide according to the present invention.

FIG. 2 is a schematic view illustrating the cross section of the metallographic structure of a sintered valve guide having a higher ratio of iron-phosphorus-carbon compound phase.

FIGS. 3A-3D are graphs showing the relationships of the phosphorus content in the whole composition with the ratio of iron-phosphorus-carbon compound phase (FIG. 3A), the ratio of copper-tin alloy phase (FIG. 3B), the amount of free graphite phase (FIG. 3C), and the ratio of ferrite phase (FIG. 3D).

FIGS. 4A-4D are graphs showing the relationships of the phosphorus content in the whole composition with the thickness of the iron-phosphorus-carbon compound phase (FIG. 4A), wear amount (FIG. 4B), machinability index (FIG. 4C), and radial crushing strength constant (FIG. 4D).

FIGS. 5A-5D are graphs showing the relationships of the tin content in the whole composition with the ratio of iron-phosphorus-carbon compound phase (FIG. 5A), the ratio of copper-tin alloy phase (FIG. 5B), the amount of free graphite phase (FIG. 5C), and the ratio of ferrite phase (FIG. 5D).

FIGS. 6A-6D are graphs showing the relationships of the tin content in the whole composition with the thickness of the iron-phosphorus-carbon compound (FIG. 6A), wear amount (FIG. 6B), machinability index (FIG. 6C), and radial crushing strength constant (FIG. 6d).

FIGS. 7A-7D are graphs showing the relationships of the copper content in the whole composition with the ratio of iron-phosphorus-carbon compound phase (FIG. 7A), the ratio of copper-tin alloy phase (FIG. 7B), the amount of free graphite (FIG. 7C), and the ratio of ferrite phase (FIG. 7D).

FIGS. 8A-8D are graphs showing the relationships of the copper content in an entire composition with the thickness of the iron-phosphorus-carbon compound phase (FIG. 8A), wear amount (FIG. 8B), machinability index (FIG. 8C), and radial crushing strength constant (FIG. 8D).

FIGS. 9A-9D are graphs showing the relationships of the carbon content in the whole composition with the ratio of iron-phosphorus-carbon compound phase (FIG. 9A), the ratio of copper-tin alloy phase (FIG. 9B), the amount of free graphite (FIG. 9C), and the ratio of ferrite phase (FIG. 9D).

FIGS. 10A-10D are graphs showing the relationships of the carbon content in the whole composition with the thickness of the iron-phosphorus-carbon compound phase (FIG. 10A), wear amount (FIG. 10B), machinability index (FIG. 10C), and radial crushing strength constant (FIG. 10D)

FIGS. 11A-11D are graphs showing the relationships of the sintering temperature with the ratio of iron-phosphorus-carbon compound phase (FIG. 11A), the ratio of copper-tin alloy phase (FIG. 11B), the amount of free graphite (FIG. 11C), and the ratio of ferrite phase (FIG. 11D).

FIG. 12A-12D are graphs showing the relationships of the sintering temperature with the thickness of the iron-phosphorus-carbon compound phase (FIG. 12A), wear amount (FIG. 12B), machinability index (FIG. 12C), and radial crushing strength constant (FIG. 12D).

FIGS. 13A-13D are graphs showing the relationships of the sintering time with the ratio of iron-phosphorus-carbon compound phase (FIG. 13A), the ratio of copper-tin alloy phase (FIG. 13B), the amount of free graphite phase (FIG. 13C), and the ratio of ferrite phase (FIG. 13D).

FIGS. 14A-14D are graphs showing the relationships of the sintering time with the thickness of the iron-phosphorus-carbon compound phase (FIG. 14A), wear amount (FIG. 14B), machinability index (FIG. 14C), and radial crushing strength constant (FIG. 14D).

FIGS. 15A-15D are graphs showing the relationships of the cooling rate with the ratio of iron-phosphorus-carbon compound phase (FIG. 15A), the ratio of copper-tin alloy phase (FIG. 15B), the amount of free graphite phase (FIG. 15C), and the ratio of ferrite phase (FIG. 15D).

FIGS. 16A-16D are graphs showing the relationships of the cooling rate with the thickness of iron-phosphorus-carbon compound phase (FIG. 15A), wear amount (FIG. 15B), machinability index (FIG. 15C), and radial crushing strength constant (FIG. 15D).

FIGS. 17A-17D are graphs showing the relationships of the oxide content in iron powder with the ratio of iron-phosphorus-carbon compound phase (FIG. 17A), the ratio of copper-tin alloy phase (FIG. 17B), the amount of free graphite phase (FIG. 17C), and the ratio of ferrite phase (FIG. 17D).

FIGS. 18A-18D are graphs showing the relationships of the oxide content in iron powder with the thickness of the iron-phosphorus-carbon compound phase (FIG. 18A), wear amount (FIG. 18B), machinability index (FIG. 18C), and radial crushing strength constant (FIG. 18D).

FIGS. 19A-19D are graphs showing the relationships of the amount of atomized iron powder with the ratio of iron-phosphorus-carbon compound phase (FIG. 19A), the ratio of copper-tin alloy phase (FIG. 19B), the amount of free graphite phase (FIG. 19C), and the ratio of ferrite phase (FIG. 19D).

FIGS. 20A-20D are graphs showing the relationships of the amount of atomized iron powder with the thickness of the iron-phosphorus-carbon compound phase (FIG. 20A), wear amount (FIG. 20B), machinability index (FIG. 20C), and radial crushing strength constant (FIG. 20D).

FIGS. 21A-21D are graphs showing the relationships of the amount of machinability-improving component powder with the ratio of iron-phosphorus-carbon compound phase (FIG. 21A), the ratio of copper-tin alloy phase (FIG. 21B), the amount of free graphite phase (FIG. 21C), and the ratio of ferrite phase (FIG. 21D).

FIGS. 22A-22D are graphs showing the relationships of the amount of machinability-improving powder with the thickness of the iron-phosphorus-carbon compound phase (FIG. A), wear amount (FIG. 22B), machinability index (FIG. 22C), and radial crushing strength constant (FIG. 22D).

DETAILED DESCRIPTION OF THE INVENTION

Sintered alloys produced by powder metallurgy have metallographic structures different from each other, according to the composition and the particle diameters of the raw powders used and the manufacturing conditions such as heating temperature, etc., even if the whole composition is the same; and the material properties such as mechanical strength of the sintered alloys vary significantly depending on the structure of the sintered alloys. In the present invention, taking into consideration the influence on the material properties of the various phases which are present in the sintered alloy, the metallographic structure of the valve guide material is designed to impart to the sintered alloy the material properties required for the final valve guide material, and it is made a base for determining the raw materials and the manufacturing conditions.

Valve guides demands both high strength and high wear resistance. Conventional alloys for valve guides satisfy these requirements but are still insufficient in machinability, and thus there exists a strong need from the users for improvement of the nonconformity during processing. Accordingly, to meet the needs of the users, the present invention is aimed at improving the machinability of valve guides, on the basis of an alloy containing a matrix containing iron as the main component, and, for improving wear resistance, a copper-tin alloy phase, an iron-phosphorus-carbon compound phase and a free graphite phase. Hereinafter, the metallographic structure of the sintered alloy for the sintered valve guide according to the present invention will be described. Here, it is noted, the ratio of each phase in the cross section of the metallographic structure in the following description by area % is indicated as an average value.

For improvement of strength, the matrix of the sintered alloy is constructed with a pearlite structure, which is produced by diffusion of carbon into a raw iron powder at sintering of the raw iron powder mixed with a graphite powder. Since metal powders containing carbon as a solid solution are hard and less compressible, an iron powder and a graphite powder are used as raw materials. Lack in amount of the graphite powder leads to decrease in the amount of the carbon that binds to the base metal, and consequently to the decrease in the strength of the matrix due to a large amount of the ferrite (a-iron) phase formed in the matrix. In this connection, it is noted that, as will be described below, formation of the iron-phosphorus-carbon compound phases is often accompanied by a small amount of ferrite phases produced around the steadite phase. However, if pearlite occupies 90% or more, by area ratio, of the matrix, production of ferrite in the balance portion is still in the allowable range, because the deterioration in the strength of matrix is still small.

The iron-phosphorus-carbon compound phase is dispersed in the pearlite matrix. By blending an iron-phosphorus alloy powder into the raw iron powder with the graphite powder and sintering the raw powder mixture, the iron-phosphorus-carbon compound precipitates in the thin film form on the grain boundary of the pearlite phases, forming hard iron-phosphorus-carbon compound phase and thus raising the wear resistance of the sintered alloy. The increase in wear resistance is more significant when the ratio of the iron-phosphorus-carbon compound phases appeared in the cross section of the metallographic structure is 0.1 area % or more. On the other hand, increase in the amount of the iron-phosphorus-carbon compound phases produced leads to increase in the thickness of the film and growth into plate-shaped iron-phosphorus-carbon compound phase, and consequently to drastic deterioration in the machinability of sintered alloy. Thus, it is important to reduce the amount of the iron-phosphorus-carbon compound phases produced and disperse the iron-phosphorus-carbon compound phases in a thin form for prevention of the decrease in the machinability of sintered alloy. Specifically, in the cross section of metallographic structure, the ratio of the iron-phosphorus-carbon compound phase in the cross section of the metallographic structure should be 3 area % or less, and the ratio of the a portion of the iron-phosphorus-carbon compound phase having a thickness of 15 microns or more should be 10 area % or less in the whole iron-phosphorus carbon compound phase. More specifically, the ratio of a portion of the iron-phosphorus-carbon compound phase having a thickness of 15 microns or more is preferably 0.1 area % or less in the entire iron-phosphorus-carbon compound phase; the ratio of another portion having a thickness of 5 microns or more and less than 15 microns is preferably 10 to 40 area %; and the other portion of the iron-phosphorus-carbon compound phase are less than 5 microns in thickness, and such a structure is preferably formed. The iron-phosphorus-carbon compound phase often deprives the pearlite matrix of carbon during generation, causing production of some ferrite phase surrounding the iron-phosphorus-carbon compound phase. The ferrite phase, which is lower in strength, may be allowed to produce in an amount in the range of 10 area % or less in the matrix of sintered alloy, but presence thereof in a greater amount is undesirable. That is, an excessively large phosphorus content of raw powder may lead to production of thick iron-phosphorus-carbon compound phase and, at the same time, by depriving the pearlite matrix of carbon, to dispersion of the ferrite phases lower in strength into the matrix in a greater amount. Thus, the amount of the iron-phosphorus-carbon compound phase produced should be controlled carefully so as to prevent the above problems, and specifically, the ratio of the iron-phosphorus-carbon compound phase in the cross section of the metallographic structure should be kept in the range of 0.1 to 3 area %. Accordingly, the amount of the iron-phosphorus alloy powder to use should be adjusted so as to make the phosphorus content in sintered alloy 0.04 to 0.15 mass %. If attaching importance to strength or wear resistance, the amount of iron-phosphorus alloy powder is 0.1 to 0.15 mass %, and, if machinability is more important, 0.04 to 0.1 mass %.

Copper-tin alloy phase is dispersed in the sintered alloy according to the invention. The copper-tin alloy phase is soft and effective in raising the wear resistance of the sintered alloy by improving the compatibility with valves or sliding counterparts. The effects of the copper-tin phase are more significant when it is dispersed in the matrix at a ratio of 1 area % or more of the cross section area of the metallographic structure, but the dimensional stability of the product at sintering is impaired by the expansion of the copper during sintering when this phase is present at a ratio of more than approximately 3 area %, and thus, the blending ratio in the raw powder is preferably adjusted so that the ratio of the copper-tin alloy phase in the cross section area is 1 to 3 area %. The copper-tin alloy phase possibly produces by using a simple copper powder and a simple tin powder as raw materials, but in that case, the composition and the distribution of the alloy phase produced from them in the sintered alloy fluctuate more widely, deteriorating the dimensional stability and the wear resistance of sintered alloy. Therefore, use of a copper-tin alloy powder is desirable. When the copper-tin alloy phase is dispersed in the pearlite matrix in a fine form such as has a maximum grain size of 20 microns or less and is present at a ratio of 80 area % or more in the entire copper-tin alloy phase, uniformity in the dispersion of copper-tin alloy phase becomes high and it is more effective from the view point of compatibility. Depending on the bridging state of the powder particles on the preparation of green compact, the copper-tin alloy powder may remain undiffused and a copper-tin alloy phase having a grain size of 150 microns or more may be produced by the undiffused copper-tin alloy powder. However, if its ratio is 5 area % or less with respect to the entire copper-tin alloy phase in the cross section, such an amount of presence is allowable. The copper-tin alloy powder produces a liquid phase during sintering, which contributes to accelerate sintering, and copper and tin respectively reinforce the matrix as they are dissolved in the matrix as solid solution. However, an excessive amount of copper contained as solid solution leads to the significant deterioration in dimensional stability due to expansion of copper and an excessive amount of tin to embrittlement of the matrix. For prevention of the problems above and more favorable dispersion of the copper-tin alloy phases, the tin content in the copper-tin alloy powder used is preferably 8 to 11 mass %. As a result, the composition of the copper-tin alloy phase in the sintered alloy also becomes in the range similar to above. Accordingly, the suitable contents of copper and tin in the whole alloy composition are respectively 3.5 to 5.0 mass % and 0.3 to 0.6 mass %, as determined based on the composition of copper-tin alloy powder and the area ratio of the copper-tin alloy phases in the cross section.

It is preferable but not essential that a ttace amount of metal oxide phase is dispersed additionally in the pearlite matrix. This metal oxide phase is composed of oxide of at least one metal selected from the group consisting of aluminum, silicon, magnesium, iron, titanium and calcium, and the oxide works as an easy-cutting component and improves machinability. However, an excessive content of the metal oxide leads to embrittlement of the matrix, and the metal oxide in an amount of 0.46 to 1.41 mass % in the whole composition is preferably dispersed in the pearlite matrix. It is important to disperse such metal oxide phase uniformly in the matrix, and use of an iron powder containing the metal oxide(s) as the raw powder is desirable. Iron powders containing the metal oxide(s) at a total amount of 0.5 to 1.5 mass % are preferable, and such iron powders include ore reduced iron powders. Atomized iron powders and mill scale reduced iron powders commonly used have a lower content of metal oxide.

In addition, the free graphite phase is dispersed in the metallographic structure. This phase originates from the raw graphite powder and works as a solid lubricant to contribute to improvement of the machinability and the wear resistance of sintered alloy. Although it is difficult to determine the ratio of free graphite phase in sintered alloy accurately from the microscopic photograph of alloy cross section because of the loss during sample preparation, it is possible to determine the mass ratio of free graphite according to the method of quantitatively determining free carbon specified by JIS-G1211 “Method for determination of carbon content”, and, based on the graphite content obtained above and the specific gravity of graphite, the effect of the free graphite phase in relation to the ratio of free graphite phase is obtainable. According to the above, the effect of the free graphite phase is more significant when the ratio of free graphite phase is approximately 0.8 area % or more, and production of a sintered alloy containing the free graphite phase in an amount of more than 3.2 area % leads to precipitation of hard cementite (Fe3C) in the matrix and consequently to deterioration in machinability. An excessive amount of graphite powder also impairs the compressibility of powder and reduces the ratio of the matrix in the sintered alloy, resulting in deterioration of the strength of sintered alloy. Accordingly, the ratio of free graphite phase is preferably approximately 0.8 area % or more and 3.2 area % or less in the alloy cross section.

For further improvement of the machinability, at least one of manganese sulfide (MnS) powder and magnesium silicate mineral powder can be blended as a solid lubricant in a total amount of 1.6 mass % or less with respect to the raw powder, and it is dispersed in the pores or at the powder grain boundary of the sintered alloy. The MnS phase and the magnesium silicate phases improve wear resistance and are particularly effective in improving machinability. In addition, the MnS phase protects the blade of cutting tool and extends the tool lifetime, while the magnesium silicate minerals, which have a cleavage property that cleave easily during cutting, are effective in reducing machining force. Both components have a chip-breaking action and are effective in preventing the blade of tool from being heated and elongating the tool lifetime by breaking chips into smaller pieces.

The sintered alloy described above having free graphite phase, iron-phosphorus-carbon compound phase, copper-tin alloy phase, metal oxide phase and, as occasion requires, solid lubricant phase in the metallographic structure can be produced by the method of production described below, and sintered valve guides are manufactured, by settling the compacting step in the production of the sintered alloy so that the sintered alloy obtained has a shape suitable as valve guides. As a result, the sintered alloy and the sintered valve guide produced have the following composition of copper: 3.5 to 5 mass %, tin: 0.3 to 0.6 mass %, phosphorus: 0.04 to 0.15 mass %, carbon: 1.5 to 2.5 mass %, the metal oxide: 0.46 to 1.41 mass %, and iron: the balance. The amount of the solid lubricant, if used, is 1 mass % or less with respect to the whole alloy composition.

Valve guides commonly used have a density of 6.3 to 6.9 g/cm3 and contain pores in an amount of approximately 4 to 15% by density ratio. The sintered valve guide according to the present invention described above is the same from this point.

In producing the sintered alloy and the sintered valve guide, a powder mixture is prepared first. Raw materials including a graphite powder, an iron-phosphorus alloy powder, a copper-tin alloy powder, an iron powder, and, as occasion requires, a solid lubricant powder are blended uniformly, to give a powder mixture. The raw powders will be described respectively in detail below.

An Example of the raw iron powder for the pearlite matrix is an atomized iron powder having a particle size approximately of −150 to −65 meshes (minus sieve of screen of 65 to 150 meshes, maximum particle size: 104 to 200 microns). Alternatively, an ore reduced iron powder having a particle size of −150 to −65 meshes that contains the metal oxide(s) described above at a total metal oxide amount of 0.5 to 1.5 mass % is favorably used, because it improves machinability. Ore reduced iron powders contain metal oxides in a greater amount inherently because of the preparative methods, and the metal oxides are effective in improving machinability. A metal oxide content below the above range may lead to decrease in the improvement of machinability. A metal oxide content higher than the above range leads to hardening and deterioration in compressibility of the powder and it is thus undesirable. Ore reduced iron powders are porous and absorb the copper-tin alloy liquid phase generated during sintering by capillary action, and use of them makes the distribution of the components in the sintered alloy more uniform. Ore reduced iron powders having a larger particle size make it difficult to increase the density of powder, while those having a smaller particle size reduce the flowability of powder, and thus, powders having a size approximately of −150 to −65 meshes are favorable. However, ore reduced iron powders, which contain a greater amount of metal oxides, are slightly harder and lower in compressibility than atomized iron powders or the like, and thus, the strength of the sintered valve guides obtained is slightly lower than that of the atomized iron powders. Thus, the raw iron powder for sintered valve guide is selected according to one of the requirements in properties: strength or machinability. Alternatively, combined use of a powder mixture consisting of an ore reduced iron powder and an atomized iron powder, for example a mixture containing the atomized iron powder at a ratio of 10 mass % or more, improves the compressibility of the powder and the strength of the sintered alloy obtained. In such a case, an atomized iron powder substitution ratio of more than 30 mass % may result in irregular distribution of metal oxides and prohibition of the improvement in machinability. Therefore, the substitution ratio is preferably 10 to 30 mass % when an atomized and ore reduced iron powder are used in combination. The particle size of the atomized iron powder is settled to be the same as or less than that of the ore reduced iron powder.

The iron-phosphorus alloy powder is a material for supplying phosphorus, which is used in the form of alloy for safer-handling because phosphorus is unstable and flammable. Phosphorus diffuses into the iron base, raising the strength of the pearlite matrix metal and the wear resistance thereof by producing the iron-phosphorus-carbon compound phased. Iron-phosphorus alloys having a phosphorus content of approximately 10 to 13 mass % generate the liquid phase of the iron-phosphorus alloy in the temperature range of 950 to 1,050° C. A greater amount of liquid phase is undesirable as it impairs the dimensional stability during sintering, but a suitable amount of liquid phase promotes growth of the necks and improves the strength of the sintered alloy. Therefore, an iron-phosphorus alloy powder having a phosphorus content of 15 mass % or more is favorably used for proper control of generation of the liquid phase. Phosphorus in the iron-phosphorus alloy powders having a phosphorus content of 15 mass % or more diffuses into the iron powders during sintering; and when the phosphorus content reaches the range above in some areas, the alloy powders generate liquid phases. The liquid phase wets the surface of iron powder, phosphorus therein diffuses rapidly from the liquid phase into the iron powder, decreasing the phosphorus content in the liquid phase to less than the range described above and making the liquid phase solidify. Accordingly, the iron-phosphorus alloy improves the strength by facilitating the growth of necks among iron particles and prevents extreme deterioration in dimensional stability by restricting liquid phase generation partially and making the liquid phase solidify in a shorter period of time. When the phosphorus content of the iron-phosphorus alloy powder used is less than 15 mass %, the composition of the iron-phosphorus alloy reach the range above generating liquid phase by diffusion of phosphorus during sintering, and as a result, the liquid phase is formed more rigorously, which leads to deterioration of dimensional stability and reduction in the amount of the iron-phosphorus-carbon compound phase generated due to the diffusion of phosphorus into the entire matrix. On the other hand, when the phosphorus content of iron-phosphorus alloy powder is more than 21 mass %, the iron-phosphorus alloy powder becomes harder and the compressibility of powder mixture is reduced so that the green compact and sintered alloys become lower in density, resulting in lower strength of the sintered valve guides obtained. In addition, the iron-phosphorus-carbon compound phase generated becomes thicker, reducing the machinability of the sintered alloy. Thus, use of an iron-phosphorus alloy powder having a phosphorus content of 15 to 21 mass % is preferable, and the amount thereof used is preferably approximately 0.27 to 0.7 mass % with respect to the total powder mixture amount. Use of an iron-phosphorus alloy powder having a particle size of approximately −250 to −150 meshes (maximum particle size: 61 to 104 microns) is preferable from the point of the compressibility of powder mixture. The iron-phosphorus alloy powder used may contain inevitable amounts of impurities, for example, carbon, silicon, manganese and others in a total amount in the range of 1.5 mass % or less.

The copper-tin alloy powder, which is used for reduction in size and increase in dispersion homogeneity of the copper-tin alloy phase in sintered alloy, preferably has a particle size smaller than that of the iron powder. For prevention of the deterioration in dimensional stability by expansion of copper and the embrittlement of matrix by excessive solid solution of tin, use of a copper-tin alloy powder having a tin content of 8 to 11 mass % is preferable, and the amount thereof used is preferably approximately 3.93 to 5.44 mass % with respect to the total powder mixture amount. According to the above, the copper content and the tin content in powder mixture become respectively, approximately 3.5 to 5.0 mass % and 0.3 to 0.6 mass %. A copper-tin alloy powder having a particle size smaller than that of the iron powder is preferably used for uniform dispersion of the copper-tin alloy powder in powder mixture and generation of fine copper-tin alloy phase uniformly dispersed in sintered alloy. Preferably, a copper-tin alloy powder having a particle diameter of −250 to −400 meshes (max. particle size: 35 to 61 microns) is used. The copper-tin alloy powder may contain inevitable amounts of impurities.

The graphite powder forms a pearlite structure and generates an iron-phosphorus-carbon compound by bonding to the iron powder and the iron-phosphorus alloy powder during sintering, and the residual graphite forms a free graphite phase. The suitable carbon content in the sintered alloy is 1.5 to 2.5 mass %, but, considering the consumption for reduction of the metal oxides contained in iron powder and loss by reaction with water in environmental air, the amount of the graphite powder is preferably, approximately 1.7 to 2.7 mass % with respect to the total powder mixture amount. However, an excessive amount of graphite causes precipitation of cementite in the pearlite matrix and also reduces the compressibility of powder mixture, resulting in decrease in the density of green compact and sintered alloy and the strength of valve guides. When the particle size of the graphite powder used is extremely smaller, the free graphite phases remaining after sintering becomes scarce, while the graphite powder having a excessively larger particle size makes the compressibility of powder mixture significantly smaller, the distribution of the components in the sintered alloy drastically irregular, and ferrite phase tends to generate surrounding the iron-phosphorus-carbon compound phases.

As described above, the solid lubricants used are a MnS powder and/or a magnesium silicate mineral powder, and examples of the magnesium silicate minerals include magnesium metasilicate mineral and magnesium orthosilicate mineral. Typical examples of the magnesium metasilicate minerals include enstatite, clinoenstatite, enstenite, hypersten and the like, and examples of magnesium orthosilicate minerals include forstelite, chrysolite and the like. When using the solid lubricant, the amount of it is preferably 1 mass % or less with respect to the total powder mixture, for prevention of deterioration of the strength of sintered alloy.

The powder mixture obtained by uniformly mixing the raw powders as described above is then compression molded into a green compact, using a mold. A mold of a suitable shape for a desirable product is used during the compacting, and for use in preparing valve guides, it has a long round tubular cavity. Specifically, a mold consisting of a die having a cylindrical hole, a cylindrical core rod that forms a long tubular cavity with the die as it is placed at the center of the die cavity, and upper and lower punches having a ring-shaped cross section to be inserted into the cavity is used. A green compact is prepared by placing the lower punch in the cavity, filling the powder mixture in the cavity, placing the upper punch over the powder mixture, and compressing the powder mixture between the upper and lower punches in the axial direction. At this time, the compression pressure is preferably adjusted so that the resulting green compact has a green density approximately of 6.5 to 7.1 g/cm3.

In connection with the compacting step described above, since the shape of the compact is long in the axial direction, the compression pressure may not reach the middle part of the compact in the axial direction, giving a green compact having a compression density smaller in the middle part in the axial direction than those at the both ends. In such a case, the strength of the sintered valve guide obtained becomes lower in the middle region in the axial direction. For overcoming the problem, it is effective to incline at least one of the cavity wall of die and the circumferential surface of core rod so that the cavity becomes tapered slightly with respect to the punch-moving direction. If the taper angle is small, the effect on the green compact dimension can be compensated by the spring-back effect of powder particles, and thus it becomes possible to make uniform the compression density, by making the compression pressure reach the middle part in the axial direction without a substantially adverse influence on the dimension of the product. The taper rate is preferably, approximately 1/5,000 to 1/1,000, and a rate of less than 1/5,000 results in insufficient improvement in the compression density of the middle part, while a ratio of more than 1/1,000 in apparent variation in diameter between two ends of the green compact.

The molded green compact is sintered at 350 to 1,050° C. in a nonoxidative atmosphere and then cooled. When sintered at a temperature in the range above, graphite reacts with the iron powder, forming a pearlite structure. In addition, a part of the iron-phosphorus alloy powder forms a liquid phase facilitating formation of the sintering bonds by diffusion of the powders, while phosphorus diffuses in the iron powder and reacts with graphite, forming iron-phosphorus-carbon compound phase, and the iron-phosphorus-carbon compound phase precipitate in the cooling process after sintering. The copper-tin alloy powder forms a liquid phase in the heating process during sintering, accelerating sintering and diffusing copper and tin into the iron base. The sintered alloy after sintering and cooling contains iron-phosphorus-carbon compound phase dispersed and precipitated in the thin film form and minute copper-tin alloy phase dispersed and precipitated from the liquefied copper and tin, in the pearlite structure. The metallographic structure is formed in a period of approximately five minutes, but elongation of the sintering time improves the strength of the sintered product by further growth of the necks among iron powder particles, and thus the sintering time is preferably 20 minutes or more, more preferably 45 minutes or more, from the point of strength. However, sintering at a sintering temperature of more than 1,050° C. or for a sintering time of more than 90 minutes advances the diffusion of graphite into. the matrix further, reducing the residual free graphite amount, increasing the amount of precipitated iron-phosphorus-carbon compound, and increasing the thickness of the iron-phosphorus-carbon compound phase. Consequently, it leads to drastic deterioration in machinability. On the other hand, at a sintering temperature of less than 950° C., sintering does not proceed sufficiently, giving an undesirable metallographic structure and reducing the strength markedly. Since the speeds of generation and diffusion of the liquid phase vary in accordance with the heating temperature during sintering, the green compact is preferably sintered at a higher temperature for a shorter sintering period or at a lower temperature for a longer period, for obtaining a favorable metallographic structure. Because slower cooling leads to increase in the amount of the iron-phosphorus-carbon compound phase and the ferrite phase precipitated and also in the thickness thereof, the cooling rate is preferably approximately 8° C./min or more, and more preferably approximately 10° C./min or more.

A crude sintered valve guide is obtained by the sintering above, and subsequent step for high-precision mechanical processing of the internal face thereof with a reamer gives a finished sintered valve guide. The present invention achieves improvement of the machinability of the sintered alloy for sintered valve guide, and thus shortens the period needed for mechanical processing with a reamer and reduces processing defects.

Immersion of the crude sintered valve guide in oil, which causes adsorption of the oil in the pores by capillary action, is effective for raising the air tightness of the sintered valve guide. The oil also functions as a lubricant oil during the mechanical processing, improving the machinability. The oil may be embedded into the pores of the crude sintered valve guide forcibly by vacuum deaeration during immersion. In addition, dispersion of molybdenum disulfide or the like in the oil is favorable for improving machinability and wear resistance.

The cross section of the metallographic structure in the sintered valve guide according to the present invention obtained as described above is shown schematically in FIG. 1. The metallogrpahic structure comprises a matrix, pores P and graphite phase G dispersed in the pores P, and the matrix contains: pearlite phase PE that may contain metal oxide phase MO; copper-tin alloy phase CS; and an iron-phosphorus-carbon compound phase FPC. The iron-phosphorus-carbon compound phase FPC are dispersed thinly, surrounded by a very small amount of ferrite phase F.

FIG. 2 is a schematic view illustrating the cross section of the metallographic structure of a conventional sintered alloy described, for example, in the document of JP-B No. 55-34858 that has a greater phosphorus content. In this case, the ratio of a thick portion of the iron-phosphorus-carbon compound phases, having a thickness of 15 microns or more, is high and a ferrite phase which is formed by carbon deficiency is present in a greater amount, surrounding the iron-phosphorus-carbon compound phase. Sintered alloys having such a metallographic structure are poorer in machinability and strength than those having the structure shown in FIG. 1. Namely, the sintered alloy described in document of JP-B No. 55-34858 above contains the thick iron-phosphorus-carbon compound phase shown in FIG. 2.

EXAMPLES

Hereinafter, the present invention will be described in more detail with reference to Examples.

Example 1

(Samples 1 to 27)

Using an ore-reduced iron powder (metal oxide content: 0.1 mass %) or an atomized iron powder (metal oxide content: 0.2 mass %) as the iron powder, powder mixture samples 1 to 27 were prepared, respectively, by blending the iron powder with an iron-phosphorus alloy powder, a copper-tin alloy powder and a graphite powder in the mixing ratios shown in Table 1. The composition of each powder mixture sample is shown in Table 2. The particle diameters of the powders used were respectively as follows: ore-reduced iron powder (150 microns or more: 5%, 45 to 150 microns: 75%, and less than 45 microns: 20%), atomized iron powder (150 microns or more: 17%, 45-150 microns: 58%, and less than 45 microns: 25%), iron-phosphorus alloy powder (63 microns or more: 3%, 45 to 63 microns: 10%, and less than 45 microns: 87%), and copper-tin alloy powder (150 microns or more: 7%, 45-150 microns: 73%, and less than 45 microns: 20%) and graphite powder (average particle diameter: 0.6 to 0.8 mm).

Each powder mixture sample was compressed at a pressure of 550 MPa into a round tubular green compact (for abrasion test and machinability test) having an outer diameter of 11 mm, an inner diameter of 6 mm and a length of 40 mm, and a ring-shaped green compact (for radial crushing test) having an outer diameter of 18 mm, an inner diameter of 10 mm and a length of 10 mm, both of which were sintered at a temperature of 1,000° C. for 60 minutes in a nonoxidative atmosphere, then cooled from 1,000 to 600° C. at a rate of 12° C./min and allowed to cool to room temperature, to obtain each of sintered compact samples 1 to 27.

The cross section of the metal structure in each of the sintered compact samples 1 to 27 was observed under microscope (×340), and the ratios of the iron-phosphorus-carbon compound phase and the copper-tin alloy phase in the cross section of metallographic structure (area %), the ratio of the ferrite phase in the matrix (area %), the ratio of the free graphite (mass %), and the ratios (area %) of the regions, respectively, having a thickness of: less than 5 microns; 5 microns or more and less than 15 microns; and 15 microns or more, in the iron-phosphorus-carbon compound phase were determined. Results are summarized in Table 3.

Separately, each of the sintered compact samples 1 to 27 was subjected to the following abrasion, machinability and radial crushing tests for determination of the wear amount, the machinability index and the radial crushing strength constant of the samples. Results are summarized in Table 3.

(Abrasion Test)

The abrasion test of each tubular sintered compact was carried out in a vertical valve guide abrasion tester. In the abrasion test, a valve stem was connected to the bottom end of a piston having its shaft line in the vertical direction, and a valve was inserted inside the sintered compact and moved back and forth in an exhaust gas atmosphere at 500° C. while applying a lateral load of 3 MPa to the piston. The stroke speed was then 3,000 rpm and the stroke width was 8 mm. After reciprocating motion for 30 hours, the wear amount (micron) of the inner bore face of sintered compact was determined.

(Machinability Test)

The inner bore face of a tubular sintered compact was reamed by using a reamer made of sintered hard alloy, and the period necessary for machining the sintered compact to a depth of 10 mm in the axial direction was measured. The necessary period of each sample was converted to an index value with respect to 100 of the necessary period of the sample 13 (having the alloy composition equivalent to that described in JP-B No. 55-034858, hereinafter referred to as conventional alloy). Here, a smaller index indicates that the sintered compact is more easily machinable and has a shorter processing time, i.e., that it is superior in machinability.

(Radial Crushing Test)

A ring-shaped sintered compact was pressed in the diametrical direction gradually with an increasing pressure load until the sintered compact broke down according to the method regulated by JIS Z2507 “Sintering bearing—Determination of radial crushing strength”. The radial crushing strength constant K (N/mm2) was calculated from the maximum pressure load above in accordance with the following formula (1) (wherein, F: maximum load (N) when the compact broke down; L: length of ring-shaped compact (mm); D: outer diameter of ring-shaped compact (mm); and e: wall thickness of ring-shaped compact (mm)).
K=F(D−e)/(L×e2) (1)

TABLE 1
Mixing Ratio mass %
Fe—P AlloyCu—Sn Alloy
PowderPowder
Comp.Comp.
SampleFe Powdermass %mass %Graphite
NoTypeFePCuSnPowderComments
0192.45ore reduced0.1580.0020.005.0090.0010.002.40Fe—P add.: <min.; P cont.: <min.
0292.33ore reduced0.2780.0020.005.0090.0010.002.40Fe—P add.: min.
0392.25ore reduced0.3580.0020.005.0090.0010.002.40
0492.20ore reduced0.4080.0020.005.0090.0010.002.40
0592.10ore reduced0.5080.0020.005.0090.0010.002.40Fe—P add.: max.
0691.90ore reduced0.7080.0020.005.0090.0010.002.40
0791.60ore reduced1.0080.0020.005.0090.0010.002.40Fe—P add.: >max.; P cont.: >max.
0892.33ore reduced0.2785.0015.005.0090.0010.002.40Fe—P add.: min.; P in Fe—P: min.; P cont.: min.
0991.90ore reduced0.7079.0021.005.0090.0010.002.40Fe—P add.: max.; P in Fe—P: max.; P cont.: max.
1091.20atomized1.4080.0020.005.0090.0010.002.40corresp. to alloy in JPB55-34858
1192.46ore reduced0.4080.0020.004.7495.005.002.40Sn in Cu—Sn: <min.; Sn cont.: <min.
1292.31ore reduced0.4080.0020.004.8992.008.002.40Sn in Cu—Sn: min.
1392.14ore reduced0.4080.0020.005.0689.0011.002.40Sn in Cu—Sn: max.
1491.91ore reduced0.4080.0020.005.2985.0015.002.40Sn in Cu—Sn: >max.; Sn cont.: >max.
1594.20ore reduced0.4080.0020.003.0090.0010.002.40Cu—Sn add.: <min.; Cu in Cu—Sn: <min.
1693.20ore reduced0.4080.0020.004.0090.0010.002.40Cu—Sn add.: <min.
1792.70ore reduced0.4080.0020.004.5090.0010.002.40
1891.80ore reduced0.4080.0020.005.4090.0010.002.40Cu—Sn add.: >max.
1991.20ore reduced0.4080.0020.006.0090.0010.002.40Cu—Sn add.: >max.; Cu in Cu—Sn: >max.
2093.27ore reduced0.4080.0020.003.9392.008.002.40Cu—Sn add.: min.; Sn in Cu—Sn: min.;
Sn cont.: min.
2191.76ore reduced0.4080.0020.005.4489.0011.002.40Cu—Sn add.: max.; Sn in Cu—Sn: max.;
Sn cont.: max.
2293.27ore reduced0.4080.0020.003.9389.0011.002.40Cu—Sn add.: min.; Sn in Cu—Sn: max.;
Cu cont.: min.
2391.76ore reduced0.4080.0020.005.4492.008.002.40Cu—Sn add.: max.; Sn in Cu—Sn: min.;
Cu cont.: max.
2493.40ore reduced0.4080.0020.005.0090.0010.001.20graphite add.: <min.; C cont.: <min.
2592.90ore reduced0.4080.0020.005.0090.0010.001.70graphite add.: min.; C cont.: min.
2691.90ore reduced0.4080.0020.005.0090.0010.002.70graphite add.: max.; C cont.: max.
2791.60ore reduced0.4080.0020.005.0090.0010.003.00graphite add.: >max.; C cont.: >max.

TABLE 2
Composition mass %
SampleMetal
NoFePCuSnCOxideComments
0191.650.034.500.502.200.92Fe—P add.: <min.; P cont.: <min.
0291.620.054.500.502.200.92Fe—P add.: min.
0391.610.074.500.502.200.92
0491.600.084.500.502.200.92
0591.580.104.500.502.200.92Fe—P add.: max.
0691.540.144.500.502.200.92
0791.480.204.500.502.200.92Fe—P add.: >max.; P cont.: >max.
0891.640.044.500.502.200.92Fe—P add.: min.; P in Fe—P: min.; P cont.: min.
0991.530.154.500.502.200.92Fe—P add.: max.; P in Fe—P: max.; P cont.: max.
1092.320.284.500.502.200.18corresp. to alloy of JPB-55-34858
1191.860.084.500.242.200.92Sn in Cu—Sn: <min.; Sn cont.: <min.
1291.710.084.500.392.200.92Sn in Cu—Sn: min.
1391.540.084.500.562.200.92Sn in Cu—Sn: max.
1491.310.084.500.792.200.92Sn in Cu—Sn: >max.; Sn cont: >max.
1593.580.082.700.302.200.94Cu—Sn add.: <min.; Cu in Cu—Sn: <min.
1692.590.083.600.402.200.93Cu—Sn add.: <min.
1792.090.084.050.452.200.93
1891.200.084.860.542.200.92Cu—Sn add.: >max.
1990.610.085.400.602.200.91Cu—Sn add.: >max.; Cu in Cu—Sn: >max.
2092.660.083.620.312.200.93Cu—Sn add.: min.; Sn in Cu—Sn: min.; Sn cont.: min.
2191.160.084.840.602.200.92Cu—Sn add.: max.; Sn in Cu—Sn: max.; Sn cont.: max.
2292.660.083.500.432.200.93Cu—Sn add.: min.; Sn in Cu—Sn: max.; Cu cont.: min.
2391.160.085.000.442.200.92Cu—Sn add.: max.; Sn in Cu—Sn: min.; Cu cont.: max.
2492.790.084.500.501.000.93graphite add.: <min.; C cont.: <min.
2592.290.084.500.501.500.93graphite add.: min.; C cont.: min.
2691.300.084.500.502.500.92graphite add.: max.; C cont.: max.
2791.000.084.500.502.800.92graphite add.: >max.; C cont.: >max.

TABLE 3
Ratio in Met. Section area %
Fe—P—CEstimation Items
Compound PhaseFerriteRadial
ThicknessCu—SnFreePhase inWearMachin-Crushing
SampleμmAlloyGraphiteMatrixAmountabilityStr. Con.
No<55-15≧15Phasemass %area %μmIndexMPaComments
010.1091902.21.880.08116.3486Fe—P add.: <min.; P cont: <min.
021.00881112.21.660.05422.5536Fe—P add.: min.
031.50841332.31.520.04725.0579
041.90811542.31.421.24628.8596
052.20761862.21.332.64330.0613Fe—P add.: max.
062.90682482.21.225.74033.8638
075.904535202.41.0219.13491.3637Fe—P add.: >max.; P cont.: >max.
080.30901002.31.670.05617.5517Fe—P add.: min.; P in Fa—P: min.; P cont.: min.
093.006525102.31.246.04033.8635Fe—P add.: max.; P in Fe—P: max.; P cont.: max.
1010.00145542.30.7126.130100.0602corresp. to alloy of JPB55-34858
112.10761772.11.421.26113.8472Sn in Cu—Sn: <min.; Sn cont.: <min.
122.10801552.21.421.25225.6584Sn in Cu—Sn: min.
132.00841422.31.421.24630.6600Sn in Cu—Sn: max.
142.10851412.51.421.24436.9529Sn in Cu—Sn: >max.; Sn cont.: >max.
152.50741880.61.421.67240.0411Cu—Sn add.: <min.; Cu in Cu—Sn: <min.
162.30791651.21.421.45832.5547Cu—Sn add.: <min.
172.10801641.71.421.35130.0581
181.80831522.81.420.04425.0601Cu—Sn add.: >max.
191.60851414.11.420.04021.3587Cu—Sn add.: >max.; Cu in Cu—Sn: >max.
202.40761681.11.421.47610.6488Cu—Sn add.: min.; Sn in Cu—Sn: min.;
Sn cont.: min.
211.50851412.81.420.04232.5608Cu—Sn add.: max.; Sn in Cu—Sn: max.;
Sn cont.: max.
222.30761771.01.421.55736.3518Cu—Sn add.: min.; Sn in Cu—Sn: max.;
Cu cont.: min.
231.90841513.01.420.04520.6609Cu—Sn add.: max.; Sn in Cu—Sn: min.;
Cu cont.: max.
241.3090912.20.710.68246.3621graphite add.: <min.; C cont: <min.
251.50851322.31.020.95537.5607graphite add.: min.; C cont: min.
262.70722172.31.703.64023.8527graphite add: max.; C cont.: max.
273.905232162.22.036.03721.9394graphite add.: >max.; C cont.: >max.

The content of phosphorus in the whole composition changes among samples 1 to 9; the content of phosphorus in the iron-phosphorus alloy powder is kept constant among samples 1 to 7; and the contents of phosphorus both in the entire composition and in the iron-phosphorus alloy powder change among samples 8 and 9. The relationships among the phosphorus content in the whole composition and the ratios of the respective phases in these samples and sample 10 (the sintered alloy disclosed in JP-B No. 55-034853, referred to below as conventional alloy) are shown in FIGS. 3A-3D, and the relationship of the phosphorus content with the thickness of the iron-phosphorus-carbon compound phase and the material properties (wear amount, machinability index and radial crushing strength constant) are shown in FIGS. 4A-4D.

As apparent from FIGS. 3A, 3C and 3D and FIG. 4A, the ratio of copper-tin alloy phase does not vary according to the increase in the phosphorus content (FIG. 3B), but the ratio and thickness of iron-phosphorus-carbon compound phase increase drastically according to increase in the phosphorus content, and the amount of free graphite phase decreases and that of ferrite phase increases in the phosphorus content range of more than 0.15 mass %. The results indicate that formation of iron-phosphorus-carbon compound is accelerated by the increase in phosphorus content but formation of the iron-phosphorus-carbon compound consumes free graphite and the carbon dissolved in the matrix, consequently-increasing the ratio of ferrite phase. In addition, as shown in FIGS. 4B, 4C and 4D, the sintered compacts having a phosphorus content of less than 0.04 mass % are superior in machinability, but have a greater wear amount and a smaller radial crushing strength constant. When the phosphorus content is 0.04 mass % or more, the radial crushing strength increases and the wear amount decreases, but the machinability index decline, according to increase in phosphorus content. In particular when the phosphorus content is 0.15 mass % or more, the machinability index increases significantly.

As shown in FIGS. 3D, 4A and 4C, fluctuation in machinability index has significant correlations with the ratio of the iron-phosphorus-carbon compound phase produced and the ratio of its portion having a thickness of 15 microns or more. In the phosphorus content range of 0.15 mass % or more, the ratio of the iron-phosphorus-carbon compound phase having a thickness of 15 microns or more becomes 10 area % or less and the machinability index becomes smaller at 35 or less. That is, fineness or reduction in dimensions of the iron-phosphorus-carbon compound phase leads to improvement in machinability.

On the other hand, fluctuation of the radial crushing strength has a correlation with the ratio of iron-phosphorus-carbon compound phase. The radial crushing strength constant increases along with increase in the ratio of iron-phosphorus-carbon compound phase, and the sintered compacts have a sufficiently high radial crushing strength constant of approximately 500 MPa or more in the phosphorus content range of 0.04 mass % or more, in which the ratio of iron-phosphorus-carbon compound phase becomes 0.1 area % or more, and a radial crushing strength higher than that of the conventional alloy is exhibited at a phosphorus content of 0.1 mass % or more. However, the ratio of ferrite phase, which is lower in strength, increases along with the increase in phosphorus content, and a phosphorus content of 0.2 mass % or more on the contrary leads to deterioration in strength.

Fluctuation in wear amount has a correlation with the ratio of the iron-phosphorus-carbon compound phase generated, the wear amount of sintered compacts decreases drastically in the range where the phosphorus content is 0.04 mass % or more and the ratio of iron-phosphorus-carbon compound phase is 0.1 area. That is, the wear amount decreases along with the increase in phosphorus content and in the ratio of the iron-phosphorus-carbon compound phase produced.

The results above indicate that valve guides favorable in all of radial crushing strength, machinability and abrasion resistance can be obtained, when the phosphorus content in the whole composition is 0.04 to 0.15 mass %, the ratio of iron-phosphorus-carbon compound phase is 0.1 to 3 area %, and the ratio of a portion of the iron-phosphorus-carbon compound phase having a thickness of 15 ,,m or more relative to the whole iron-phosphorus-carbon compound phase is 10% or less as area ratio.

In samples 4 and 11 to 23, the tin content and/or the copper content in the whole composition change. More specifically, the copper content is constant among samples 4 and 11 to 14, the composition of the copper-tin alloy powder used is the same among samples 4 and 15 to 19, and the tin and copper contents in the whole composition and the composition of copper-tin alloy powder change among samples 20 to 23. For these samples and sample 13, the relationship of the tin content in the whole composition and the ratios of respective phases is shown in FIGS. 5A-5D, and the relationships of the tin content with the thickness of the iron-phosphorus-carbon compound phase and material properties (wear amount, machinability index and radial crushing strength constant) are shown in FIGS. 6A-6D. In addition, the relationship between the copper content in the whole composition and the ratios of respective phases is shown in FIGS. 7A-7D, and the relationships of the copper content with the thickness of the iron-phosphorus-carbon compound phase and material properties (wear amount, machinability index and radial crushing strength constant) in FIGS. 8A-8D.

The change in the ratio of copper-tin alloy phase shown in FIG. 5B varies significantly according to the manner of adding tin, and, considering these results and those in FIG. 7B, the ratio of copper-tin alloy phase seems to be more dependent on copper content rather than on tin content. In addition, FIGS. 5A, 5C, 5D and FIG. 6A also indicate the tin content exerts a smaller influence on the configuration of the phases in metallographic structure.

In contrast, the wear resistance and the radial crushing strength increase while the machinability decreases, along with increase in tin content. It is because of increase in the amount of the tin dissolved in the matrix. When the tin content is 0.3 mass % or more in the whole composition, the sintered compacts have favorable wear resistance and radial crushing strength, i.e., a wear amount of 70 microns or less and a radial crushing strength constant of 500 MPa or more. From the point of machinability, a tin content of 0.6 mass % or less is preferable.

According to FIG. 7A and FIG. 8A, increase in copper content leads to decrease in the ratio and the thickness of iron-phosphorus-carbon compound phase, resulting in improvement in the machinability as shown in FIG. 8C. It is because the copper improves the hardenability of the matrix and makes fast the apparent cooling speed (influence by cooling speed will be described below in detail.). In addition, FIG. 7B indicates that increase in copper content raises the ratio of copper-tin alloy phase. The presence of a copper-tin alloy phase, which is soft and superior in compatibility, improves the wear resistance and facilitates processing, and the sintered compacts show a favorable machinability index of 35 or more in the copper content range of 3.5 mass % or more (FIG. 8C). Here, it is noted that the copper-tin alloy also improves the radial crushing strength (FIG. 8D), but reduces the strength at a content of 5.0 mass % or more because of the excessive softness of the copper-tin alloy phase. From these results, the content of tin is favorably in the range of 0.3 to 0.6 mass %, and that of copper in the range of 3.5 to 5.0 mass %.

In samples 4 and 24 to 27, the carbon content in the whole composition is altered. The relationships between the carbon content in the whole composition and the ratios of respective phases of these samples and sample 13 are shown in FIGS. 9A-9D, and the relationships of the carbon content with the thickness of the iron-phosphorus-carbon compound phase and the material properties (wear amount, machinability index and radial crushing strength constant) are shown in FIGS. 10A-10D.

It is understood from FIGS. 9A and 10A, the ratio of iron-phosphorus-carbon compound phase and the thickness thereof increase along with the increase in carbon content. The amount of free graphite phase also increases (FIG. 9C). However, FIG. 9D shows that the ratio of ferrite phase increases drastically at a carbon content of more than 2.5 mass %, indicating that the production of the iron-phosphorus-carbon compound phase is accompanied with removal of carbon from the matrix.

The machinability index shown in FIG. 10C is influenced by the thickness of the iron-phosphorus-carbon compound phase as well as the ratio of free graphite phase. That is, the machinability index becomes minimal in the carbon content range of 1.5 to 2.5 mass %, in combination of: the reduction in machinability index (improvement in machinability) caused by the increase of the free graphite phase; and the increase in machinability index (deterioration in machinability) caused by the increase and growth of the iron-phosphorus-carbon compound phase. As to the wear amount, it is regarded to decline with increase in the ratio of iron-phosphorus-carbon compound phase. However, the radial crushing strength constant, which is expected to be improved along with the increase in the ratio of iron-phosphorus-carbon compound phase, decreases as shown in FIG. 10D, especially drastically at a carbon content of over 2.5 mass %. It is because the increase in the carbon powder content leads to deterioration of the compressibility of powder mixture and consequently to deterioration in the strength of both green and sintered compacts. From these results, the range of the carbon content most preferable from the points of radial crushing strength, machinability and wear resistance is 1.5 to 2.5 mass %.

Example 2

(Samples 28 to 38)

Sintered compact samples 28 to 31 were prepared by using a powder mixture having the same composition as that for sample 4 of Example 1 and by repeating the operations (preparation of powder mixture, formation of green compact, sintering, and cooling) similar to those for sample 4, except that the sintering temperature was changed to 900° C. (sample 28), 950° C. (sample 29), 1,050° C. (sample 30), or 1,100° C. (sample 31) as shown in Table 4. The composition of the entire powder mixture of each sample is shown in Table 5.

Separately, sintered compact samples 32 to 36 were prepared by repeating the operations similar to those for sample 4 (preparation of powder mixture, formation of green compact, sintering, and cooling), except that the sintering period was changed to 10 minutes (sample 32), 20 minutes (sample 33), 45 minutes (sample 34), 90 minutes (sample 35), or 120 minutes (sample 36).

Yet separately, sintered compact samples 37 and 38 were prepared by repeating the operations similar to those for sample 4 (preparation of powder mixture, formation of green compact, sintering, and cooling), except that the cooling rate after sintering was changed to 8° C./min (sample 37) or 4° C./min (sample 33).

By using each of the sintered compact samples 28 to 38, the ratios of respective phases in the cross section of the metallographic structure of the sintered compact were determined by cross sectional area observation, and the abrasion, machinability and radial crushing tests thereof were performed in a similar manner to samples 1 to 27. The results are summarized in Table 6.

TABLE 4
Mixing Ratio mass %
OreFe—20PCu—10SnSint.Sint.Cooling
SampleReducedAlloyAlloyGraphiteTemp.TimeRate
NoFe Pwd.PowderPowderPowder° C.min.° C./min.Comments
0492.200.405.002.4010006012
2892.200.405.002.409006012sint. temp.: <min.
2992.200.405.002.409506012sint. temp.: min.
3092.200.405.002.4010506012sint. temp.: max.
3192.200.405.002.4011006012sint. temp.: >max.
3292.200.405.002.4010001012
3392.200.405.002.4010002012
3492.200.405.002.4010004512
3592.200.405.002.4010009012
3692.200.405.002.40100012012sint. time: >max.
3792.200.405.002.401000608cooling rate: min.
3892.200.405.002.401000604cooling rate: <min.

TABLE 5
Composition mass %
SampleMetal
NoFePCuSnCOxideComments
0491.600.084.500.502.200.92
2891.600.084.500.502.200.92sint. temp.: <min.
2991.600.084.500.502.200.92sint. temp.: min.
3091.600.084.500.502.200.92sint. temp.: max.
3191.600.084.500.502.200.92sint. temp.: >max.
3291.600.084.500.502.200.92
3391.600.084.500.502.200.92
3491.600.084.500.502.200.92
3591.600.084.500.502.200.92
3691.600.084.500.502.200.92sint. time: >miax.
3791.600.084.500.502.200.92cooling rate: min.
3891.600.084.500.502.200.92cooling rate:
<min.

TABLE 6
Ratio in Met. Section area %
Fe—P—CEstimation Items
Compound PhaseFerriteRadial
ThicknessCu—SnFreePhase inWearCrushing
SampleμmAlloyGraphiteMatrixAmountMachinabilityStr. Con.
No<55-15≧15Phasemass %mass %μmIndexMPaComments
041.90811542.31.421.24628.8596
280.000002.02.0396.69510.0297sint. temp.: <min.
290.29851502.21.528.96815.0509sint. temp.: min.
302.93553692.21.223.84335.0615sint. temp.: max.
314.443943182.20.5126.74042.5476sint. temp.: >max.
321.81821802.11.929.46611.9387
331.93732612.21.724.65818.8542
342.11702912.31.521.55025.6588
352.86573852.41.226.24234.4584
363.873845172.50.6119.43841.9567sint. time: >max.
372.96563592.31.424.34330.0591cooling rate: min.
385.333842202.31.4218.63850.6582cooling rate: <min.

The relationships between the sintering temperature and the ratios of respective phases in the whole composition in samples 4 and 28 to 31 are shown in FIGS. 11A-11D, and the relationships of the sintering temperature with the thickness of the iron-phosphorus-carbon compound phase and the material properties (wear amount, machinability index and radial crushing strength constant) are shown in FIGS. 12A-12D.

According to FIG. 11B, the ratio of the copper-tin alloy phase is independent of the sintering temperature. On the other hand, the iron-phosphorus-carbon compound phase is formed when it is higher than 900° C., and the ratio and the thickness thereof increase along with elevation in the sintering temperature (FIGS. 11A and 12A), while the ratio of free graphite decreases on the contrary (FIG. 11C). It is because solubilization and diffusion of graphite progress more rapidly at higher sintering temperature, and the high ratios of ferrite phase at lower sintering temperature are due to difficulty in diffusion of the graphite grains. However, the ratio of ferrite phase increases at a sintering temperature of over 1100° C. because of the deficiency of carbon caused by excessive formation of the iron-phosphorus-carbon compound phase (FIG. 11D).

Definite correlations of the thickness of the iron-phosphorus-carbon compound phase and the ratio of free graphite phase with the machinability index are possibly observed also in FIGS. 11A-11D and 12A-12D. The sintering temperature leading to a machinability index of 35 or less is in the range of 1,050° C. or less, wherein the ratio of iron-phosphorus-carbon compound phase becomes approximately 3 area % or less and the ratio of the iron-phosphorus-carbon compound phase having a thickness of 15 microns or more becomes 10% or less.

For the radial crushing strength, it has a correlation with the ratio of iron-phosphorus-carbon compound phase as well as with the ratio of ferrite phase, and the radial crushing strength decreases at a higher ratio of ferrite phase. The sintering temperature leading to a radial crushing strength of 500 MPa or more is in the range of 950 to 1,050° C., wherein the ratio of iron-phosphorus-carbon compound phase becomes approximately 0.2 to 3 area % and that of ferrite phase approximately 9 area % or less.

As described above, even if the composition is the same, the metallographic structure formed and the material properties thereof differ significantly depending on the sintering temperature. The sintered valve guides exhibiting favorable material properties, which are produced at a practical sintering temperature of 950° C. to 1,050° C., have the iron-phosphorus-carbon compound phase at a ratio of approximately 0.2 to 3 area %, the iron-phosphorus-carbon compound phase having a thickness of 15 microns or more at a ratio of 10% or less of the whole iron-phosphorus-carbon compound phase, and the ferrite phase at a ratio of approximately 9 area % or less.

For samples 4 and 32 to 36, the relationships between the sintering time and the ratios of respective phases in the whole composition and 32 to 36 are shown in FIGS. 13A-13D, and the relationships of the sintering time with the thickness of the iron-phosphorus-carbon compound phase and the material properties (wear amount, machinability index and radial crushing strength constant) are shown in FIGS. 14A-14D.

As apparent from FIGS. 13A, 13C and 13D, solid solution and diffusion of carbon and formation of the iron-phosphorus-carbon compound phase progress according as the sintering time is made long, and the ratio of the ferrite phase increases when the carbon diffusion is insufficient or when the iron-phosphorus-carbon compound phase is formed excessively. These results resemble the influences by sintering temperature.

Thus, the correlations of the ratio of iron-phosphorus-carbon compound phase and the thickness and the ratio of free graphite phase with the machinability index are possibly observed also in FIGS. 13A-13D and 14A-14D, and the sintering time leading to a machinability index of 35 or less is in the range of 90 minutes or less, wherein the ratio of iron-phosphorus-carbon compound phase becomes approximately 3 area % or less, the ratio of the iron-phosphorus-carbon compound phase having a thickness of 15 microns or more becomes approximately 10% or less of the whole iron-phosphorus-carbon compound phase, and the ratio of free graphite phase becomes 1 area % or more.

In addition, the sintering time leading to a radial crushing strength of 500 MPa or more is in the range of 20 min or more. Even at a sintering time of 20 min or less, the iron-phosphorus-carbon compound phase is formed in an amount sufficient for exhibiting a particular strength, and thus the metallographic structure per se is formed. The increase in the strength of sintered compacts caused by elongation of sintering time is due to growth of the necks between iron particles, and the sintering time should be decided properly according to desirable radial crushing strength and wear resistance.

For samples 4, 37, and 38, the relationships between the cooling rate and the ratios of respective phases in the whole composition are shown in FIGS. 15A-15D and the relationships of the cooling rate with the thickness of the iron-phosphorus-carbon compound phase and the material properties (wear amount, machinability index and radial crushing strength constant) are shown in FIGS. 16A-16D.

As apparent from FIGS. 15A-15D and 16A, the amounts of copper-tin alloy phase and free graphite phase do not vary according to the cooling rate, but the amounts of iron-phosphorus-carbon compound phase and ferrite phase declines, and thus the thickness of the iron-phosphorus-carbon compound phase also decreases at larger cooling rate.

Among material properties, machinability is most affected by the change of the cooling rate. In general, according as the cooling rate increases, the precipitation produced at solidification of a liquid matter become smaller in dimensions at a higher cooling rate. The iron-phosphorus-carbon compound phase precipitated in the sintered valve guide becomes thinner and the amount thereof smaller, accompanied by decrease in the ratio of ferrite phase. As a result, the machinability of the resulting compacts is improved. The precipitation of copper-tin alloy phase precipitated from liquid phase also become smaller. The cooling rate leading to a machinability index of 35 or less is in the range of 8° C./min or more, wherein the ratio of iron-phosphorus-carbon compound phase becomes 3 area % or less, the ratio of the iron-phosphorus-carbon compound phase having a thickness of 15 microns or more becomes 10% or less of the whole iron-phosphorus-carbon phase, and the ratio of ferrite phase becomes 5 area % or less.

Example 3

(Samples 39 to 49)

Sintered compact samples 39 and 42 were prepared in accordance with the powder mixture composition for each sample shown in Table 7, by repeating the operations similar to those for sample 4 (preparation of powder mixture, formation of green compact, sintering and cooling), except that the content of oxides in the raw iron powder was changed to 0.2 mass % (mill scale reduced iron powder, sample 39), 0.5 mass % (sample 40), 1.5 mass % (sample 41), or 2.0 mass % (sample 42). The composition of the entire powder mixture for each sample is shown in Table 8.

Separately, sintered compact samples 43 to 49 were prepared by repeating the operations similar to those for sample 4 (preparation of powder mixture, formation of green compact, sintering, and cooling), except that a part or all of the ore reduced iron powder is replaced with an atomized iron powder (oxide content: 0.2 mass %) and the ratio thereof in the composition of the entire powder mixture was changed to 5 mass % (sample 43), 10 mass % (sample 44), 15 mass % (sample 45), 20 mass % (sample 46), 30 mass % (sample 47), 40 mass % (sample 48), or 92.2 wt % (sample 49).

By using each of the sintered compact samples 39 to 49, the ratios of respective phases in the cross section of the metallographic structure of the sintered compact were determined by cross sectional area observation, and the abrasion, machinability and radial crushing tests thereof were performed in a similar manner to samples 1 to 27. The results are summarized in Table 9.

TABLE 7
Mixing Ratio mass %
Ore Reduced
Iron PowderAtomizedMillscaleFe—20PCu—10Sn
SampleOxideIronReducedAlloyAlloyGraphite
Nomass %PowderIron Pwd.PowderPowderPowderComments
0492.201.000.405.002.40
1091.201.405.002.40corresp. to alloy of JPB55-34858
390.2092.200.405.002.40oxide in Fe pwd.: <min., using millscale reduced Fe
4092.200.500.405.002.40oxide in Fe pwd.: min.
4192.201.500.405.002.40oxide in Fe pwd.: max.
4292.202.000.405.002.40oxide in Fe pwd.: >max.
4387.201.005.000.405.002.40
4482.201.0010.000.405.002.40
4577.201.0015.000.405.002.40
4672.201.0020.000.405.002.40
4762.201.0030.000.405.002.40
4852.201.0040.000.405.002.40
4992.200.405.002.40using atomized Fe pwd.

TABLE 8
Composition mass %
SampleMetal
NoFePCuSnCOxideComments
0491.600.084.500.502.200.92
1092.320.284.500.502.200.18corresp. to alloy of JPB55-34858
3992.520.084.500.502.200.18oxide in Fe pwd.: (min., using millscale reduced Fe
4092.060.084.500.502.200.46oxide in Fe pwd.: min.
4191.140.084.500.502.201.38oxide in Fe pwd.: max.
4290.680.084.500.502.201.84oxide in Fe pwd.: >max.
4391.650.084.500.502.200.88
4491.700.084.500.502.200.84
4591.750.084.500.502.200.80
4691.800.084.500.502.200.76
4791.900.084.500.502.200.68
4892.000.084.500.502.200.60
4992.520.084.500.502.200.18using atomized Fe pwd.

TABLE 9
Ratio in Met. Section area %
Fe—P—CEstimation Items
Compound PhaseRadial
ThicknessCu—SnFreeFerriteWearCrushing
SampleμmAlloyGraphitePhase inAmountMachinabilityStr.Con.
No<55-15≧15Phasemass %MatrixμmIndexMPaComments
041.90811542.31.421.24628.8596
1010.00145542.30.7126.130100.0602corresp. to alloy of JPB55-34858
392.26682932.21.421.24141.3627oxide in Fe pwd.: <min., using millscale Fe
402.26673032.21.421.24435.0611oxide in Fe pwd.: min.
412.27663132.31.421.25318.8567oxide in Fe pwd.: Max.
422.27673032.21.421.27112.5471oxide in Fe pwd.: >max.
432.26683112.31.421.24630.6604
442.26673122.31.421.24431.9607
452.27653322.31.421.24233.1612
462.28663222.31.421.24234.4615
472.26683112.31.421.24136.3620
482.27673212.31.421.24136.9625
492.27673032.31.221.24038.1650using atomized Fe pwd.

The relationships between the oxide content in iron powder and the ratios of respective phases in the whole composition of samples 4, 10 and 39 to 42 are shown in FIGS. 17A-17D, and the relationships of the oxide content in iron powder with the thickness of the iron-phosphorus-carbon compound phase and the material properties (wear amount, machinability index and radial crushing strength constant) are shown in FIGS. 18A-18D.

As shown in FIGS. 17A-17D and 18A, the oxide content in iron powder has almost no influence on formation of other phases such as iron-phosphorus-carbon compound phase and copper-tin alloy phase, and the oxides are present singly in the metallographic structure. On the other hand, the machinability index decreases as the oxide content increases (FIG. 18C). However, the radial crushing strength constant decreases and the wear amount increases at the same time (FIG. 13D). Thus, from FIGS. 18B-18D, the oxide content in iron powder leading to a machinability index of 35 or less, a radial crushing strength constant of 500 MPa or more and a wear amount of 360 microns or less is in the range of approximately 0.5 to 1.5 mass %.

The relationships between the blending ratio of atomized iron powder and the ratios of respective phases in the whole composition of samples 4, 10, and 43 to 49 are shown in FIGS. 19A-19D, and the relationships of the blending ratio of atomized iron powder with the thickness of the iron-phosphorus-carbon compound phase and the material properties (wear amount, machinability index and radial crushing strength constant) are shown in FIGS. 20A-20D.

There is no observable influence of the blending ratio of atomized iron powder on formation of other phases, also in FIGS. 19A-19D and 20A. As apparent from FIGS. 20B-20D, it has a small influence on material properties, and specifically, the radial crushing strength constant increases as the amount of the atomized iron powder blended increases, and the strength is highest when only an atomized iron powder is used. In addition, the machinability index decreases when the amount of the atomized iron powder blended is reduced, and in particular, the improvement in machinability is more effective when the amount of the atomized iron powder added is 30 mass % or less. The machinability index when a mill scale reduced iron powder is used is almost identical with that when an atomized iron powder is used, but is better than that when an ore reduced iron powder is used, while the radial crushing strength constant is slightly lower,compared to that when an atomized iron powder is used. Thus, if higher strength is more desirable for the sintered valve guide, use of atomized iron powder is recommended, and use of an ore-reduced iron powder is recommended when better machinability is more desirable. If an ore reduced iron powder and an atomized iron powder are used in combination, the mixing ratio of the atomized iron powder is preferably approximately 30 mass % or less, at which the improvement in machinability is more significant.

Example 4

(Samples 50 to 66)

Sintered compact samples 50 and 66 were prepared in accordance with the powder mixture composition for each sample shown in Table 10, by repeating the operations similar to those for sample 4 (preparation of powder mixture, formation of green compact, sintering, and cooling), except that, as a machinability-improving component, manganese sulfide powder was used additionally in an amount of 0.2 to 2.0 mass % (samples 50 to 55 and 62 to 66) or magnesium silicate powder in an amount of 0.2 to 2.0 mass % (samples 56 to 61 and 62 to 66). The composition of the entire powder mixture for each sample is shown in Table 11.

By using each of the sintered compact samples 50 to 66, the ratios of respective phases in the cross section of the metallographic structure of the sintered compact were determined by cross sectional area observation, and the abrasion, machinability and radial crushing tests thereof were performed in a similar manner to samples 1 to 27. The results are summarized in Table 12.

TABLE 10
Mixing Ratio mass %
OreFe—20PCu—10Sn
SampleReducedAlloyAlloyGraphiteMnSMgSiO3
NoFe Pwd.PowderPowderPowderPwd.PowderComments
0492.200.405.002.40
5092.000.405.002.400.20
5191.700.405.002.400.50
5291.400.405.002.400.80
5391.000.405.002.401.20
5490.600.405.002.401.60
5590.200.405.002.402.00MnS add.: >max.
5692.000.405.002.400.20
5791.700.405.002.400.50
5891.400.405.002.400.80
5991.000.405.002.401.20
6090.600.405.002.401.60
6190.200.405.002.402.00MgSiO3 add.: >max.
6292.000.405.002.400.100.10
6391.600.405.002.400.300.30
6491.000.405.002.400.600.60
6590.600.405.002.400.800.80
6690.200.405.002.401.001.00MgS & MgSiO3 add.: >max.

TABLE 11
Composition mass %
SampleMetal
NoFePCuSnCOxideMnSMgSiO3Comments
0491.600.084.500.502.200.92
5091.400.084.500.502.200.920.20
5191.100.084.500.502.200.920.50
5290.810.084.500.502.200.910.80
5390.410.084.500.502.200.911.20
5490.010.084.500.502.200.911.60
5589.620.084.500.502.200.902.00MnS add.: >max.
5691.400.084.500.502.200.920.20
5791.100.084.500.502.200.920.50
5890.810.084.500.502.200.910.80
5990.410.084.500.502.200.911.20
6090.010.084.500.502.200.911.60
6189.620.084.500.502.200.902.00MgSiO3 add.: >max.
6291.400.084.500.502.200.920.100.10
6391.000.084.500.502.200.920.300.30
6490.410.084.500.502.200.910.600.60
6590.010.084.500.502.200.910.800.80
6689.620.084.500.502.200.901.001.00MnS & MgSiO3 add.: >max.

TABLE 12
Ratio in Met. Section area %Estimation Items
Fe—P—CRadial
Compound PhaseCu—SnFreeFerriteWearCrushing
SampleThickness μmAlloyGraphitePhase inamountMachinabilityStr. Con.
No<55-15≧15Phasemass %MatrixμmIndexMPaComments
041.90811542.31.421.24628.8596
502.24673032.31.422.64823.8564
512.16653142.31.424.05021.3538
522.09643242.31.525.45117.5527
532.05613542.31.627.25415.0512
541.95603552.31.729.46812.5490
551.72583662.31.9228.2818.8374MnS add.: >max.
562.26653232.31.422.35122.5587
572.20643332.31.423.75320.6561
582.10633432.31.524.75816.9552
592.01613542.31.626.06313.8537
601.97613542.31.727.36512.5509
611.70573852.31.9227.1818.8397MgSiO3 add.: >max.
622.26653322.31.422.44624.4593
632.26643422.31.424.45221.9574
642.26633432.31.427.16314.4510
652.25623532.31.428.86712.5473
662.25583752.31.4247.47811.3394MnS & MgSiO3 add.: >max.

For samples 4, 10, 50 to 66, the relationships between addition of machinability-improving component powder to the powder mixture and the ratios of respective phases in the whole composition are shown in FIGS. 21A-21D, and the relationships of the amount of the machinability-improving component powder added to the powder mixture with the thickness of the iron-phosphorus-carbon compound phase and the material properties (wear amount, machinability index and radial crushing strength constant) are shown in FIGS. 22A-22D.

According to FIGS. 21C and 22C, the ratio of free graphite phase gradually increases and the machinability index gradually decreases, according as the amount of the machinability-improving component powder in the powder mixture increases. However, the effect of the machinability-improving powder is relatively mild, and the radial crushing strength constant gradually declines on the contrary as the addition amount increases, such that, at the addition amount of exceeding 1.6 mass %, the wear amount increases drastically due to brittleness of the matrix caused by sintering inhibition (diffusion inhibition). Therefore, it is difficult to improve the machinability dramatically only with the machinability-improving component powder. Accordingly, it is important to optimize the influence of the ratio and the thickness of iron-phosphorus-carbon compound phase, the ratio of free graphite phase and others on the machinability, and it is necessary to determine the composition and the manufacturing condition by considering the balance between radial crushing strength and abrasion resistance at the same time.

It must be understood that the invention is in no way limited to the above embodiments and that many changes may be brought about therein without departing from the scope of the invention as defined by the appended claims.